ABSTRACT Title of Dissertation: ATOMISTIC MODELING OF SOLID INTERFACES IN ALL-SOLID-STATE LI-ION BATTERIES Yizhou Zhu, Doctor of Philosophy, 2018 Dissertation Directed By: Assistant Professor Yifei Mo, Department of Materials Science and Engineering All-solid-state Li-ion battery based on solid electrolyte is a promising next- generation battery technology, providing intrinsic safety and higher energy density. Despite the development of solid electrolyte materials with high ionic conductivity, the high interfacial resistance and interfacial degradation at the solid electrolyte–electrode interfaces limit the electrochemical performance of the all-solid-state batteries. Fundamental understanding about the solid-solid interfaces is essential to improve the performance of all-solid-state batteries. In this dissertation, I perform first principles computation to bring new understanding about these solid interfaces. Using our developed computation approach based on large materials database, I calculated the intrinsic electrochemical stability window of solid electrolytes and predicted interphase decomposition products. I revealed the effects of different types of interphase layers on the interface stability and battery performance, and also provided interfacial engineering strategies to improve interface compatibility. Lithium metal anode can provide significantly higher energy density of Li-ion batteries. However, only a limited number of materials are known to be stable against lithium metal due to its strong reducing nature. Using first-principles calculations and large materials database, I revealed the general trend of lithium reduction behavior in different material chemistry. Different from oxides, sulfides, and halides, nitride anion chemistry exhibits unique stability against lithium metal, which is either thermodynamically intrinsic or a result of stable passivation. Therefore, many nitrides materials are promising candidate materials for lithium metal anode protection. Since solid electrolytes in all-solid-state batteries are often polycrystalline, the grain boundaries can have an important impact on the ion diffusion in solid electrolytes. I performed molecular dynamics simulations to study the ion diffusion at grain boundaries in solid electrolyte materials, and showed the distinct diffusion behavior at grain boundaries different from the facile ion transport in the bulk. In addition, I studied the order-disorder transition induced by mechanical strain in lithium garnet. Such transition can lead to orders of magnitude change in ionic diffusivity. This series of work demonstrated that computational modeling techniques can help to gain critical fundamental understandings of the solid interfaces in all-solid- state Li-ion battery, and to provide practical engineering strategies to improve the battery performance. ATOMISTIC MODELING OF SOLID INTERFACES IN ALL-SOLID-STATE LI-ION BATTERIES by Yizhou Zhu Dissertation submitted to the Faculty of the Graduate School of the University of Maryland, College Park, in partial fulfillment of the requirements for the degree of Doctoral of Philosophy 2018 Advisory Committee: Professor Yifei Mo, Chair Professor Peter Chung Professor Liangbing Hu Professor Eric Wachsman Professor Chunsheng Wang © Copyright by Yizhou Zhu 2018 Acknowledgements First of all, I would like to express my sincere gratitude to my Ph.D. advisor Professor Yifei Mo. He is both a great researcher and a great mentor, and these two add up to be a great advisor. Research side, the most important things I learned from him are: how to pick important scientific questions, how to think critically, and how to get insights from computation and modeling. Mentoring side, he put huge amount of time and efforts on training me. Just pick one example, when I was writing my very first paper in his group, we had more than 40 iterations on revising the manuscript. During the past four years, I never ever regretted joining his group and working with him. For me, this is a great four-year journey. I would also thank Prof. Chunsheng Wang, Prof. Liangbing Hu, Prof. Eric Wachsman, and Prof. Peter Chung. I learnt a lot from them during the interactions and collaborations with them. This dissertation would not have been possible without the advice from these outstanding researchers. I also want to thank my labmates and collaborators: Xingfeng He, Qiang Bai, Adelaide Mei-Chun Nolan, Yunsheng Liu, Alex Epstein, Dr. Fudong Han, Dr. Kun Fu, Dr. Wei Luo, Dr. Xiulin Fan, etc. I enjoyed the collaborations and discussions with them, and I appreciate all the help they offered to me during the past few years. Additional thanks to the coffee machine in our office, this work (literally) cannot be finished without its continuous coffee support. I also thank my friends Dr. Yubing Zhou, Jianan Li, Yan Ding, Juan He, Dr. Liangfang Zhao, Lijuan Liu, Yanxue Hong, Xuan Xiao, Wenying Li, Bowen Liang, Tongfei Chen, Yuwang Chen, Dr. Lei Zhang, Dr. Bi-cheng Zhou. They offered me ii help, suggestions, and encouragement in the past few years, especially during my hard times. Special thanks to Dr. Shuze Zhu. Thanks for all the support, understanding, and accompany, which gives me the courage to pursue my dream. Last and most importantly, I want to thank my parents and my grandparents, for their unconditional love and support throughout my whole life. iii Table of Contents Acknowledgements ................................................................................................... ii Table of Contents ..................................................................................................... iv Chapter 1. Introduction .............................................................................................. 1 1.1 The Stability of Solid Electrolyte – Electrode Interfaces and Engineering Strategies. .............................................................................................................. 3 1.2 Material Design and Interface Engineering to Enable Lithium Metal Anode .... 4 1.3 Ionic Diffusion at the Grain Boundary in Solid Electrolyte Materials ............... 6 1.4 Strain Effect on Ionic Diffusion in Lithium Garnet........................................... 7 Chapter 2. Computational Methods............................................................................ 9 2.1. Interface Equilibrium Criteria ......................................................................... 9 2.2. Compositional Phase Diagram ...................................................................... 10 2.3. Grand Potential Phase Diagram ..................................................................... 11 2.4. Evaluate the Chemical Stability of Interfaces ................................................ 12 2.5. Evaluate the Electrochemical Stability of Interfaces ...................................... 14 2.6. Density Functional Theory (DFT) Calculations ............................................. 15 Chapter 3. Electrochemical and Chemical Stability of the Solid Electrolyte–Electrode Interfaces ................................................................................................................. 17 3.1 Stability of Solid Electrolyte Materials against Lithium Metal ....................... 19 3.2 Stability of Solid Electrolyte Materials under High Potentials ........................ 25 3.3 Chemical Stability of the Solid Electrolyte against non-Li Elements .............. 27 3.4 Chemical Stability of the Solid Electrolyte-cathode Interfaces ....................... 30 3.5 Electrochemical Stability of the Solid Electrolyte-Cathode Interfaces ............ 34 3.6 Improved Interfacial Stability by Applying Coating Layers ............................ 38 3.7 Critical Role of Interphases, Interface Types and Interfacial Engineering Strategies ............................................................................................................. 43 Chapter 4. Strategies Based on Nitride Materials Chemistry to Stabilize Li Metal Anode...................................................................................................................... 53 4.1. General Trend of Materials Stability against Li Metal Reduction. ................. 55 4.2. Cation Effect on Lithium Metal Stability and Passivation ............................. 58 4.3. Anion Effect on Lithium Metal Stability and Passivation. ............................. 60 4.4 Stability and Interphase Passivation of Nitrides .............................................. 63 4.5 Stability and Interphase Passivation of Mixed-anion Nitrides ......................... 65 4.6 Discussion ..................................................................................................... 70 4.7 Conclusion ..................................................................................................... 76 Chapter 5. Ionic Transport in Grain Boundaries of Lithium Garnet .......................... 77 5.1 Computational Methods ................................................................................. 79 5.2 Diffusion and Structure Properties of Grain Boundaries in Lithium Garnet .... 83 5.3 Discussion ..................................................................................................... 87 Chapter 6. Strain Induced Order-disorder Transition in Lithium Garnet ................... 90 6.1 Effect of Mechanic Deformation on the Ionic Diffusion in lithium garnet ...... 91 6.2 Strain-induced Order-Disorder Transition in Lithium Garnet ......................... 92 6.3 Discussion ..................................................................................................... 95 iv Chapter 7. Conclusion and Future Work .................................................................. 97 7.1 Conclusions ................................................................................................... 97 7.2 Future Work .................................................................................................. 98 Publications, Patents, and Presentations ................................................................. 100 Appendices ............................................................................................................ 104 Appendix A ....................................................................................................... 104 Appendix B ....................................................................................................... 119 Appendix C ....................................................................................................... 144 Bibliography ......................................................................................................... 145 v Chapter 1. Introduction The continued drive for high energy density Li-ion batteries has imposed ever- stricter requirements on the electrolyte materials. Current organic liquid electrolytes in commercial Li-ion batteries are flammable, causing notorious safety issues. Moreover, the limited electrochemical window of the organic liquid electrolytes limits the choice of electrode materials and hence the achievable energy density of the Li-ion batteries. All-solid-state Li-ion batteries (ASLiBs) based on ceramic solid electrolyte (SE) materials are a promising next-generation energy storage technology with high energy density and enhanced cycle life.1,2 ASLiBs benefit from extra safety provided by intrinsically non-flammable ceramic solid electrolytes materials, thus is considered to be a promising ultimate solution to the notorious safety issue of Li-ion batteries. Thanks to the recent discovery and development of Li ionic conductor materials, such as Li lithium thiophosphates1,3,4 and Li garnet-type materials,5,6 high Li ionic conductivity of 1 to 10 mS/cm comparable to the organic liquid electrolytes has been achieved in the solid electrolyte materials. Moreover, the claimed outstanding stability of ceramic solid electrolyte materials may enable Li metal anode and high-voltage cathodes,1,2 which can significantly increase the energy density of Li-ion batteries.7-9 However, despite the high bulk ionic conductivity achieved in the solid electrolyte materials, the electrochemical performance of ASLiBs is still not comparable to those assembled with liquid electrolytes. High interfacial resistance and interfacial degradation are the key limiting factors of the ASLiB performance, including poor cyclability, low coulombic efficiency, etc. Therefore, to resolve the high interfacial resistance and to further improve the ASLiBs performance, fundamental understanding of the origin of high 1 interfacial resistance is required. In addition, specific interfacial engineering and material design strategies are needed to guide the future development of ASLiBs with improved interfacial stability and interfacial ionic transport. Figure 1.1 Schematic illustration of all-solid-state Li-ion batteries.10 Ion transport at multiple interfaces, including the interfaces between solid electrolytes and electrodes, grain boundaries in solid electrolytes, are involved during battery operation. The properties of these interfaces, especially the ionic transport properties and interface stabilities, are critical to the battery performance. Computational modeling can provide fundamental understanding of the properties and mechanisms at these buried interfaces, which are difficult to access and characterize in experiments. In this dissertation, I will show how we use computational modeling to investigate the interfacial properties in ASLiBs. Using first principles calculations and molecular dynamics simulations, we gain the fundamental understanding of materials interfaces and provide design strategies to develop interfaces with improved stability for ASLiBs. 2 1.1 The Stability of Solid Electrolyte – Electrode Interfaces and Engineering Strategies. Currently, one of the key bottleneck problem of ASLiBs is the high interfacial resistance at the electrolyte–electrode interface. Previously, the high interfacial resistance at these solid-solid interfaces was attributed to multiple factors, including poor physical contact, interfacial degradation, carrier depletion caused by space-charge effect. However, the intrinsic chemical and electrochemical stability of these interfaces were neglected. Li Metal Anode Ceramic Solid Electrolyte Cathode LMO Fast Li+ transport High interfacial resistance Figure 1.2. Despite the fast bulk ionic transport achieved in bulk solid electrolyte materials, the high interfacial resistance is the limiting factor in ASLiB performance. Early reports on the electrochemical stability window of solid electrolyte materials were based on conventional cyclic voltammetry measurements. For example, Li10GeP2S12 (LGPS) and Li7La3Zr2O12 (LLZO) were reported to be stable from 0 to 5 V. 1,9,11,12 However, later experimental and computational studies questioned the claimed stability of solid electrolyte materials against Li metal and at high voltages.13,14 It is critical to assess the intrinsic electrochemical stability window of the solid electrolyte materials. Cycling outside of the intrinsic electrochemical window of the solid electrolyte materials would lead to the 3 spontaneous formation of an interphase layer, and these interphase layers can have significant effects on the interfacial resistance and the overall performance of the ASLiBs. The continued decompositions at the solid electrolyte–electrode interfaces may lead to interfacial degradation and poor coulombic efficiency of the ASLiBs. In addition, the interphase layers may be poor ionic conductor, which would result in high interfacial resistance. In chapter 3, we use computation to investigate the thermodynamics of the interfaces between solid electrolytes and electrodes in ASLiBs. We found that the intrinsic electrochemical windows of solid electrolytes are limited, which leads to the spontaneous formation of interphase layer. I will discuss how the properties of the formed interphase affect the battery performance. Based on their properties, we classify different solid electrolyte–electrode interfaces into different types, and provide strategies to improve the interface stability and battery performance. 1.2 Material Design and Interface Engineering to Enable Lithium Metal Anode Lithium metal is one of the most promising anode materials. It has the highest theoretical specific capacity and the lowest standard potential, which can significantly boost the energy density of Li-ion batteries. However, strong reducing nature of lithium metal limits its application, causing incompatibility and undesired reaction with most electrolytes. The lack of long-term stability between lithium metal anode and electrolyte leads to quick battery capacity fade and cell failure. In addition, the undesired growth of Li dendrite during cell cycling causes short circuiting in the cell, leading to catastrophic cell failure and safety 4 issues. Despite decades of research efforts, the goal of stabilizing lithium metal anode over many charge-discharge cycles and long battery life time has not been fully realized. Figure 1.3. Due to the strong reducing nature of lithium metal, identify potential protection coating materials to protect electrolytes against lithium metal reduction is crucial for the development of lithium batteries. We propose to use nitride-based materials chemistry, which has unique intrinsic stability against Li metal reduction, to engineer the anode interface and protect electrolytes from reduction. To enable Li metal anode, identify materials that are intrinsically stable or can be used as protection coating is critical. However, despite that many lithium solid electrolytes were previously reported to be stable against Li metal, multiple experimental and first principles studies confirmed that most solid electrolytes such as Li10GeP2S12 (LGPS), NASICON-type Li1.3Al0.3Ti1.7(PO4)3 (LATP), and perovskite Li0.33La0.56TiO3 (LLTO), are not thermodynamically stable against Li metal and will be reduced at the interface. 13-20 Therefore, developing Li-stable materials to protect the solid electrolytes from lithium metal reduction is critical. Previous discoveries on Li-metal-stable materials are mostly based on a trial-and- error approach. Only a few materials, such as lithium binaries Li2O, Li2S, Li3N, and LiF, 5 are known to be stable against Li metal. A more comprehensive understanding on the general trend of material stability against Li metal reduction is needed. In chapter 4, I will show how we use first principles calculations powered by large- scale materials database, to uncover lithium-stable materials out of a variety of chemistry across the periodic table. we discovered that nitrides exhibit thermodynamically intrinsic stability against lithium metal. Such unique stability is lacking in other materials chemistry, such as oxides, sulfides, and halides, which are commonly used for lithium protection. The intrinsic stability of many nitride compounds against lithium reduction makes them ideal candidate materials to protect lithium metal anode over long term. 1.3 Ionic Diffusion at the Grain Boundary in Solid Electrolyte Materials The recent advancement in all-solid-state Li-ion batteries is mainly driven by the discovery and development of novel solid electrolyte materials. Despite that they are solid materials, some of them can provide liquid-like high ionic conductivity. The fast bulk ionic transport in these has drawn a lot of research effort to understand the underlying mechanism of its fast ion diffusion.21-25 However, the employed solid electrolytes in all-solid-state Li-ion batteries are usually polycrystalline materials. Therefore, besides the bulk ionic transport, the ion diffusion at grain boundaries can also have a significant impact on the overall ionic transport and thus battery performance. 26,27 Grain boundaries usually exhibit different structures and chemistry from bulk, which can lead to distinct ionic diffusion behavior. Moreover, some recent studies revealed 6 that the grain boundaries are related to lithium dendrite growth, as lithium dendrite show preference growth inside the grain boundary of solid electrolyte.28,29 Despite the extensive study on the bulk properties, current understanding on the grain boundaries in solid electrolytes are still limited. We select Li-stuffed garnet Li7La3Zr2O12 (LLZO) as a model system to study the grain boundary diffusion in super ionic conductor materials. In chapter 5, I perform molecular dynamics simulation to investigate the ionic diffusion at grain boundaries in lithium garnet. We found that grain boundaries in lithium garnet exhibit slower ion transport compared to the bulk. Through analyzing the atomistic structure of grain boundaries, we revealed that the framework shows limited change and that lithium ions exhibit different configuration at the grain boundaries compared to bulk. 1.4 Strain Effect on Ionic Diffusion in Lithium Garnet Mechanical strain is can significantly affect the ion diffusion in ionic conductor materials.30,31. The strain on the solid electrolytes may be caused by the volume change of electrode materials during cycling or by the applied pressure during the assembling of ASLiBs. Lithium garnet is a promising solid electrolyte for all-solid-state Li-ion batteries. However, there is little information about the effect of mechanical strain on the ion diffusion of lithium garnet. In chapter 6, I will show our simulation results about the strain effects on the ionic diffusion in lithium garnet. We found that there is a strong coupling between the lattice parameters and lithium sublattice ordering. Applying mechanical strain can induce an order- disorder transition in lithium sublattice of lithium garnet. This strain-induced order-disorder 7 transition of lithium sublattice can have a significant effect on the Li-ion diffusion in lithium garnet. 8 Chapter 2. Computational Methods In this chapter, I introduce our computation method to evaluate the interface stability in the all-solid-state Li-ion batteries. 2.1. Interface Equilibrium Criteria To evaluated the stability of the interfaces, we consider the following equilibrium criteria. These equilibrium criteria must be satisfied for the interfaces at thermodynamic equilibrium. 1) Equilibrium of neutral Li Stable interface should reach the equilibrium of neutral Li (under applied potential) across the interfaces. Through this equilibrium, the Li chemical potential µLi should be equal at the contact points between solid electrolyte and electrode materials. This criterion is equivalent to the equilibrium of the materials (i.e., the electrochemical stability) against different electrostatic potential ϕ. The equilibration of µLi and ϕ are facilitated by the good mobility of Li ions in the solid electrolytes and electrode materials. 2) Equilibrium of non-Li elements Stable interface should also reach the equilibrium of neutral non-Li elements across the interfaces. Similar to the equilibrium of Li, the equilibrium of any non-Li element M at the interface requires the equivalence of chemical potential µM at the contact points between the solid electrolytes and electrode materials, though non-Li elements, such as Co and S, are usually not as mobile as Li. 3) The full thermodynamic equilibrium This criterion is to evaluate the chemical stability of the interface, which is 9 determined by whether the two materials mixed by a certain ratio can have an exothermic reaction to form other phases as the phase equilibria at the same composition. Such reaction would not exist if two materials were chemically stable against each other, where the phase equilibria would be equivalent to two original materials. If the phase equilibria were different from the combination of original two materials, we calculate the decomposition energy to evaluate the driving force. This criterion considered the full equilibrium of all elements, while only one element is allowed to equilibrate in criteria 1 and 2. 4) The equilibrium of the two materials in contact under an applied potential. This criterion is to evaluate the electrochemical stability of the interface, which is different from the chemical stability in the absence of an applied potential as in criterion 3. We determined the thermodynamic phase equilibria of the interface as a function of the applied potential. Similar to criterion 3, an exothermic reaction to from other phases at the applied potential suggests thermodynamically favorable decomposition reaction. We use phase diagram calculations to evaluate the aforementioned criteria. Specifically, we calculate the phase equilibria under different conditions and then evaluate the phase/interface stability and thermodynamic driving force of decomposition. 2.2. Compositional Phase Diagram The compositional phase diagrams were constructed to evaluate the phase equilibria of a given solid electrolyte or electrode phase with the composition C. The phase equilibria were determined by constructing the convex energy hull of all relevant phases in the compositional phase diagram.32 The phase equilibria at the composition C corresponding to the energy minimum Eeq(C) were identified by comparing the energy of all relevant phases 10 in its compositional space. The phase stability of the investigated phase was evaluated using the decomposition energy ΔED, Δ𝐸#(phase) = 𝐸,-(𝐶) − 𝐸(phase), Eq. 1 of a given phase to its phase equilibria. ΔED is the negative of energy above hull. Li O2 LiO2 Li2O2 Li2O2 Li2O Li2O Li2O2 Li2O LiO2 O Energy 2 above hull Li Molar fraction nO / (nLi + nO) Figure 2.1. An example phase diagram showing the phase stability and electrochemical stability in Li-O system. The convex hull was connected by four stable phases: Li, Li2O, Li2O2, and O2, while LiO2 is an unstable phase with a positive energy above hull (intersect). The left and right color bar shows the electrochemical stability of the stable phases. Each color block corresponds to the most stable phase under the given chemical potential of Li or O. 2.3. Grand Potential Phase Diagram Grand potential phase diagrams were constructed to evaluate the stability of a material in equilibrium with an external environment.13,32 The grand potential phase 11 Chemical potential μLi (eV) Chemical potential μO (eV) diagram identifies the phase equilibria Ceq(C, µM) of a given phase with the composition C in equilibrium with the chemical potential µM of element M. The given phase is stable within a certain range of µM. Outside the µM stable range of the phase, the composition of the phase equilibria Ceq(C, µM) have a different number of element M from the original composition C, where the number of element M is changed by ΔnM. The decomposition reaction energy at the chemical potential µM is calculated as Δ𝐸01,2# (phase, 𝜇5) = 𝐸,- 6𝐶,-(𝐶, 𝜇5)7 − 𝐸(phase) − Δ𝑛5 ∙ 𝜇5. Eq. 2 Using the same scheme in previous studies,13 the electrode potential ϕ was considered as a part of the Li chemical potential µLi, 𝜇 =:;(𝜙) = 𝜇:; − 𝑒𝜙, Eq. 3 where 𝜇=:; was the chemical potential of Li metal. In this study, 𝜇5 was referenced to the elementary state 𝜇=5 , and the applied potential ϕ was referenced to Li metal. The electrochemical window of the phase was estimated as the range of ϕ, where the phase is neither oxidized nor reduced. The decomposition reaction energy at applied voltage ϕ was calculated as Δ𝐸01,2? (phase, 𝜙) = Δ𝐸 01,2 ? (phase, 𝜇:;(𝜙)). Eq. 4 2.4. Evaluate the Chemical Stability of Interfaces In this study, we considered the interface as a pseudo-binary33 of the solid electrolyte and the electrode, which has a composition 𝐶@2A,BCD,(𝐶EF, 𝐶FG,DAH0I,, 𝑥) = 𝑥 ∙ 𝐶EF + (1 − 𝑥) ∙ 𝐶FG,DAH0I, Eq. 5 where CSE and CElectrode are the compositions (normalized to one atom per formula) of SE 12 and electrode materials, respectively, and x is the molar fraction of SE which varies from 0 to 1. The energy of the interface pseudo-binary, 𝐸@2A,HBCD,(SE,Electrode, 𝑥) = 𝑥 ∙ 𝐸(SE) + (1 − 𝑥) ∙ 𝐸(Electrode) Eq. 6 was set to a linear combination of the electrolyte and electrode energies. The decomposition energy of the interface pseudo-binary was calculated similar to Eq. 1 as Δ𝐸#(SE, Electrode, 𝑥) Eq. 7 = 𝐸,-U𝐶;2A,HBCD,(𝐶EF, 𝐶FG,DAH0I,, 𝑥)V − 𝐸@2A,HBCD,(SE, Electrode, 𝑥) ΔED(SE, Electrode, x) includes the decomposition energy ΔED from Eq. 1 if the SE or the electrode are not thermodynamically stable. We defined ΔED,mutual as the reaction energy between phase equilibria of SE and electrode materials, Δ𝐸#,WXAXCG(SE, Electrode, 𝑥) = Δ𝐸#(SE, Electrode, 𝑥) − 𝑥 ∙ Δ𝐸#(SE) − (1 − 𝑥) ∙ Eq. 8 Δ𝐸#(Electrode). The ΔED,mutual describes the mutual reaction between SE and electrode materials excluding the decomposition energy ΔED (Eq. 1) of the SE and electrode themselves. Since the phase equilibria and the reaction energies vary with the pseudo-binary composition, we identified the minimum of the mutual reaction energy, Δ𝐸#,W;2,WXAXCG(SE, Electrode) = min _Δ𝐸#,WXAXCG(SE, Electrode, 𝑥)`, \∈(=,^) Eq. 9 which reaches the minimum at x = xm. At the same interface pseudo-binary composition xm, we calculated the total decomposition energy as Δ𝐸#,W;2,A0ACG(SE, Electrode) =Δ𝐸#(SE, Electrode, 𝑥W). Eq. 10 It is worth noting that the identified xm corresponds to the most exothermic decomposition reactions, while the actual interphase layer may differ from the most favorable 13 thermodynamic phase equilibria and may have a distribution of elemental profile and materials compositions across the interfaces. 2.5. Evaluate the Electrochemical Stability of Interfaces The electrochemical stability of the interface was evaluated for the interface pseudo- binary (Eq. 5) using the grand potential phase diagram described in section 2.3 and the previous study33. The decomposition reaction energy Δ𝐸01,2# (SE, Electrode, 𝑥, 𝜙) at applied voltage ϕ was calculated as Δ𝐸01,2# (SE,Electrode, 𝑥, 𝜙) = 𝐸,- 6𝐶,-(𝐶;2A,HBCD,(𝐶EF,𝐶FG,DAH0I,, 𝑥), 𝜇5)7 − 𝐸;2A,HBCD,(SE,Electrode, 𝑥) − Eq. 11 Δ𝑛:; ∙ 𝜇:;(𝜙). The decomposition reaction energy Δ𝐸01,2# was normalized to the number of non- Li atoms because the number of Li changes with the phase equilibria at different applied voltage. Similar to section 2.4, we defined the mutual reaction energy Δ𝐸01,2#,WXAXCG as Δ𝐸01,2#,WXAXCG(SE, Electrode, 𝑥, 𝜙) Eq. 12 = Δ𝐸openD (SE, Electrode, 𝑥, 𝜙) − 𝑥 ∙ Δ𝐸 open D (SE, 𝜙) − (1 − 𝑥) ∙ Δ𝐸 open D (Electrode, 𝜙). to evaluate the reaction energy of the “mutual” reaction between electrolyte and electrode excluding the decomposition energy Δ𝐸01,2# (SE,𝜙) and Δ𝐸 01,2 # (Electrode,𝜙) of the SE and electrode themselves, respectively. Since the phase equilibria and reaction energies are dependent on the pseudo-binary composition, we identified the minimum of mutual reaction energy at a given applied voltage ϕ as Δ𝐸01,2#,W;2,WXAXCG(SE, Electrode, 𝜙) = min _Δ𝐸 01,2 #,WXAXCG(SE, Electrode, 𝑥, 𝜙)` Eq. 13 \∈(=,^) for the mutual reaction between the SE and electrode materials under the applied voltage ϕ, 14 similar to the previous study33. If Δ𝐸01,2#,W;2,WXAXCG(SE, Electrode,𝜙) ≠ 0, we calculated the total decomposition reaction energy as Δ𝐸01,2 01,2#,W;2,A0ACG(SE, Electrode, 𝜙) = 𝛥𝐸# (SE, Electrode, 𝑥W(𝜙), 𝜙), Eq. 14 where xm(ϕ) is at the minimum point of the mutual reaction energy Δ𝐸01,2#,W;2,WXAXCG identified in Eq. 13. If Δ𝐸01,2#,W;2,WXAXCG(SE, Electrode,𝜙) = 0, the minimum point xm of ΔED,min,mutual in Eq. 9 is used in Eq. 14. For the SE–LCO interfaces, we considered the voltage range of ϕ from 2 V to 5 V for Li-ion batteries. 2.6. Density Functional Theory (DFT) Calculations Most compounds used in this work to construct phase diagrams were from the Materials Project (MP)34 and Inorganic Crystal Structure Database (ICSD). We also added additional lithium ternary compounds, which were predicted using substitution prediction algorithm developed by Hautier et al. The energies of most materials were obtained from the MP database. For materials that are not available in the MP database, density functional theory (DFT) computation based on Perdew–Burke–Ernzerhof (PBE) generalized gradient approximation (GGA) functionals 35 described by the projector augmented-wave approach36 as implemented in the Vienna Ab initio Simulation Package (VASP)37 were performed. The parameters of DFT calculations, such as the plane-wave energy cutoff and k-points density, were consistent with those used for the MP.38 The energy correction schemes for transition metal, O, S, N, and F elements were applied as in the MP.39,40 The calculated reaction energies and voltages were based on DFT energies as in previous studies,13,32 where the contributions of PV terms 15 and entropy terms were neglected. We also require the electronic conductivity property of the formed interphases to evaluate the interface properties. Due to the lack of electronic conductivity data from other sources, we used the valence state of transition metal cations and the band gap calculated in GGA to judge whether a phase is electronic insulating. Phases that have transition metal cations at their highest valence state and that have a calculated band gap wider than 0.7 eV in GGA are considered electronic insulating. Though the band gap calculated in GGA is systematically underestimated, we chose our cut-off value as 0.7 eV, which is the GGA calculated gap of Li3P, a known poor electronic conductor.19,41 16 Chapter 3. Electrochemical and Chemical Stability of the Solid Electrolyte–Electrode Interfaces The high interfacial resistance at the electrolyte–electrode interface is a crucial problem in ASLiBs, limiting the power and rate performances of the ASLiBs.2 The high interfacial resistance is attributed to poor interfacial contact,42 the mechanical failure of the contacts,43 interfacial degradation due to mutual diffusion,44 or the formation of lithium- depleted space-charge layer.2 The lithium-depleted space-charge layer formed at the interface due to the large chemical potential differences between the sulfide solid electrolytes and the oxide cathode materials is often blamed as the cause of high interfacial resistance.2 The possibility of the decomposition at the interfaces is often overlooked since the solid electrolytes are claimed to have excellent stability.1 The claimed outstanding stability of the solid electrolyte materials is based on the widely reported electrochemical window of 0-5 V from cyclic voltammetry (CV) measurements.1,9,11,12 However, some recent experimental and computational studies questioned the claimed stability of solid electrolyte materials against Li metal and at high voltages. For example, the reduction and oxidation of Li10GeP2S12 (LGPS) at low and high potentials, respectively, in contrast to the originally claimed 0-5 V electrochemical window, have been demonstrated by first principles computation13 and experiments.14 Recent in-situ This chapter has been published in Y. Zhu, X. He, Y. Mo, Origin of outstanding stability in the lithium solid electrolyte materials: insights from thermodynamic analyses based on first principles calculations. ACS Applied Materials & Interfaces, 7, 23685- 23693, (2015) Y. Zhu, X. He, Y. Mo, First principles study on electrochemical and chemical stability of solid electrolyte– electrode interfaces in all-solid-state Li-ion batteries. Journal of Materials Chemistry A, 4, 3253-3266, (2016) 17 X-ray photoelectron spectroscopy (XPS) experiments also observed the interfacial decomposition of LiPON,45 lithium lanthanum titanate,18 and NASICON-type17 solid electrolyte materials against Li metal. Transmission electron microscopy (TEM) and energy dispersive X-ray (EDX) studies identified interphase layers of tens to a hundred nanometers at the LiCoO2 electrode-electrolyte interfaces with Li2S-P2S5 46 or LLZO.44 These interphase layers were attributed to the structural disordering and the mutual diffusion of non-Li elements, such as Co and S, across the interface. And they have significant effects on the interfacial resistance and the overall performance of the ASLiBs. For example, continued decompositions may lead to interfacial degradation and poor coulombic efficiency of the ASLiBs. In addition, the interphase layers may be poor ionic conductor, which would result in high interfacial resistance. However, the formation mechanisms of such interphase layers and the effects on the performance of ASLiBs have rarely been discussed. Currently, there is limited knowledge about the formation mechanisms of the interphase layers at the buried solid electrolyte–electrode interfaces in ASLiBs. The formation of interphase layers in ASLiBs may be caused by three mechanisms: 1) The reduction or oxidation of the solid electrolyte materials under an applied potential due to their limited electrochemical windows; 2) The chemical reactions between the solid electrolyte and the electrode materials caused by the chemical incompatibility between the solid electrolyte and the electrodes; and 3) The electrochemical reactions of the solid electrolyte–electrode interfaces during the cycling of the ASLiBs. The thermodynamics for these three mechanisms, which correspond to the chemical 18 and electrochemical stability of the interfaces, are well defined. To the best of our knowledge, such thermodynamics information has not been available for the solid electrolyte and electrode materials in the context of ASLiBs. Understanding the origin of the interfacial decomposition and the formation mechanisms of interphase layers are critical for resolving the issue of high interfacial resistance in ASLiBs and for guiding the development of ASLiBs. In this chapter, we employ a computational scheme based on first principles calculations to obtain the thermodynamics of the interfaces between commonly used solid electrolytes and electrodes in ASLiBs. Our results confirmed the strong thermodynamic driving force for the decomposition at interfaces in ASLiBs due to the limited electrochemical window of the solid electrolyte materials and the poor chemical compatibility between the solid electrolyte and the electrodes. In addition, some of the formed interphases and the applied coating layers were demonstrated to improve the interface stability and the performance of the ASLiBs. The interfaces based on different solid electrolyte and interphases were classified into different types, and the strategies to address the interfacial problems were proposed and demonstrated. 3.1 Stability of Solid Electrolyte Materials against Lithium Metal We first evaluated the electrochemical stability of solid electrolyte materials against Li metal and at low voltages. The crystal structures and the phase equilibria of the commonly used solid electrolyte materials investigated in this study are summarized in Appendix. The phase equilibria, i.e., the phases with the lowest energy, in equilibrium with Li metal were identified by the Li grand potential phase diagrams. 19 Anode / Solid Reductant Electrolyte μ~Li+ μ~e- Reduction potential Solid electrolyte μLi Reduction (a) (b) Figure 3.1. (a) Electrochemical window (solid color bar) of solid electrolyte and other materials. The oxidation potential to fully delithiate the material is marked by the dashed line. (b) Schematic diagram about the change of Li chemical potentials µLi (black line) and the electrochemical potential 𝜇e:;f (blue dash line) and 𝜇e,g (red dash line) across the interface between the anode and the solid electrolyte. Since the actual profile of 𝜇e,g determined by the charge carrier distribution may be complicated,47,48 the profiles of chemical and electrochemical potential shown here are schematic and may not be linear. The vertical scale is for the electrostatic potential or the voltage referenced to Li, and is reversed for the chemical potential or electrochemical potential. The solid electrolyte materials are not thermodynamically stable against Li metal (Table 3.1) and are reduced at low voltages with highly favorable decomposition energy (Fig. 3.1a and Table 3.1). In contrast, the Li binary compounds, such as LiF, Li2O, Li2S, Li3P, and Li3N, are thermodynamically stable against Li metal (Fig. 3.1a). The lithiation and reduction of Li10GeP2S12 (LGPS) starts at 1.71 V, and the LGPS in equilibrium with Li 20 Higher ϕ / Higher μ Lower μ Interphase metal is eventually lithiated into the phase equilibria consisting Li15Ge4, Li3P, and Li2S. The Li reduction of the LGPS into these reaction products has a highly favorable reaction energy of -1.25 eV/atom (-3014 kJ/mol of LGPS) at 0 V (Table 3.1). In agreement with our computation, the reduction of LGPS starting at 1.71 V and the formation of Li-Ge alloy after the reduction have been demonstrated in CV and XPS experiments, respectively.14 Other sulfides materials, such as Li3PS4, Li4GeS4, Li6PS5Cl, and Li7P2S8I, are reduced at a similar voltage of ~1.6-1.7 V. The reduction potential is mostly governed by the reduction of P and Ge in the materials, and the reduction products include Li3P and Li2S at 0 V. For those materials containing Ge, Cl, and I elements, Li-Ge alloy, LiCl, and LiI are formed, respectively, as a part of phase equilibria at 0 V. The Li7P3S11 is reduced at a voltage of 2.28 V into Li3PS4 with a small decomposition energy, and the majority of the reduction starts at 1.71 V due to the lithiation of Li3PS4. The decomposition energy for all these solid electrolyte decreases with the potential to ~ -1 eV/atom at 0 V (Table 3.1), indicating the highly favorable reduction reactions of the sulfide solid electrolytes. The reduction of oxide solid electrolyte materials Li0.33La0.56TiO3 (LLTO) and Li1.3Ti1.7Al0.3P3O12 (LATP) starts at a voltage of 1.75 and 2.17 V, respectively. Our predicted reduction potential of LLTO is in good agreement with the value of 1.7 - 1.8 V reported in the CV experiments.20,49 The calculations also found the reduction of Ti4+ in LLTO and LATP into Ti3+ or lower valences at low voltages (Table 3.1). The reduction of Ti is a widely known problem and is observed at the interfaces of LLTO18 and LATP17 with Li metal by in-situ XPS spectroscopy. In addition, the reduction of Ge-containing oxide materials Li1.5Al0.5Ge1.5(PO4)3 (LAGP) and Li3.5Zn0.25GeO4 (LISICON) starting at 2.70 and 1.44 V respectively, and Li-Ge alloys are formed at low voltages (Fig. 3.1a and Table 3.1). 21 The reductions of LAGP and LISICON are consistent with the experiment studies.50-52 The good agreements between our computation results and many experiments demonstrated the validity of our computation scheme. Table 3.1 Reduction reaction of the solid electrolyte materials at Li metal. ED Phase equilibria with Li metal (eV/atom) Li2S Li2S (stable) 0 LGPS Li15Ge4, Li3P, Li2S -1.25 Li3PS4 Li3P, Li2S -1.42 Li4GeS4 Li15Ge4, Li2S -0.89 Li7P3S11 Li3P, Li2S -1.67 Li6PS5Cl Li3P, Li2S, LiCl -0.96 Li7P2S8I Li3P, Li2S, LiI -1.26 LiPON Li3P, Li3N, Li2O -0.66 LLZO Zr (or Zr3O), La2O3, Li2O -0.021 LLTO Ti6O, La2O3, Li2O -0.35 LATP Ti3P, TiAl, Li3P, Li2O -1.56 LAGP Li9Al4, Li15Ge4, Li3P, Li2O -1.99 LISICON Li15Ge4, LiZn, Li2O -0.77 Our calculations found the Li reduction of the solid electrolyte materials that are thought to be stable against Li. For example, LiPON, which is calculated using Li2PO2N as a representative of the material class, shows a reduction potential of 0.69 V. The final decomposition products of LiPON in equilibrium with Li metal are Li3N, Li2O, and Li3P (Table 3.1), which are consistent with the in-situ XPS observations.45 Although the calculated decomposition energy of LiPON is as large as -0.66 eV/atom at 0 V (Table 3.1), 22 LiPON is known to be compatible with Li metal as demonstrated by many experimental studies.12,53 Similarly, Li3PS4 and Li7P2S8I, which are reported to be compatible with Li metal anode,9,54,55 are reduced against Li metal and at low voltages (Table 3.1 and Fig. 3.1). Therefore, the stability of these solid electrolyte materials against Li metal is not thermodynamically intrinsic. The decomposition products, which form an interphase between the solid electrolyte and electrode, passivate the solid electrolyte and inhibit the continuous decomposition. For example, the decomposition products of LiPON, Li3PS4, and Li7P2S8I are Li binary compounds, such as Li2O, Li2S, Li3P, Li3N, and LiI, formed at the Li reduction. The interphase consisting of these decomposition products are stable against the high 𝜇:; of Li metal (Fig. 3.1a), which is beyond the reduction potential (cathodic limit) of the solid electrolyte (Fig. 3.1b). At the equilibrium, the redistribution of Li+ and other charged carriers (such as electron e-) are formed at the interface to account for the potential drop across the electrode-electrolyte interface.56 The electrochemical potential of the highly mobile Li+, 𝜇e:;f, which includes the electrostatic potential energy, is constant across the interface. In contrast, the electrochemical potential of the electronic carrier 𝜇e,g (red line in Fig. 3.1b) decreases significantly in the interphase from the anode to the solid electrolyte, since these interphases have poor electronic mobility and conductivity. Therefore, the Li chemical potential µLi (black line in Fig. 3.1b), which equals to the sum of 𝜇e:;f and 𝜇e,g, decreases in the interphase from the anode to the solid electrolyte. The high 𝜇:; from the anode decreases to be within the electrochemical window of the solid electrolyte after the passivation of the decomposition interphase. As a result, the decomposition of the solid electrolyte has no thermodynamic driving force to continue into the bulk. The solid electrolyte is stabilized by 23 the decomposition interphases, which essentially serve as solid-electrolyte-interphases (SEIs) in all-solid-state Li-ion batteries. In summary, the SEI of the decomposition interphase decreases the high Li chemical potential µLi applied on the solid electrolyte, and bridges the Li chemical potential gap between Li metal and the solid electrolyte. This passivation mechanism explained the observed Li metal compatibility of LiPON, Li3PS4, and Li7P2S8I. The passivation mechanism relies on the electronic insulating properties of the decomposition interphase layers to stabilize the solid electrolyte, and is not active if the interphase layer is electronically conductive. For example, the reduction of LGPS, LAGP, and LISICON with Li metal forms electronically conductive Li-Ge alloys, and the lithiation of LLTO and LATP forms titanates with Ti of 3+ or lower valences. The decomposition interphases for these solid electrolytes at Li reductions are mixed electronic and ionic conductors. The electronic conductivity in the interphase cannot account for the drop of 𝜇e,g across the interface regardless of the specific electron transport mechanism being metallic, band, or polaronic conduction. These mixed conductor interphases cannot account for the 𝜇:; drop as the change of both 𝜇e:;f and 𝜇e,g would be small across the interphase. As a result, the solid electrolyte is still exposed to the high 𝜇:; of the anode, and the reduction reaction continues into the bulk. In addition, the mixed electronic and ionic conductor interphase facilitate the kinetic transport of Li ion and electrons for the decomposition reactions.18 The absence of the passivation mechanism explains the lithiation and reduction of LGPS, LLTO, LATP, LAGP, and LISICON observed in the CV experiments. It is worth noting that garnet LLZO shows the lowest reduction potential of as low as 0.05 V against Li and the least favorable decomposition reaction energy of only 0.021 24 eV/atom (49 kJ/mol of LLZO) at 0 V among all solid electrolyte materials examined (Table 3.1). Given such small reaction energy, the Li reduction of garnet is likely to be kinetically inhibited, and the reduction products of Li2O, Zr3O and La2O3 (Table 3.1) may provide passivation to the material. These explain the widely reported 0 - 5 V window of garnet from the CV measurements in the literature.5,57 The reduction of garnet at 0.05 V forms Zr3O (Fig. 3.1), following another plateau at 0.004 V to form Zr based on the DFT GGA energies. Since these small values of energy and voltage is beyond typical accuracy of DFT and the approximations of the scheme, it is inconclusive whether the garnet LLZO is reduced to Zr3O or Zr at 0 V or against Li metal. Nevertheless, the formation of Zr would be thermodynamically favorable at a potential significantly lower than 0 V, which corresponds to applying high current density at the Li-LLZO interface. Recent report of instability of garnet against Li at elevated temperatures of 300 °C may be an indication of the limited stability of garnet against Li metal,58 as the diffusion and phase nucleation are facilitated at high temperatures. 3.2 Stability of Solid Electrolyte Materials under High Potentials The oxidation reactions of the solid electrolyte materials were investigated using the same method. The LGPS material is delithiated and oxidized starting at 2.42 V (Fig. 3.1), and the final oxidation products of P2S5, GeS2, and S are formed at the equilibrium oxidation potential of 2.48 V (Table 3.2). The oxidation potential of the LGPS is confirmed by the CV experiment.14 Similar to Li2S, all sulfide solid electrolytes such as Li3PS4, Li4GeS4, Li7P3S11, and Li7P2S8I are oxidized at ~ 2.5 V to form S (Table 3.2). The other elements, such as P and Ge, are usually oxidized into P2S5 and GeS2, respectively. The oxidation 25 reactions of sulfide solid electrolytes are highly favorable at 5 V as described by the decomposition energy (Table 3.2). Table 3.2. Oxidation reaction of the solid electrolyte materials at 5 V. ED Phase equilibria at 5V (eV/atom) Li2S S -1.77 LGPS GeS2, P2S5, S -1.04 Li3PS4 S, P2S5 -0.97 Li4GeS4 GeS2, S -1.12 Li7P3S11 S, P2S5 -0.90 Li6PS5Cl P2S5, S, PCl3 -1.27 Li7P2S8I P2S5, S, I2 -1.00 LiPON PNO, P2O5, N2 -0.69 LLZO O2, La2O3, La2Zr2O7 -0.53 LLTO O2, TiO2, La2Ti2O7 -0.15 LATP O2, TiP2O7, Ti5P4O20, AlPO4 -0.065 LAGP Ge5O(PO4)6, GeP2O7, AlPO4, O2 -0.056 LISICON Zn2GeO4, GeO2, O2 -0.57 Oxide solid electrolyte materials have higher oxidation potential than sulfides. The oxidation of LLZO, LISICON, and LLTO starts at 2.91, 3.39, and 3.71 V, respectively. The NASICON materials, LATP and LAGP, show the best resistance to oxidation with the highest oxidation potential of 4.21 and 4.28 V, respectively (Fig. 3.1), and the lowest decomposition energy of only ~ -0.06 eV/atom at 5 V (Table 3.2). The delithiation reactions continue at higher voltages, and the O2 gas release during the oxidation at high voltages for all oxide solid electrolytes (Table 3.2). The oxidation of these solid electrolyte materials is 26 not surprising, given that Li2O is oxidized at 2.91 V and that the O2 gas is released by the further oxidation of Li2O2. LiPON starts oxidation at 2.63 V with the N2 gas release. Our computation results are consistent with the experiments by Yu et al.12, in which the onset of LiPON oxidation at ~2.6 V in the I-V measurements and the micro-sized gas bubbles in the LiPON material were observed after applying a high voltage of 6 V. A significant overpotential to the calculated thermodynamic equilibrium potential is expected for the oxidation reaction processes, which are likely to have a slow kinetics. The kinetic limitations of the oxidation reactions may come from multiple aspects. Most decomposition products at high voltages (Table 3.2) are electronic insulating, and the diffusion of non-Li elements is usually slow in solids. Furthermore, the nucleation and release of O2 and N2 gas molecules are likely to have sluggish kinetics. For example, a significant overpotential of > 1 V is often observed in the oxygen evolution reactions in metal-air batteries.59 Therefore, the overpotential of the decomposition reactions may provide a higher nominal oxidation potential of > 5 V and a wider nominal electrochemical window observed in the CV experiments.1,9,11,12 3.3 Chemical Stability of the Solid Electrolyte against non-Li Elements In addition to the equilibrium of Li (and ϕ) evaluated in section 3.1-3.2, the equilibrium of non-Li elements such as Co, O and S should also be achieved between the solid electrolyte and the electrode, though non-Li elements are usually less mobile than Li/Li+ in the solid electrolyte and electrode materials. The equilibria evaluated in this section correspond to chemical stability instead of the electrochemical stability, although the same scheme is used to evaluate electrochemical stability as in section 3.1 27 and 3.2. The µO stability window of LCO and L0.5CO electrode materials (Fig. 3.2a) corresponds to the range of possible µO values in the cathode material at the discharged and charged states, respectively. The oxide SEs, such as LLZO, LLTO, and LATP, have the µO window overlapping with those of LCO and L0.5CO (Fig. 3.2a). This overlapping of the µO window suggests that a common value of µO could be achieved across the interface of solid electrolyte and LCO/L0.5CO materials where the equilibrium µO is in the stable window for both materials. Table 3.3. Phase equilibria of the solid electrolyte and LiCoO2(LCO) /Li0.5CoO2(L0.5CO) materials at different chemical potential of O and S, respectively. anion µ Decomposition ΔE openD Corresponding environment (eV) products (eV/atom) LGPS µO = -0.73 O-poor limit of L0.5CO GeP2O7, SO3, Li2SO4 -3.68 µO = -2.64 O-poor limit of LCO LiPO3, S, GeS2, Li2SO4 -0.28 µO = -3.00 Oxidation onset Li4GeS4, S, Li3PO4 -0.02 Li3PS4 µO = -0.73 O-poor limit of L0.5CO LiPO3, SO3, Li2SO4 -3.73 µO = -2.64 O-poor limit of LCO LiPO3, S, Li2SO4 -0.24 µO = -3.00 Oxidation onset Li3PO4, S 0 LiPON0.14 µO = -0.73 O-poor limit of L0.5CO LiNO3, Li3PO4, Li4P2O7 -0.09 µO = -2.64 O-poor limit of LCO N2, Li3PO4, Li4P2O7 -0.02 µO = -4.08 Oxidation onset Li2PO2N, Li3PO4, LiN3 0 LiPON0.46 µO = -0.73 O-poor limit of L0.5CO LiNO3, Li3PO4, Li4P2O7 -0.36 µO = -2.64 O-poor limit of LCO N2, Li3PO4, Li4P2O7 -0.11 µO = -4.08 Oxidation onset Li2PO2N, Li3PO4, LiN3 0 LCO µS = -1.24 S-poor limit of LGPS Co9S8, Li2SO4 -0.28 µS = -2.72 Decomposition onset of LCO CoO, Li6CoO4, Li2SO4 0 L0.5CO µS = -1.24 S-poor limit of LGPS Co9S8, CoSO4, Li2SO4 -0.40 µS = -7.35 Decomposition onset of L0.5CO Li2SO4, Co3O4, LiCoO2 -0.03 28 Figure 3.2. The stability window of solid electrolyte and electrode LiCoO2/Li0.5CoO2 (LCO/Li0.5CO)with respect to the chemical potential of (a) O, (b) S, and (c) Co. In contrast, the µO windows of the sulfide solid electrolytes, LGPS and Li3PS4, have a significant gap with those of LCO and L0.5CO (Fig. 3.2a). Therefore, no common value of µO at the interface can simultaneously satisfy the µO equilibrium criterion between sulfide SEs and LCO materials. The equilibrium µO would be beyond the stability window of one or both materials, which would decompose simultaneously as a result. The µS stability window of the sulfide solid electrolytes does not overlap with that of LCO/L0.5CO either (Fig. 3.2b). These gaps of µO and µS windows suggest that the sulfide SE–LCO interfaces cannot satisfy the criteria (criterion 2 in Chapter 2) for the equilibrium of S and O across the interface. The high µO of the LCO/L0.5CO tends to oxidize LGPS and Li3PS4 into phases including Li3PO4 and Li2SO4, and the high µS of sulfide solid electrolytes tends to reduce LCO into phases including cobalt sulfides (Fig. 3.2 and Table 3.3). The reactions between sulfide electrolytes and LCO cathode material are highly thermodynamically favorable (Table 3.3), and the 29 formation of interphase layer has been reported in several experimental studies.46,60 The observed distribution of Co and S across the interface, which was previously interpreted as the mutual diffusion of S and Co,46 is likely the formation of cobalt sulfides interphase layer. Similar to sulfides, LiPON also has a µO window gap with LCO (Fig. 3.2a), but its decomposition energy is significantly smaller than sulfide solid electrolytes (Table 3.3). The µCo stability windows of LCO and L0.5CO overlap with all solid electrolytes investigated, suggesting that the equilibrium of Co can be achieved at the interface without going beyond the µCo window of the SE or LCO materials. This result suggests that the incompatibility between sulfide SE and LCO are mostly originated from the discrepancy of O and S anion chemistries rather than that of Co. The distribution and mutual diffusion of Co across the interfaces observed in the previous EDX experiments46 are likely due to the formation of decomposition interphases such as cobalt sulfides. 3.4 Chemical Stability of the Solid Electrolyte-cathode Interfaces The equilibria with respect to only one element evaluated in above sections correspond to physical situations where only one mobile Li reaches equilibrium. Full thermodynamic equilibria evaluated in this section allow the simultaneous equilibria with respect to all elements and describe the chemical stability of the SE–electrode interfaces. The chemical stability of the interface evaluated in this section is about whether two materials would react exothermically without any applied voltage. This chemical stability of the interface is important for the heat treatment and sintering during the cell preparation. 30 Table 3.4. The phase equilibria and decomposition energies of the SE–LCO and SE-L0.5CO interfaces. ΔED,min,mutual ΔED,min,total CSE CCathode xm Phase equilibria at xm (meV/atom) (meV/atom) LGPS LCO 0.42 Co9S8, Li2S, Li2SO4, -340 -349 Li3PO4, Li4GeO4 L0.5CO 0.40 Co9S8, Li2S, Li2SO4, -499 -527 Li3PO4, Li4GeO4 Li3PS4 LCO 0.41 Co9S8, Li2S, Li2SO4, -405 -405 Li3PO4 L0.5CO 0.39 Co9S8, Li2S, Li2SO4, -564 -584 Li3PO4 LiPON0.14 LCO 0.93 CoN, Li3PO4, Li6CoO4, N2 -35 -35 L0.5CO 0.59 Co3O4, LCO, LiNO3, -65 -65 Li3PO4 LiPON0.46 LCO 0.81 CoN, Li2O, Li3PO4 -99 -99 L0.5CO 0.62 CoO, LCO, Li3PO4, N2 -154 -154 LLZO LCO 0.96 La2O3, Li6Zr2O7, Li5CoO4 -1 -8 L0.5CO 0.47 La2O3, La2Zr2O7, -39 -60 Li7Co5O12, O2 LLTO LCO 0.64 Co3O4, La2Ti2O7, Li2TiO3, -0.5 -44 L0.5CO L0.5CO - LLTO, L0.5CO (stable) 0 - LATP LCO 0.32 L0.5CO, Co3O4, Li3PO4, -53 -63 LiAl5O8, TiO2 L0.5CO - LATP, L0.5CO (stable) 0 - Our calculations found that the interfaces between sulfide solid electrolytes and LCO/L0.5CO cathode are not thermodynamically stable (Table 3.4 and Fig. 3.3). The mutual reactions of LGPS with LCO and L0.5CO have favorable decomposition energies ΔED,min,mutual of -340 and -499 meV/atom, respectively. Similarly, the decomposition 31 energies of Li3PS4 with LCO and L0.5CO are -405 and -564 meV/atom, respectively. The phase equilibria of both interfaces include the formation of Co9S8, Li2SO4, and Li3PO4. The formation of cobalt sulfides at the interface between LCO and Li2S-P2S5 electrolyte was reported in the experiments.46 This reaction corresponds to a valence change of Co from 3+ to 2+. The cobalt sulfide phases with lower valences are known to be electronically conductive,61 which are detrimental to the stability of the interfaces.15 In addition, the mutual reaction energies for both LGPS and Li3PS4 with L0.5CO are lower than those with LCO, indicating larger thermodynamic driving force for the interfacial decomposition reactions at the charged state of the battery. Figure 3.3. Calculated mutual reaction energy ΔED,mutual of SE–LCO (solid lines) and SE- L0.5CO (dashed lines) interfaces. The mutual reaction energy of LLZO–LCO, LLTO–LCO, LLTO-L0.5CO, and LATP-L0.5CO have zero or near-zero values (minimum values provided in Table 4), which overlap at ΔED,mutual = 0. 32 The interfaces between LCO cathode and oxide solid electrolytes have significantly better chemical stability with ΔED,min,mutual of zero to tens of meV/atom compared to the sulfide electrolytes (Table 3.4 and Fig. 3.3). For example, the minimum reaction energy ΔED,min,mutual of garnet LLZO with LCO and L0.5CO are only -1 and -39 meV/atom, respectively. In addition, LLTO and LATP (Table 3.3 and Fig. 3.4) are thermodynamically stable against the charged state cathode L0.5CO, and their interfaces with LCO have small decomposition energies ΔED,min,mutual of only -1 and -53 meV/atom, respectively. The interfacial decomposition reactions may be kinetically inhibited. The good stability of LLTO and LLZO against LCO has been observed in the experiments.42,62 However, the sintering of the interfaces at high temperature may enhance the formation of the interphase layers.63 The main decomposition products of LLZO-L0.5CO interface include La2O3, La2Zr2O7, O2 (Table 3.4), which are similar to the delithiation products of garnet LLZO at high voltage (Table 3.2). The decomposition of the LLZO-L0.5CO interface is mainly due to the delithiation of the LLZO by L0.5CO. Our predicted phase equilibria at the LATP–LCO interface includes Co3O4, which is also observed as a reaction product at LCO-LATP interface in an experimental study.64 The formation of Co3O4 and L0.5CO corresponds to the Co valence increase from 3+ to 4+. The increase of Co valence at oxide solid electrolyte interfaces is opposite from that at the sulfide solid electrolyte interfaces, indicating the different nature of interfacial decompositions and the resulting properties between oxide and sulfide solid electrolytes. In addition, the decomposition interphase layers including Li3PO4 and LiAl5O8 can potentially passivate the interface and provide decent Li ionic conductivity across the interface.65 In summary, the oxide solid electrolyte materials, LLZO, LLTO, and LATP, have significantly better chemical compatibility with 33 the LCO cathode materials compared to sulfide solid electrolytes. LiPON also shows thermodynamically favorable reactions with LCO and L0.5CO (Table 3.4 and Fig. 3.3). The phase equilibria and decomposition energy are highly dependent on the LiPON composition. At low N content of 0.14, the mutual reaction energy ΔED,min,mutual of LiPON is only -35 and -65 meV/atom against LCO and L0.5CO, respectively. At higher nitrogen content of 0.46, the minimum mutual reaction energies with LCO and L0.5CO of -106 and -153 meV/atom, respectively, indicate more favorable decompositions. The LCO chemical compatibility of LiPON is better than sulfide solid electrolytes but worse than oxide solid electrolytes. The formation of an interphase layer and the change of chemical structures at the LiPON/LiCoO2 interfaces were observed by in situ XPS experiments.66 The good compatibility of the LiPON with the LCO cathode material observed in the experiments67-70 may be explained by the effect of the interphase layers, including Li3PO4, which is a well-known coating material for cathode71-73 and is the dominant decomposition product at the LiPON interfaces. 3.5 Electrochemical Stability of the Solid Electrolyte-Cathode Interfaces In this section, we evaluated the phase equilibria of the solid electrolyte–LCO interfaces at the applied potential ϕ on the basis of the equilibrium criterion 4 (section 2.5). These phase equilibria correspond to the interphase evolution, such as lithiation or delithiation, in response to the applied potential, and the interfacial decomposition described by the reaction energies may become more favorable at certain applied potentials. The interface between LCO and Li3PS4 was found to have poor stability over the entire range of the applied voltage from 2 to 5 V. The interfacial mutual reaction energy ∆𝐸01,2#,W;2,WXAXCG is 34 in the range of [-737, -594] meV/atom for LPS from 2 to 5 V (Fig. 3.4 and Table 3.5). The total decomposition energy Δ𝐸01,2#,W;2,A0ACG of LGPS or Li3PS4 reaches -1.25 eV/atom at 5 V (Fig. 3.4), suggesting highly favorable decompositions at high voltages. Previous experimental studies also reported the formation of interphases of tens nanometers including cobalt sulfides between LCO and Li2S-P2S5 SE after charging.46 The growing differences between ∆𝐸01,2#,W;2,WXAXCG and Δ𝐸 01,2 #,W;2,A0ACG at higher voltages (Fig. 3.4) are mainly due to the increasing contribution from the delithiation of the SEs to the total decomposition energy Δ𝐸01,2#,W;2,A0ACG. The oxide solid electrolytes–LCO interfaces generally have significantly better stability than the sulfide SEs during the whole voltage range. For example, the LLZO–LCO interface has the minimum interfacial mutual reaction energy ∆𝐸01,2#,W;2,WXAXCG of only -33 meV/atom (Table 3.5). At high voltage above 4 V, the reaction products, La2Zr2O7 and LaCoO3, are likely poor Li ionic conductor materials. The predicted phase LaCoO3 was observed experimentally for the LLZO–LCO interface after the heat treatment of over 1000 °C,74 which causes Li loss similar to the delithiation at high voltage. In addition, LaCoO3 may be electronic conductive,75 leading to the formation of thick interphase layers and potential degradation at the interface. However, the final phase equilibria of the interface at high voltage involve O2 gas release, which is likely to have poor kinetics as in the oxygen evolution reactions in Li-air batteries. A significant amount of overpotential, yielding high oxidation potential observed in many CV experiments, is expected for such oxidation reactions. Similarly, the LCO interfaces with LLTO and LATP show good stability with zero mutual reaction energy ∆𝐸01,2#,W;2,WXAXCG at voltage higher than 3.34 V and 4.53 V, 35 respectively (Fig. 3.4 and Appendix). The good electrochemical stability of the interfaces between LLTO/LATP and LCO at high voltages may partly due to high oxidation potential (anodic limit) of these solid electrolytes and the good chemical stability between the solid electrolytes and LCO cathode. Figure 3.4. The mutual reaction energy Δ𝐸01,2#,W;2,WXAXCG (solid lines) and the total reaction energy Δ𝐸01,2#,W;2,A0ACG (dashed lines) at the SE–LCO interfaces under applied potential ϕ in a 2 V – 5 V range. The mutual reaction energy of the LiPON0.14–LCO interface is comparable to oxide SEs, though LiPON0.14 reacts favorably with LCO in the whole 2 V- 5 V range (Fig. 4 and Table 5). The formation of Li3PO4 interphase and the small decomposition energy may explain the widely observed stability of LiPON with LCO.67-69,76 The interfacial reaction between LiPON and LCO also involves N2 and O2 gas release at above 4.23 V, respectively. The void formation at LiPON–LiCoO2 interface after battery cycling was reported in an in- situ TEM experimental study.69 36 Table 3.5. Phase equilibria and decomposition energy of SE–LCO interfaces under applied potential ϕ. Applied Phase equilibria at x = xm(ϕ) Δ𝐸 01,2 01,2 #,W;2,WXAXCG Δ𝐸#,W;2,A0ACG Interface potential ϕ (V) under ϕ (meV/atom) (meV/atom) Li3PS4– 2.00 - 2.91 Co3S4, Li2SO4, Li3PO4 [-641, -596] [-737, -596] LCO 2.91 - 2.99 Co3S4, Li2SO4, Li4P2O7 [-596, -594] [-750, -737] 2.99 - 3.20 Co3S4, Li2SO4, LiCoPO4 [-607, -594] [-779, -730] 3.20 - 3.26 Co3S4, Li2SO4, Co3(PO4)2 [-613, -608] [-785, -773] 3.26 - 5.00 Co3S4, CoSO4, Co3(PO4)2 [-737, -616] [-1273, -763] LLZO– 2.00 - 2.57 La2O3, Li6Zr2O7, Li5CoO4 [-2, -2] [-9, -9] LCO 2.57 - 2.81 La2O3, Li6Zr2O7, Li7Co5O12 [-7, -2] [-13, -8] 2.81 - 3.50 La2O3, La2Zr2O7, Li7Co5O12 [-33, -2] [-72, -11] 3.50 - 3.99 La2O3, La2Zr2O7, O2, L0.5CO 0 [-137, -53] 3.99 - 5.00 LaCoO3, La2Zr2O7, O2 [-5, 0] [-656, -284] LiPON0.14– 2.00 - 2.12 CoN, N2, Li3PO4 [-53, -53] [-55, -55] LCO 2.12 - 2.70 CoO, N2, Li3PO4 [-64, -53] [-70, -55] 2.70 - 3.01 Co3O4, N2, Li3PO4 [-59, -53] [-84, -70] 3.01 - 3.89 Co3O4, LiNO3, Li3PO4 [-67, -35] [-139, -77] 3.89 - 4.16 LiCoPO4, LiNO3, Li3PO4 [-33, -11] [-190, -165] 4.16 - 4.18 LiCoPO4, LiNO3, Li4P2O7 [-9, -9] [-221, -221] 4.19 - 4.23 CoPO4, LiNO3, Li4P2O7 [-9, -8] [-221, -215] 4.23 - 4.54 CoPO4, LiNO3, O2 [-53, -9] [-288, -169] 4.54 - 5.00 CoPO4, Co(NO3)4, O2 [-116, -56] [-482, -294] 37 Our results show that the stability of the interface and the formed phase equilibria are dependent on the applied potential, as the delithiation at high voltages provides additional thermodynamic driving force for the interfacial decompositions. Among the SE– LCO interfaces investigated, the interfaces of the SEs (such as oxides) that have good electrochemical window and good chemical stability with the electrode in general show better stability during electrochemical cycling. Therefore, the SE with a wide electrochemical window and the good chemical stability with electrodes are desired to achieve intrinsic interface stability during electrochemical cycling. For those SE (such as sulfides) and interfaces that cannot satisfy the above criteria, an interfacial coating layer material can be applied to resolve the stability problems at the interface. 3.6 Improved Interfacial Stability by Applying Coating Layers Currently, the interfacial resistance has become a critical problem for the performance of all-solid-state Li-ion batteries. Engineering of the interface, such as the application of interfacial coating layers, is used to improve interfacial protection and to reduce interface resistance. In this section, we investigated the electrochemical stability of the coating layer materials, such as Li4Ti5O12,77,78 LiTaO3,79 LiNbO3,80,81 Li2SiO3,82 and Li3PO4,83 which were demonstrated to suppress the mutual diffusion of non-Li elements and to reduce the interfacial resistance at the solid electrolyte-cathode interfaces in all-solid-state Li-ion batteries.2,44,46,79 38 Solid Cathode / Electrolyte Oxidant μ~Li+ μLi Solid electrolyte oxidation Oxidation potential μ~e- (a) (b) Figure 3.5. (a) Electrochemical stability window (color bars) of commonly used coating layer materials. The oxidation potential to fully delithiate the material is marked by the dash line. The blue solid line at 3.9 V represents the equilibrium voltage of the LiCoO2 cathode material. (b) Schematic diagram of Li chemical potentials µLi (black line) and the electrochemical potential 𝜇e:;f (blue dashed line) and 𝜇e,g (red dashed line) across the solid electrolyte-cathode interface. Our calculations show that these coating layer materials have an electrochemical window from the reduction potential of 0.7-1.7 V to the oxidation potential of 3.7-4.2 V (Fig. 3.5a). Therefore, the coating layer materials are stable between 2 V to 4 V, the usual voltage range during the cycling of Li-ion batteries. In addition, the coating layers materials have poor electronic conductivity and can serve as artificial SEIs to passivate the solid electrolyte (Fig. 3.5b). Given that the sulfide solid electrolyte materials are oxidized at as low as 2.5 V and are not thermodynamically stable at the voltage of 4 V, the coating layers serve as critical passivation through the passivation mechanism. The coating layers mitigate the low Li chemical potential µLi from the cathode material applied on the solid electrolyte 39 Higher ϕ / Higher μ Lower μ Coating Layer materials. As a result, the oxidation and delithiation of the solid electrolyte at the cathode interface is stopped, and the oxidation potential (anodic limit) of the solid electrolyte is extended by the artificial coating layer. Therefore, the coating layer effectively extended the anodic limit of the sulfide solid electrolyte from ~2.5 V to ~4 V. The overpotential to oxidize the coating layers may further extend the nominal stability window. Similar strategy of applying artificial coating layers has been employed at the anode side for the protection and stabilization of Li metal anode. For example, Polyplus84 has applied coating layers between Li metal and LATP electrolyte to protect the LATP materials against Li metal. The passivation mechanism of the coating layer at the anode side is the same as the decomposition interphase demonstrated in section 3.1. Table 3.6. The mutual reaction energy ΔED,min,mutual (in meV/atom) of the coating layer materials with the SE or LCO materials. Li3PS4 LLZO LCO L0.5CO Li4Ti5O12 -80 -75 -1 0 LiNbO3 -155 -76 0 0 Li2SiO3 -19 -29 0 0 LiTaO3 -49 -68 0 0 Li3PO4 0 0 0 0 The coating layer material lying between the original SE–LCO interfaces forms two new interfaces with the SE and LCO, respectively. By applying the same computational scheme as in section 2.4, we investigated interfacial chemical compatibility of these two interfaces with the coating layer. The previously demonstrated coating layer materials, such as Li4Ti5O12, LiTaO3, LiNbO3, Li2SiO3, and Li3PO4, have excellent chemical stability with 40 the LCO and L0.5CO with zero or negligible decomposition energy ΔED,min,mutual (Table 3.6). In addition, all coating layer materials show relatively better stability with the sulfide SE comparing to the original sulfide SE–LCO interfaces, which have the interfacial reaction energy ΔED,min,mutual of ~ -500 meV/atom. As a result, the interface with the coating layer has significantly improved stability and suppresses the formation of thick interphase layer. This result of the coating layer was observed as reduced mutual diffusion across the interface in the experiments.77,79,82 In addition to stabilizing the interface, the coating layer of only a few nanometers is significantly thinner than the decomposition interphase layer of ~10 to 100 nm. The reduced thickness of the coating layer may significantly reduce the high interfacial resistance caused by the thick decomposition interphase layer. Figure 3.6. Electrochemical window of potential coating layer materials. The dashed box marks the equilibrium voltage to fully delithiate the materials. 41 Table 3.7. The mutual reaction energy Δ𝐸#,W;2,WXAXCG (in meV/atom) of the coating layer materials with the SE or LCO materials. Li3PS4 LLZO LCO L0.5CO Li4TiO4 -125 0 0 -30 Li2TiO3 -75 -5 0 0 Li8Nb2O9 -147 0 0 -20 Li3NbO4 -132 -4 0 0 LiNb3O8 -173 -115 -16 0 Li8SiO6 -177 0 -3 -50 Li4SiO4 -81 -1 0 -12 Li2Si2O5 -10 -69 -4 0 Li5TaO5 -117 0 0 -32 Li3TaO4 -64 -3 0 0 LiTa3O8 -64 -105 -22 0 Li4P2O7 -9 -101 -44 -3 LiPO3 -32 -201 -76 -19 In addition, we also computationally investigated other compounds (listed in Fig. 3.6 and Table 3.7) based on the same cations, such as Ti, Nb, Si, Ta, and P, as potential coating layer materials. All these lithium metal oxide materials have a wide electrochemical window (Fig. 3.6). The compounds with higher Li content generally show lower reduction potential, and the compounds with lower Li content or higher O content show higher oxidation potential. Most of these compounds have excellent chemical stability against LCO 42 and L0.5CO cathode materials with zero or small Δ𝐸#,W;2,WXAXCG (Table 3.7). All these coating layer materials significantly improve the stability of sulfide SE– LCO interfaces. The chemical stability between LLZO–coating layer interface varies significantly with the compositions of the coating layer materials. It is interesting to note that the coating layers that are the most stable with sulfide SE may not necessarily be the most stable ones with LLZO. For example, LiNbO3, a demonstrated coating layer for sulfide SEs, is not as stable as Li3NbO4 and Li8Nb2O9 with LLZO according to the decomposition energy Δ𝐸#,W;2,WXAXCG (Table 3.7). In the previous study, the application of Nb at the LLZO–LCO interfaces may form amorphous lithium niobates including Li8Nb2O9 or Li3NbO4-like phases,81 which stabilize the LLZO–LCO interface and hence reduce the interfacial resistance. Similarly, the other Li-rich coating layers, such as Li4TiO4, Li2TiO3, Li8SiO6, Li4SiO4, Li5TaO5, and Li3TaO4 may work better with LLZO than those previously demonstrated for sulfide SEs, such as Li4Ti5O12, LiTaO3, and Li2SiO3. 3.7 Critical Role of Interphases, Interface Types and Interfacial Engineering Strategies Our thermodynamic analyses based on first principles calculations indicate that most solid electrolyte materials have limited electrochemical window. In contrast to the widely held perception about the outstanding stability of the solid electrolyte materials, the solid electrolyte materials are reduced and oxidized at low and high potentials, respectively, and are not thermodynamically stable against Li metal. The sulfide solid electrolytes based on thio-phosphates are reduced at ~1.6-1.7 V and oxidized at ~2.5 V. The stability window of oxide solid electrolytes varies greatly from one material to another. Although some 43 oxides have high reduction potential as sulfides, most oxide solid electrolytes have a significantly higher oxidation potential, and are not oxidized until > 3 V. In particular, the NASICON materials, LATP and LAGP, are thermodynamically stable up to ~4.2 V. Among all these oxides investigated, the Li garnet materials have the best resistance to Li reduction. Overall, the oxide solid electrolyte materials have significantly wider electrochemical window than sulfides. The reduction and oxidation potentials as well as the decomposition products of solid electrolytes predicted from our calculations are in good agreement with prior experimental studies, confirming that our computation method based on the Li grand potential phase diagram is a valid scheme in evaluating the electrochemical stability of materials. Our calculation results demonstrated that the good stability of the solid electrolyte materials is originated from the kinetic stabilizations. First, the wide, nominal electrochemical window observed in many CV experiments can be partially attributed to the significant overpotential of the sluggish kinetics during the decomposition reactions (Fig. 3.7). The decomposition reactions though kinetically sluggish are still thermodynamically favorable at the applied overpotential and may happen over an extended period of time, leading to the deterioration of the batteries. This kinetic stabilization from the sluggish kinetics of the reactions is different from the passivation mechanisms. The passivation mechanism of the interphases is the origin of the outstanding stability in the solid electrolyte. The decomposition interphases with good stability and poor electronic transport are effectively the SEIs in the all-solid-state Li-ion batteries to passivate the solid electrolytes (Fig. 3.7). The interphases, which are stable against solid electrolytes and electrodes, mitigate the Li chemical potential discrepancy between the electrolyte and 44 electrode at the interfaces. As a result, the anodic/cathode limits and the electrochemical window of the solid electrolyte are significantly extended by the extra electrochemical window provided by the interphases (Fig. 3.6). The effective electrochemical window of the solid electrolyte materials is its own intrinsic electrochemical window plus the electrochemical window of the interphases (Fig. 3.7). Higher ϕ / Lower μLi Overpotential ϕ=5V Interphase/ Cathode coating Intrinsic Oxidation potential Electrochemical Window Solid Extendedelectrolyte Electrochemical Reduction potential Window Interphase Anode ϕ=0V Figure 3.7. Schematic diagram of the electrochemical window (color bars) and the Li chemical potential profile (black line) in all-solid-state Li-ion battery. The profile of chemical potential is schematic in this plot and may not be linear. The high µLi in anode (silver) and low µLi in cathode (blue) are beyond the stability window of the solid electrolyte (green). The observed nominal electrochemical window is extended by the overpotential (dash line) and by the interphases (orange and yellow), which account for the gap of µLi between solid electrolyte and electrodes across the interfaces. Our computational study revealed that the electrochemical and chemical stabilities of the solid electrolyte–electrode interfaces in the ASLiBs are intrinsically limited. Most solid electrolyte–electrode interfaces are not thermodynamically stable before, after, and during electrochemical cycling. In particular, the sulfide solid electrolytes are neither 45 chemically nor electrochemically stable against the LCO cathode, due to the anion chemical potential discrepancy between sulfide solid electrolytes and LCO. The oxide solid electrolytes have significantly better chemical and electrochemical stabilities with the LCO cathode than sulfide SEs. The interfacial reactions between oxide solid electrolyte and LCO, though may be kinetically limited, are also thermodynamically favorable, suggesting potential interfacial degradations after thermal processing or during battery cycling. Type 1 Type 2 Type 3 Interface Interface Interface Solid Electrolyte Solid Electrolyte Solid Electrolyte Li+ + -No interphase Li MIEC e Li+ e- Interphase SEI Electrode Electrode Electrode Figure 3.8. Schematic illustration of the three types of solid electrolyte–electrode interfaces. Type 1: ideal stable interface; Type 2: interface formed with interphase layers that are mixed ionic and electronic conductors (MIEC); Type 3: interfaces formed with stable solid- electrolyte interphase (SEI) These results about the interfacial thermodynamics suggest the ubiquitous formation of the interphase layers at the solid electrolyte–electrode interfaces, which have significant implications for the research and development of the ASLiBs. The properties of these interphase layers would be critical to the performance of the ASLiBs. In particular, the distinctive chemistry nature of sulfide and oxide solid electrolytes lead to very different interphase properties for the SE–LCO cathode interfaces and hence require different interfacial engineering strategies. Following the definitions by Wenzel et al.,18 we 46 distinguish three different types of interfaces based on the interface stability and interphase properties. (Fig. 3.8) Type 1 interface - Stable interfaces with no decomposition or interphase layers. Type 1 interfaces are either thermodynamically intrinsic stable or kinetically stabilized during the cycling voltage range. Type 1 interfaces do not have interphase layer between the solid electrolyte and electrode, and is expected to have decent interfacial ionic conductivity, because both the solid electrolyte and the electrode are good Li ionic conductors. However, our thermodynamic analyses demonstrated that few solid electrolyte materials have thermodynamically intrinsic stabilities against Li metal anode or high- voltage cathodes, or during the entire range of cycling voltages. Among five classes of SE materials investigated, LLTO and LATP have best electrochemical stability against LCO cathode materials at high voltages, and LLZO has the best stability against Li metal. Some oxide SEs may form kinetically stabilized Type 1 interfaces with the electrode materials,42,62 though these Type 1 interfaces may still degrade and convert to other types during sintering or electrochemical cycling due to the limited thermodynamic stability. Type 1 interface requires the SEs with wide electrochemical window and good chemical compatibility with the electrode materials. Since the electrochemical window of the solid electrolytes and the chemical compatibility of the SE–electrode interfaces are limited, it is unrealistic to have a solid electrolyte with both Type 1 interfaces with Li metal anode and high-voltage cathode. Type 2 interface - Interfaces formed with interphase layers that are mixed ionic and electronic conductors (MIEC). The formation of the interphase layer between the solid electrolyte and the electrode is expected for most Li-SE and SE–LCO interfaces as suggested by the interfacial 47 thermodynamics results from our computation. The electronic conductivity of the interphase, regardless of specific electronic conductive mechanisms, determines whether the interface is Type 2 or Type 3. If the formed interphase is MIEC, the simultaneous transport of Li ions and electrons would enable continued decomposition reactions, which are thermodynamically favorable. Therefore, the MIEC interphases cannot provide essential passivation at the interface.15 The examples of Type 2 interfaces include the Li-LGPS and Li-LLTO interfaces, which formed interphases including the electron-conducting Li-Ge alloys and lithium titanates, respectively. The formation of Type 2 interfaces in these materials explains the Li reduction of the LGPS and LLTO materials observed in the experiments.14,85 The continued decomposition may result in poor coulombic efficiency in ASLiBs. Therefore, Type 2 interfaces, which may cause continued interfacial degradation and high interfacial resistance, should be avoided in the ASLiBs. The application of artificial coating layers can be used to engineer Type 2 interfaces into Type 3 interfaces. Type 3 interface - Interfaces formed with stable solid-electrolyte interphase (SEI). In contrast to Type 2 interfaces, the interphase layers in Type 3 interfaces are electronic insulating and are stable during the electrochemical cycling, acting as a stable SEI at the SE–electrode interfaces. Type 3 interfaces with spontaneously formed SEIs have self- limiting interfacial degradation. The SEIs passivate the SE materials and stop the continued decomposition into the bulk of the solid electrolyte.15 The Li-LiPON interface is Type 3, where the formed interphases containing Li2O, Li3N, and Li3P are electronic insulating.15,45 In addition, the interphase decomposition layer, including fast Li-ion conducting Li3N and Li3P,41,86 is likely to yield low interfacial resistance. The interfacial resistance of Type 3 interfaces is highly dependent on the ionic conductivity of the SEIs. The Li-LiPON interface 48 is an ideal Type 3 interfaces which have spontaneously formed SEIs with excellent ionic conductivity. For those Type 3 interfaces with poor Li-ion conducting SEIs, the application of a coating layer with good ionic conductivity as the artificial SEI may resolve the high resistance problem of the spontaneously formed SEIs. We summarize the categorization of different combination of SE and electrode materials into the defined interface types (Fig. 3.8). Our discussion focuses on Li metal anode and LCO cathode materials (similar results are expected for other transition metal oxide cathodes), which are often desired for the ASLiBs. For Li-SE interfaces, Type 1 interfaces are rare due to the strong thermodynamic driving force for Li metal to reduce the solid electrolyte. The Li-LLZO interface may be Type 1 as a result of kinetically stabilization and small thermodynamic driving force for Li reduction. The high temperature treatment may facilitate the Li reduction of LLZO and formation of the interphase.58 The Li-LiPON and Li-Li3PS4 are typical Type 3 interfaces, where the passivation layers of high Li ionic conductivity and poor electronic conductivity are formed. The Li-LGPS, Li-LLTO, and Li-LATP are typical Type 2 interfaces, where significant reduction of the SE is observed. Cations, such as Ti and Ge, in SEs would facilitate the formation of the Type 2 interfaces, because of the formation of the electronic conducting interphase layers at Li reduction. Therefore, the use of these cations should be avoided in the design of SE materials for Li metal compatibility. In contrast, anion mixing should be a viable strategy for the SE materials design to simultaneously improve stability and Li ionic conductivity. LiPON is a successful case of using oxynitride to achieve good stability and Li ionic conductivity. Similar successes have been demonstrated in the doping of Li3PS4 with halides, such as LiCl 49 and LiI.54,87 Similarly, we expect some oxysulfide compounds as promising solid electrolytes. The interfaces between the solid electrolyte and cathode, such as LCO, may be more complicated, due to a large number of elements involved and a wide range of Li chemical potential at the charged and discharge states of the battery. The sulfide SE–LCO interfaces may also be partially Type 2, because the interphases include the cobalt sulfide binaries such as Co9S8 or Co3S4, which are electronically conductive.61,88 This formation of MIEC interphases may explain the large thickness of the interphase layers of 10-100 nanometers observed at these sulfide–LCO interfaces.46 Thick interphase layers generally result in a high interfacial resistance. However, the decomposition does not happen to the entire bulk of SEs due to the limited diffusion of Co and the drop of the Co content inside the SEs. The part of the interphase at the direct adjacent of solid electrolyte is of low content of Co sulfides and is still passivating (Type 3), which stops the further growth of the interphase layers. The variation of the interphase composition across the interface has been observed in our computation (Appendix) and the EDX experiments.60,82 Therefore, the sulfide SE-cathode interface is not strictly Type 2 (as in Li-LLTO) but rather a mixture of Type 2 and Type 3. Having some Type 2 features inside the interphase may have negative impact on the electrochemical performance of the ASLiBs. The enabled transport of both Li+ and e- through the MIEC interphase would facilitate the electrochemical cycling of these interphase layers during cycling voltages (as predicted in section 3.4). Active electrochemical cycling of these interphase layers may facilitate the growth of the decomposition interphase layers as a result of cyclic lithiation and delithation14, which lead to repetitive volume change and the failure of the mechanical contact at the interface. Such 50 special interface properties may explain the poor cyclability of the ASLiBs. Therefore, the key problem of sulfide SE–LCO interfaces is the poor stability, which leads to thick interphase layers, high interfacial resistance, and degradation over cycling. Therefore, the application of the coating layer at these interfaces is the corresponding strategy to address the stability issue by turning the interfaces into a desired Type 3 interface with thin thickness and improved interfacial conductivity. The coating layers serve as artificial SEIs to stabilize the interface and to resolve the issue of poor interfacial conductivity. In general, we found that the oxide SE and LCO interfaces may be Type 1 or Type 3. In particular, LLTO-L0.5CO and LATP-L0.5CO are thermodynamically stable as Type 1, and other oxide SE–LCO interfaces may be kinetically stabilized as Type 1. After the thermodynamically favorable decomposition between these materials, some oxide SE– LCO interfaces may turn into Type 3, because the interphases formed at oxides SE-cathode interfaces are mostly electronic insulating. However, such interphases and formed SEI layers in some Type 3 oxide solid electrolyte-cathode interfaces may have poor Li ionic conductivity, since most of the equilibrium phases have low or zero Li content as a result of the delithiation. For example, the spontaneous decomposition interphase, such as La2Zr2O7, of LLZO–LCO interface are likely poor Li ion conductor. For these Type 3 interfaces, the key issue is not poor interfacial stability but the low ionic conductivity of the formed interphase layers. Therefore, interfacial engineering for the aforementioned Type 3 interfaces is also necessary to improve interfacial ionic conductivity. The increased interfacial conductivity after applying lithium niobate interphase layer in the previous study81 may be due to the higher ionic conductivity of the coating layer than that of the spontaneous decomposition interphase. In addition, the soft and ductile Nb metal may 51 enhance the wetting and may promote the interfacial contact between the LLZO and LCO. As illustrated above, the categorization of different types of interfaces is critical to understand the problems at the interfaces and to apply corresponding interfacial engineering, which is the key to achieve desired interfacial properties and to improve electrochemical performance in ASLiBs. 52 Chapter 4. Strategies Based on Nitride Materials Chemistry to Stabilize Li Metal Anode Li metal has been long desired as the anode material with the highest theoretical specific capacity and the lowest standard potential to significantly increase the energy density in rechargeable Li-ion battery.89,90 Enabling Li metal anode has been regarded as the “Holy Grail” 91,92, but confronts many challenges. The undesired growth of Li dendrite during cell cycling causes short circuiting in the cell, leading to catastrophic cell failure and safety issues. In addition, Li metal is strongly reducing, and is not compatible with most electrolytes.93-95 The lack of long-term stability between electrolyte and Li metal anode results in low coulombic efficiency, capacity fading during cycling, and cell failure.94-96 To protect electrolytes against Li metal and to stabilize Li metal anode, electrolytes and additives are developed to form stable solid-electrolyte-interphase (SEI) layer on Li metal.96 Forming such SEI layers has led to the success of graphite anode in commercial Li-ion batteries.93,94,97,98 Besides the spontaneous formation of SEI, applying Li-stable protection materials as artificial SEI on Li metal is also demonstrated to significantly improve the cyclability and coulombic efficiency of the cells with Li metal anode.84,99-102 Recently, using solid electrolytes to assemble all-solid-state lithium-ion batteries is a promising direction to enable Li metal anode.90 For example, the solid electrolyte LiPON, an oxynitride material, was demonstrated to achieve a cycle life of over 10,000 cycles in thin-film lithium metal batteries.8 2 This chapter has been published in Y. Zhu, X. He, Y. Mo, Strategies based on nitride materials chemistry to stabilize Li metal anode. Advanced Science, 1600517, (2017) 53 To protect against the reduction of Li metal, protection coating materials or solid electrolytes that are thermodynamically stable or that form stable passivation layers against Li metal are needed. Many Li binaries such as LiF, Li2O, Li2S, Li3N, and Li3P, are thermodynamically stable against Li metal, and some are found in SEI layers.45,98,103,104 However, despite that many lithium solid electrolytes were previously reported to be stable against Li metal, multiple experimental and first principles studies confirmed the reduction of solid electrolytes, including Li10GeP2S12 (LGPS), NASICON-type Li1.3Al0.3Ti1.7(PO4)3 (LATP), and perovskite Li0.33La0.56TiO3 (LLTO).13-20 In these materials, Li metal reduces Ge and Ti cations, forming Li-Ge alloy and lithium titanate, respectively. Even for the well- demonstrated Li-compatible LiPON materials, recent computational and in situ experimental studies confirmed Li reduction, which leads to a spontaneously formed interphase layer consisting of Li3N, Li2O, and Li3P at Li-LiPON interface.15,45 This interphase layer is electronic insulating, passivating LiPON against further Li reduction. Therefore, forming passivating SEI layer against Li metal is critical to achieving good stability. Given the limited number of available materials stable against Li metal, the development of novel Li-stable protection materials or solid electrolytes is crucial to enable Li metal anode. However, the design and discovery of Li-stable materials are still mostly based on a trial-and-error approach, due to the lack of knowledge about Li metal stability of different materials chemistry and compositions. Li reduction behavior is only known for a limited number of well-studied materials. For example, the reduction of Ge and Ti is a known problem in LGPS and LATP, respectively, but it is not clear whether these cations will still be reduced in different compositions or in other anion chemistry. In addition, it is not clear 54 whether the excellent stability of LiPON is a general property of oxynitride chemistry. Such knowledge about Li stability will be crucial to guide further materials development to stabilize Li metal anode, and will enable guided engineering of materials chemistry to form stable SEI on lithium metal. In this chapter, we aim to fill this knowledge gap about Li metal stability and passivation behavior over a wide range of materials chemistry. Using a data-driven first principles computation approach based on a large-scale materials database, we studied Li metal stability and lithiation reactions in a wide range of materials chemistry and composition space. We determined the intrinsic thermodynamic stability of materials against Li metal as a function of cation and anion chemistry, and discovered that nitride anion chemistry uniquely exhibits better thermodynamic stability against Li metal compared to oxides, sulfides, and halides. In addition, the materials chemistry and composition range that can form stable passivation interphase against Li metal were identified. On the basis of newly obtained chemistry knowledge from computation, multiple novel strategies to form stable SEI on Li metal anode were proposed, providing opportunities for future research and development of lithium batteries. 4.1. General Trend of Materials Stability against Li Metal Reduction. We calculated the Li reduction potential (cathodic limit) of M-X binary compounds and Li-M-X ternary compounds in four anion chemistries, X = N, O, S, and F (Fig. 4.1), which are commonly used in solid electrolytes or coating layer materials. To illustrate our results, we first take Al-abiding compounds as examples. The binary materials AlF3, Al2S3 and Al2O3 have high cathodic limits of >1.2 V, and their lithiation reactions with Li metal 55 form Li9Al4 and Li-X binary materials with a reaction energy of -0.77, -0.74, and -0.23 eV per Li, respectively (Table 4.1). The same trend is also observed for ternary Li-Al-X (Appendix B), Li-Ge-X, and Li-P-X compounds (Table 4.1). Table 4.1. Cathodic limits and lithiation reactions for example fluorides, sulfides, oxides and nitrides. The reaction energy ED is normalized to per Li inserted. Phase equilibria Cathodic limit ref. Phase equilibria ED (eV) Compound at cathodic limit to Li/Li+ (V) with Li metal per Li AlN Li9Al4, Li3AlN2 -0.004 Li9Al4, Li3N *3 0.16 Al2O3 Al, LiAl5O8 1.23 Li9Al4, Li2O -0.23 Al2S3 Al, LiAlS2 1.60 Li9Al4, Li2S -0.74 AlF3 Al, Li3AlF6 1.29 Li9Al4, LiF -0.77 Li5GeN3 Li15Ge4, Li3N 0.20 Li15Ge4, Li3N -0.20 Li4GeO4 Ge, Li2O 1.02 Li15Ge4, Li2O -0.72 Li4GeS4 Ge, Li2S 1.62 Li15Ge4, Li2S -1.04 Li2GeF6 Ge, LiF 2.71 Li15Ge4, LiF -1.60 Li7PN4 Li3P, Li3N 0.01 Li3P, Li3N -0.01 Li3PO4 Li3P, Li2O 0.69 Li3P, Li2O -0.69 Li3PS4 P, Li2S 1.72 Li3P, Li2S -1.42 LiPF6 P, LiF 2.74 Li3P, LiF -2.06 These examples show that different anion chemistry exhibits different stability against Li metal. In general, fluorides have the highest cathodic limits, yielding poorest stability against Li metal. This is similar to the trend observed in organic liquid electrolytes, as fluorinated compounds usually show higher reduction potential.95 Oxides and sulfides * Since AlN is stable against Li metal, such phases are not phase equilibria with Li metal, but are fully lithiated products after lithiation with overpotential. 56 have lower cathodic limits than fluorides, and only a few of investigated oxides and sulfides are thermodynamically stable against Li metal. Figure 4.1. Cathodic limits (referenced to Li/Li+) for Li reduction of a) M-X binary compounds and b) Li-M-X ternary compounds in fluoride, sulfide, oxide, and nitride anion chemistries. Only cation M at its highest common valence state is considered. Each data point represents a compound. A full list of compounds and corresponding lithiation reactions are provided in the Appendix B. By contrast, nitrides show significant lower cathodic limits compared to other anion chemistry. For example, AlN shows a negative cathodic limit referenced to Li/Li+, and a positive reaction energy of 0.16 eV per Li to form Li9Al4 and Li3N after lithiation (Table 4.1). Therefore, it is thermodynamically unfavorable to lithiate AlN by Li metal. In addition, 57 Li3AlN2 also shows a negative cathodic limit of -0.32 V (Appendix), indicating its intrinsic stability against Li metal reduction. These trends are observed in both binary and ternary compounds with other cations. In general, many nitrides have negative reduction potential, and are thermodynamically stable against Li metal. Some of these nitrides that are electronic insulators and decent ionic conductors, e.g. Li3AlN2, Li3BN2, Li5SiN3,105 may be used as buffer layer materials to protect against Li metal. In summary, nitride anion chemistry shows unique electrochemical stability at low potentials and against Li metal. 4.2. Cation Effect on Lithium Metal Stability and Passivation Using Li-P-S and Li-Ge-S ternary systems as examples, we illustrate the reduction behavior of different cation chemistry and the effect of cation on Li stability and passivation. The Li-P-S composition space (Fig. 4.2a) includes many well-known solid electrolyte materials, such as Li3PS4, Li7P3S11, and Li2S-P2S5 glass. For example, the lithiation reaction of Li3PS4 (blue dashed line in Fig. 4.2a, and Fig. 4.2b) starts with the reduction of P at voltage 1.72 V (Table 4.1) and eventually leads to Li3P and Li2S, as 8 Li + Li3PS4 → Li3P + 4 Li2S (ΔH = -11.39 eV or -1099 kJ mol-1), which is highly thermodynamically favorable. Since Li3P and Li2S are the only stable phases against Li among entire Li-P-S composition space, any Li-P-S ternary compound would form the same phase equilibria after full lithiation (green bar in Fig. 4.2a), leading to the spontaneous formation of interphase consisting of Li3P and Li2S. Since both phases are electronic insulating, the formed interphase is passivating against further Li reduction.15,106 The formation of Li3P and Li2S interphase of Li7P3S11 on Li is confirmed by in situ X-ray photoelectron spectroscopy (XPS) experiments.107 58 (a) Li (b) Lithiation P2S5 S Li3P Li PS Li2S 3 4 Delithiation LiP P Li3PS Li3PS 4 4 LiP7 Li3P7 Li3P7 Li2S LiP LiP 7 P P4S9 4S3 P4S7 P2SP 5 Li S 3P Li9Al4, AlNLi3P, Li2S xS/(xP+xS) (c) Li (d) GeS2 S Lithiation Li Li2GeS3 15Ge4 Li4GeS Li9Ge 4 4 Li2S Delithiation LiGe Li4GeS4 Li GeS Li4GeS4 2 3 Ge Li2S GeS GeS2 LiGe Ge S Li9Ge 4 Li Al , ALlNi Ge , Li S Li Ge9 4 15 4 2 15 4 xS/(xGe+xS) Figure 4.2. Phase diagrams of a) Li-P-S and c) Li-Ge-S system, and the equilibrium voltage profiles and phase equilibria for lithiation and delithiation reactions of b) Li3PS4 and d) Li4GeS4. Li-stable phases that are electronic insulating (Li3P and Li2S) and that are electronic conductive (Li15Ge4) are colored green and red, respectively. The bottom bar in a) and c) represents the phase equilibria with Li metal as a function of atomic fraction xS and xM, where passivating and non-passivating ranges are colored green and red, respectively. The lithiation and delithiation paths in a) and c) are marked as blue and orange dashed lines, respectively. These lines represent constant ratio of S and M atomic fraction xS to xM (M = P, Ge) but varying Li content in the composition. Li-Ge-S system (Fig. 4.2c) is commonly used in the design of thio-LISICON electrolyte. The lithiation reaction of Li4GeS4 (Fig. 4.2d) would start with the reduction of 59 Ge at 1.62 V (Table 4.1), continue with Li-Ge alloying reactions, and eventually form Li15Ge4 and Li2S in equilibrium with Li metal (Fig. 4.2d). The entire lithiation reaction can be written as 7.75 Li + Li4GeS4 → 0.25 Li15Ge4 + 4 Li2S (ΔH = -8.04 eV or -775 kJ mol-1) In contrast to Li3P, metallic Li15Ge4 results in the interphase with mixed ionic and electronic conducting (MIEC) property. The simultaneous transport of both Li+ and e- through the MIEC interphase allows continuous, favorable Li reduction reaction, and hence cannot passivate at the interface.15-19,106,108,109 The non-passivating behavior of MIEC interphase is confirmed in previous experiments, where the interfacial layers and resistance grow significantly over a short period of time.19 Any Li-Ge-S ternary compound leads to the same phase equilibria of Li15Ge4 and Li2S on Li (red bar in Fig. 4.2c). Any Ge-containing ternary sulfide, regardless of its composition, would not be passivating against Li metal. Therefore, cations play a critical role in forming passivating or non-passivating interphase layers. Unless in nitride anion chemistry, most metal or metalloid cations will be reduced (Fig. 4.1) by Li metal to form metal or Li alloys (Table 4.1 and Appendix), leading to the formation of non-passivating MIEC interphase. The strategy to protect the reduction of metal or metalloid cations is critical. 4.3. Anion Effect on Lithium Metal Stability and Passivation. We illustrate the effects of anion chemistry on Li stability and passivation using Li- Al-X (X = O, S, F, and N) ternary systems as examples (Table 4.1, Fig. 4.3 and Appendix). The lithiation reactions of ternary oxides, sulfides and fluorides to form Li9Al4 and Li-X 60 binaries are thermodynamically favorable, leading to spontaneous interphase layer formation (Appendix). In Li-Al-O system, which is commonly used for coating layer materials,84,110,111 the phase equilibria with Li metal are Li9Al4 alloy and Li2O (Table 4.1 and Fig. 4.3a). Similarly, Li metal phase equilibria of Li-Al-S and Li-Al-F systems also include Li9Al4 (Table 4.1 and Fig. 4.3b-c). The presence of metallic Li9Al4 leads to a MIEC interphase, which may not passivate. These computation results are consistent with the experimental observation of the reduction of Al2O3 and Al2S3 at low potentials.112,113 The Li-Al-N system shows unique behavior compared to O, S, F-based systems (Table 4.1 and Fig. 4.3d). Al-abiding nitrides AlN and Li3AlN2 have negative cathodic limits and positive reaction energy for lithiation (Table 4.1 and Appendix). The phase equilibria with Li metal may consist of different combination of thermodynamically stable phases against Li metal, such as Li9Al4, AlN, Li3AlN2, and Li3N, depending on the atomic fraction of Al xAl and nitrogen xN in the material composition. When xN < xAl (red bar in Fig. 4.3d), the phase equilibria of Li reduction are Li9Al4 and AlN through the reaction Li + Li\jkAl\mnN\p → LirAls + LitN, leading to non-passivating MIEC interphase. At higher N content of xAl < xN < 2xAl, the formed Li-stable phase equilibria are AlN and Li3AlN2, through the reaction Li + Li\jkAl\mnN\p → LitAlNu + AlN. And at even higher N content xN > 2xAl, the phase equilibria are Li3AlN2 and Li3N (Fig. 4.3d), through the reaction Li + Li\jkAl\mnN\p → LitAlNu + LitN. Since Li3N, Li3AlN2 and AlN are electronic insulating,86,114,115 the formed interphase would be passivating when xN ≥ xAl (green bar in Fig. 4.3d). Such passivation 61 interphase between Li metal and Li-Al-N compounds is similar to the spontaneously formed inter-layer at the Li-LiPON interface. In addition, both Li3AlN2 and Li3N are ionic conducting materials,86,115 which facilitate interfacial ionic transport and reduce interfacial resistance. (a) Li (b) Li Li9Al Li Al4 Li O 9 4 2 Li SLi3Al2 Li3Al 2 2 LiAl Li2O2 LiAl Li5AlO4 Li5AlS4 LiAlO2 LiAlS2 Al LiAl O 5 8 Al AlAl O O2 2 S3 S 2 3 Li9Al4, Li2O Li9Al4, Li2S 0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 xO/(xAl+xO) xS/(xAl+xS) (c) Li (d) Li Li9Al4 Li Al Li N9 4 3 Li3Al2 Li3Al2 LiAl LiF LiAl Li 3AlN2 Li3AlF6 LiN3 Al AlF3 F2 Al AlN N2 Li9Al4, LiF Li9Al4, AlN AlN, Li AlN Li3N,Li3AlN3 2 2 0.0 0.2 0.4 0.6 0.8 1.0 0.0 0.2 0.4 0.6 0.8 1.0 xF/(xAl+xF) xN/(xAl+xN) Figure 4.3. Phase diagrams of a) Li-Al-O, b) Li-Al-S, c) Li-Al-F, d) Li-Al-N systems. Li-stable phases that are electronic insulating and that are electronic conductive are colored green and red, respectively. The bottom bar represents the phase equilibria with Li metal as a function of anion and cation atomic fraction xX and xM, where passivating and non-passivating ranges are colored green and red, respectively. 62 Among these four anion chemistry with Al cation, nitride is the only anion chemistry that is stable or can form passivating interphase layer at Li metal contact. This passivation mechanism is activated at high nitrogen content xN ≥ xAl, where electronic insulating nitride phases are formed to passivate and stabilize the interface. If nitrogen content is low, the formed interphases would still contain Li9Al4 alloy and cannot passivate (bottom bar in Fig. 4.3d). The passivation is enabled by the electronic insulation of (lithium) metal nitrides, such as Li3AlN2 and AlN, formed against Li metal. The stability of these nitrides are thermodynamic intrinsic and are unique to nitride. For example, the calculated cathodic limit of ternary nitride Li3AlN2 is -0.32 V (referenced to Li/Li+), which is lower than ternary oxide LiAlO2 (0.17 V), and ternary fluoride Li3AlF6 (1.06 V) (Appendix B). 4.4 Stability and Interphase Passivation of Nitrides We calculated phase equilibria with Li metal for Li-M-N compositions as a function of xM and xN (Fig. 4.4), according to the reaction Li + Li\jkM\wN\p → Phase equilibria with Li metal. The excellent electrochemical stability of metal nitrides against lithium metal is general for many cations (Fig. 4.4). For cation M = Mg, Ca, Al, Hf, Sc, B, Zr, Si, Ti, Ta, Nb, W, and V, passivating interphase may form depending on the material composition. At low N content, the formed interphase includes M metal or Li-M alloy, which are electronic conducting and may not provide passivation. At high N content, the formed interphase consists of Li3N and (lithium) metal nitrides (Fig. 4.4). If formed metal nitride is electronic insulating, the interphase would be passivating. These lithium ternary nitrides that are stable against Li metal include LiMgN, LiCaN, Li3AlN2, Li2HfN2, Li3ScN2, Li3BN2, Li2ZrN2, 63 Li5SiN3, Li5TiN3, Li4TaN3, Li7TaN4, Li7NbN4, Li6WN4 and Li7VN4. At high N content, these metal nitrides form at the interface and may passivate against Li metal. Figure 4.4. Phase equilibria of Li-M-N compositions in equilibrium with Li metal. The formed products due to Li reduction is a function of atomic fraction xN and xM (y axis) in the composition. (①: [Li4TaN3, Li7TaN4]; ②: [Zr3N2, ZrN].) The passivating and non-passivating phase equilibria regions are colored green and red, respectively. 64 By contrast, some cations, such as M = Ge, Ga, Sn, and Zn, cannot be stabilized in nitrides against Li metal regardless of the composition. Li metal phase equilibria of these Li-M-N systems always contain metal or Li-metal alloy (Fig. 4.4) forming MIEC interphase. The formed interphase layer is likely not passivating if these cations are used. The systems with only non-metal elements, such as M = P, Cl, Br, I, is in general compatible with Li metal. When in equilibrium with Li metal, the reduction products are still electronic insulating phases, including Li3P, LiCl, LiBr, and LiI. Some non-metal elements may change from cation to anion after reduction. In these nitride systems with only non-metal elements, passivating interphase would form at the Li metal interface regardless of xN (Fig. 4.4). Similarly, using only these non-metal elements also leads to Li-passivating interphases in oxides and sulfides, as observed in LiPON, Li3OCl/Li3OBr, Li9S3N, Li7P3S11, and Li7P2S8I solid electrolytes, which are Li metal compatible.3,54,70,116,117 4.5 Stability and Interphase Passivation of Mixed-anion Nitrides Since nitrides in general show good electrochemical stability against Li metal, here we explore the stability of mixed anion chemistry with nitride corresponding to the doping of nitrogen or nitride into other anion chemistry. We first studied the Li-Al-O-N quaternary system as an example, in which Li-Al-O ternary compound is doped with nitrogen. In the quaternary Li-Al-O-N system, the final phase equilibria in contact with Li metal will only consist of Li-stable phases, including Li9Al4, Li2O, Li3N, AlN, and Li3AlN2 (Fig. 4.5a). Among those phases, only Li9Al4 is MIEC, and all other phases are electronic insulating. Similar to the ternary case, the phase equilibria with Li metal are determined by atomic fraction of Al and nitrogen, xN and xAl. At low N content of xN < xAl, the formed phase 65 equilibria consist of Li9Al4, Li2O, AlN represented by the red triangle in Fig. 4.5a-b. For example, the lithiation reaction of Al3NO3 (gamma-ALON) follows the blue dashed line in Fig. 4.5a to form these products. This lithiation reaction has the voltage profile and phase equilibria shown in Fig. 4.5c. For composition at higher N content of xAl < xN < 2xAl (green triangle II in Fig. 4.5a-b), Li phase equilibria are Li3AlN2, Li2O, and AlN. At even higher N content of xN > 2xAl (green triangle I in Fig. 4.5a-b), the phase equilibria with Li are Li3N, Li2O, Li3AlN2. The phase equilibria with Li metal are a function of the Al-O-N composition, and are represented by the grand potential phase diagram in equilibrium with Li metal (Fig. 4.5b). At high N content (green triangle I and II in Fig. 4.5a-b), the formed interphase after lithiation is electronic insulating and hence passivating. These results suggest a strategy of introducing a sufficiently high amount of N to stabilize oxide compounds, which are not stable against Li metal. High-dose nitrogen doping inhibits the reduction of Al and the formation of Li-Al alloy, leading to a passivating interphase formed against Li metal. Mixing nitrogen can also stabilize Li-M-O oxides with other cation M = Mg, Ca, Al, B, Zr, Si, Ti, Ta, Nb, V and W. Similar to M = Al, having high amount of nitrogen in Li-M-O-N systems leads to forming electronic insulating nitrides against Li (Fig. 4.6, Table 4.2 and Appendix B). For example in Li-Nb-O-N system, electronic insulating phase equilibria form at high N content xN ≥ 4xNb (Fig. 4.6a and Table 2). In Li-Ca-O-N system, Li-stable phase equilibria, including electronic-insulating LiCaN nitride and CaO oxide, may form as passivation interphase at the high O and N content of xN + xO ≥ xCa (Fig. 4.6b and Table 4.2). Therefore, high-dose nitrogen doping is effective in passivating many oxides with metal cations. 66 (a) Li (b) O Li2O Li3N Li2O I Li9Al4 Li3AlN2 II Al3NO3 O N I II Al3NO3 Li3N Li9Al4 AlN Al N Li3AlN2 AlN Al (c) Al3NO3 Al3NO3 LiAl5O8 Al AlN LiAlO2 LiAl Li3Al2 Li5AlO4 Li2O Li9Al4 Figure 4.5. a) Li-Al-O-N quaternary phase diagram showing only Li-stable phases. The blue dashed line represents the lithiation path of Al3NO3 (blue square point). b) Grand potential phase diagram of Li-Al-O-N system in equilibrium with Li metal. The Gibbs triangle of the grand potential phase diagram is based on non-Li composition Al-O-N, which determines the phase equilibria after lithiation. For example, the phase equilibria of Al3NO3 (blue square) with Li metal are Li9Al4, Li2O, and AlN, corresponding to the red triangle region. The composition regions that form passivating and that form non-passivating interphases are colored green and red, respectively. c) Voltage profile and phase equilibria for the lithiation of Al3NO3. 67 (a) O (b) O Li2O Li2O CaO Li Li N 16 (NbN4)2O 3 Nb Li3N Li2Ca N Li7NbN4 Nb N LiCaN Ca2N Ca (c) O O Li (d)2O Li2O Li3N Li3P Li3N Li15Ge4 N P N Ge Figure 4.6. Grand potential phase diagrams of a) Li-Nb-O-N, b) Li-Ca-O-N, c) Li-P- O-N, d) Li-Ge-O-N systems in equilibrium with Li metal. The composition regions that form passivating and that form non-passivating interphases are colored green and red, respectively. Li-P-O-N oxynitride system has Li phase equilibria consisting of only electronic insulating phases, Li3N, Li2O, and Li3P (Fig. 4.6c), which are observed by in situ XPS experiments.45 In addition, Li3N and Li3P are fast lithium ionic conductors, which facilitate interfacial Li transport and reduce interface resistance. The formation of passivating interphase results in excellent Li compatibility of LiPON. Similarly, other non-metal elements, such as Cl, Br, and I, also always form passivating interphase at Li metal. Such 68 passivating layer, as in the case of Li-LiPON interface, may achieve good stability on Li metal anode. Table 4.2. Composition range in Li-M-N-O systems that form stable passivation interphase. Cation M Composition range forming passivation interphase layer Al xN ≥ xAl Nb xN ≥ 4xNb B xN ≥ 2xB Zr xN ≥ 2xZr Ta xN ≥ 4xTa or xN ≥ 3xTa + 2xO Ti xN ≥ 3xTi Si xN ≥ 3xSi V xN ≥ 4xV W xN ≥ 4xW Mg xN ≥ xMg Ca xN + xO ≥ xCa P, Cl, Br, I Any xN and xM Ge, Sn, Ga, Zn None Unfortunately, for oxides with some other cations, passivating interphases cannot be achieved by the introduction of nitrogen. For example, the phase equilibria of Li-Ge-O-N system with Li metal are always Li3N, Li2O, and Li15Ge4 regardless of N content (Fig. 4.6d). The presence of Li15Ge4 would lead a MIEC interphase. Similarly, materials containing Sn, Ga, or Zn have Li phase equilibria always containing metal or 69 Li-metal alloy (Appendix B and Table 4.2), and hence may not be passivated by the formation of nitrides. 4.6 Discussion Our computation results revealed the Li metal stability of materials with different cation and anion chemistry. Most oxides, sulfides, and fluorides, which are widely studied as current solid electrolytes or coating layer materials, are not stable against Li metal. Metal and metalloid cations in oxides, sulfides, and fluorides (and other halides) are reduced at low potentials in most of these compounds. The thermodynamically favorable Li reduction of metal cations would lead to the formation of mixed ionic and electronic conducting (MIEC) interphase and continuous decomposition of bulk materials, resulting in poor stability against Li metal. To avoid the detrimental effects of metal cations on stability, using only non-metal elements, such as P, Cl, Br, I, instead forms passivating interphase and enables stability against Li metal, as demonstrated in some Li-compatible solid electrolytes, such as LiPON, Li3OCl/Li3OBr, Li9S3N, Li7P3S11, and Li7P2S8I.3,54,70,116,117 However, there is a limited choice of these compounds for stabilizing Li metal. Our results provide the composition space to select, design, and discover materials that are stable against Li. The major discovery of our study is that nitride anion chemistry has significantly better stability against Li metal compared to oxide, sulfide, and halide. Many nitride materials are thermodynamically stable against Li metal. In addition, the nitrides formed at the interface during lithiation become stable, passivating SEI against Li metal. This chemistry knowledge from our computation suggests multiple strategies as follows to stabilize materials against Li metal and to improve the performance of Li metal anode. 70 1) Many nitride materials that are electronic insulators may be used as protective buffer layers on Li metal anode. These buffer layers are thermodynamically stable against Li metal, thus will protect the electrolyte from reduction and will have long-term stability on Li metal. For example, Li3N is demonstrated as an effective buffer layer material to protect liquid or solid electrolytes and to improve the cyclability and coulombic efficiency of Li metal anode.84,99,100 Besides Li3N, many Li-containing nitrides that are stable against Li metal, e.g. Li3AlN2, Li5SiN3, Li3BN2, LiMgN, LiCaN, Li2HfN2, Li3ScN2, Li2ZrN2, Li5TiN3, Li4TaN3, Li7TaN4, Li7NbN4, Li6WN4 and Li7VN4, may also be used as buffer layers on Li metal anode. In addition, developing new materials, such as Li18P6N16 nitridophosphate,118 may also be a promising direction. 2) Our results also suggest that high-dose nitrogen doping in materials can lead to in situ spontaneous passivation against Li metal. Surface modification techniques such as nitrogen ion/plasma treatment and other nitriding techniques can locally enrich nitrogen content on the surfaces of the material. The nitrogen-rich surface put into contact with Li metal would form nitride passivation interphase layers and transform into a Li-stable interface. For example, if one can introduce sufficient amount of N into the surface region of Al2O3 to the composition range xN ≥ xAl, passivation phase Li3AlN2 /AlN and Li2O would form (Fig. 4.5). The necessary nitrogen content and composition to passivate each cation are provided by our calculations (Fig. 4.6, Table 4.2 and Appendix). Various experimental approaches can be used to introduce nitrogen into the material and to stabilize the material interfaces against Li metal. For example, applying nitrogen-doped thin films on Li metal anodes has shown to improve Li metal stability and to inhibit lithium dendrite.119 Multiple 71 studies also demonstrated nitrogen doping on carbon-based anode materials greatly improves the cycling stability.120-122 3) Another way to locally enrich nitrogen is introducing Li3N to further react at the interface and to form a stable SEI layer. Li3N may be introduced as a coating layer or by nitriding Li metal. For example, Li3N can react with Al2O3 at the contact interface and form stable SEI consisting of Li3AlN2 and Li2O through an exothermic reaction, Al2O3 + 4 Li3N → 2 Li3AlN2 + 3 Li2O (ΔH = -5.76 eV or -556 kJ mol-1). A similar strategy was demonstrated on Li metal protection by Polyplus.84 4) In addition, mixing nitrogen-containing compounds, such as nitride, nitrite, and nitrate, into the materials can also stabilize the interface. For example, Al2O3 mixed with LiNO3 may also form a SEI of Li3AlN2 and Li2O in contact with Li through an exothermic reaction Al2O3 + 4 LiNO3 + 32 Li → 2 Li3AlN2 + 15 Li2O (ΔH = -58.20 eV or -5615 kJ mol-1). A similar strategy of using nitrogen-containing additives in liquid electrolytes was demonstrated in forming stable SEI layer on Li metal anode.123-127 For example, LiNO3 as additive was reported to significantly improve the cyclability and coulombic efficiency of Li metal anode by forming a SEI containing lithium nitride and oxynitrides.128-130 In agreement with our prediction, high concentration of LiNO3 is needed to achieve high nitrogen content and to form a stable SEI.125 In addition, other nitrogen-containing compounds, such as cyanide and cyanate, are demonstrated for nitriding metals. Novel nitrogen-containing compounds may be developed for nitriding materials in lithium batteries. While our computation demonstrated the thermodynamic origin of using nitrides and rich nitrogen content to stabilize Li metal anode, experimental techniques to realize 72 these new strategies in lithium batteries may need further study to address multiple potential challenges. First, nitride buffer coating layer is desired to be thin (~10 nm), dense, and uniform, to achieve low interfacial resistance and good passivation.131 Applying such coating may requires thin-film-based techniques, such as sputtering, vapor deposition, atomic layer deposition (ALD), and plasma nitriding, which may not be cost competitive.132 In addition, some nitrides may be moisture or air sensitive,131 which requires additional protection during manufacture and processing. Gas nitriding techniques based on nitrogen rich gas, such as ammonia, may be an economic and scalable approach as demonstrated in the nitriding of metals.132 However, ammonia treatment may induce H into the materials, which may be detrimental for the properties of electrolytes or electrodes. Experimental techniques to economically coat nitride layers and to effectively dope high nitrogen content at the interface need further research and development. The aforementioned strategies of using nitrogen anion chemistry to stabilize Li metal anode may offer multiple additional advantages. The thermodynamically favorable lithiation reactions provide interfacial wetting, which would promote interfacial contacts.16,106,133 Moreover, a fraction of nitrogen may be lithiated to form Li3N, a good ionic conductor, as part of the interphase to facilitate interfacial Li ion transport. In addition, many nitrides exhibit good tolerance to stress, which is a desirable mechanical property for stable interfacial layer during cycling and for suppressing Li dendrite growth.132,134-138 Therefore, the formation of stable SEI layer using nitride anion chemistry provides good stability against Li metal, reduces interfacial resistance, and may mechanically suppress Li dendrite formation. High nitrogen content as predicted by our calculations (Fig. 4.6 and Table 4.2) is necessary to form phase equilibria only consisting of passivating phases and to completely 73 inhibit the formation of reduced metallic phases, while a lower nitrogen content may still retard interfacial degradation and benefit the performance of Li metal anode as a result of kinetic stabilization. The limitations of using the thermodynamic scheme to predict interfacial phenomena should be noted as follows. In this study, we approximated the interphase as the phase equilibria, which are the most thermodynamically favorable states, whereas kinetics of reaction and diffusion are not considered. While kinetics and diffusion play crucial roles in the reactions at the interfaces, the thermodynamics provide an ultimate boundary of stability. Given facile Li diffusion, it is expected the phase equilibria would form at the immediate contact with Li metal, as demonstrated in previous in-situ XPS study.17-19,45,107 Despite possible kinetic or diffusional limitations for such reactions, the thermodynamic driving force would always exist for forming the most favorable phase equilibria. These decomposition reactions may have a significant impact on the long-term stability of the battery. It is possible that the formed interphase layers are different from thermodynamic equilibria. For example, some phases that are not stable against lithium metal may form due to kinetic stabilization. In addition, kinetic limitation of the lithiation reduction at the interface may lead to the formation of intermediate decomposition products, which may be passivating to slow down further degradation.107 As a result of such kinetic effects, some intermediate decomposition products or less than adequate nitrogen doping may still improve the stability at Li metal interfaces. Nevertheless, designing the interface with thermodynamic intrinsic stability against Li metal is still desired to realize long-term stability, cyclability, and coulombic efficiency of the lithium battery. 74 Additional approximations made in our computation scheme and their potential impact are noted as follows. We used the mixture of bulk phase properties to approximate the interphase property, and did not consider the microstructures of interphase. For example, if the formed electronic-conducting phases are not percolating from Li metal to electrolyte, the overall interphase layer would still be passivating. The effects of microstructures, local off-stoichiometry, defects accumulation, and kinetically stabilized meta-stable phases, which were not considered in our calculations, may also impact interphase properties and passivation effects. In addition, our scheme relies on known phases to predict Li reduction products of materials. There may be errors in the predicted phase equilibria if some phases that exist in nature were not included. This scenario is more likely in nitride systems, which have been relatively less studied than oxides. In addition, due to the lack of available electronic conductivity data of many materials, we used a simple cutoff based on DFT calculated band gap and the valence states of metal cations to judge whether a phase is electronic conductive or insulating. While the trend of the nitride stability is captured, using experimentally measured electronic conductivity would be more reliable in estimating the interphase properties. Some passivating interphase may still have a low electronic conductivity leading to a slow interphase layer growth and impedance increase during cycling or over time. For example, while the SEI layer on graphite anode in commercial lithium ion battery is well accepted as a passivation film, Dahn et al. showed slow growth of SEI and continuous lithium loss.139,140 The Li-compatible solid electrolyte Li7P3S11, also shows a slow interphase impedance increase over time.107 Nevertheless, these passivating interphases have significantly better stability than non-passivating interphases, and provide improved cyclability and battery performance.19 75 4.7 Conclusion In this chapter, our computation revealed that material stability against lithium metal is governed by cation and anion chemistry. Metal cations in oxides, sulfides, and fluorides are usually reduced by lithium metal, leading to poor stability against lithium metal. We discovered that nitride anion chemistry shows unique stability against Li metal. In general, metal nitrides have a significantly lower reduction potential than oxides, sulfides, and fluorides, and many nitrides are thermodynamically stable against Li metal. On the basis of this chemistry knowledge, we suggest new strategies to form stable SEI on Li metal anode. Many nitrides that have intrinsic Li metal stability can be used as buffer layers to protect materials against Li metal. In addition, high-dose nitrogen doping and nitrogen enriching at the interface can lead to spontaneous formation of a stable, passivating SEI on Li metal, transforming a lithium unstable interface into a lithium stable interface. Specifically, many cations, such as Mg, Ca, B, Al, Zr, V, W, Si, Ti, Nb, and Ta, can be protected in nitride anion chemistry systems at sufficiently high nitrogen content. However, some cations, such as Ge, Sn, Ga, and Zn, are always reduced by Li metal regardless of anion chemistry and composition. Our results provide guiding principles for the selection, design, and discovery of materials with Li metal stability, and predict interfacial engineering strategies to stabilize Li metal anodes in lithium batteries. 76 Chapter 5. Ionic Transport in Grain Boundaries of Lithium Garnet The key advancement of ASLiBs is mainly because of the discovery and development of solid electrolyte materials. Li-stuffed garnet Li7La3Zr2O12 (LLZO) is a very promising group of material.26 For example, the Al-doped cubic phase Li6.25Al0.25La3Zr2O12 exhibits bulk ionic conductivity ~ 0.4 mS/cm at room temperature.141 The fast bulk ionic transport in Li-stuffed garnet has drawn a lot of research effort to understand the underlying mechanism of its fast ion diffusion.21-25 The cubic phase lithium garnet (c-LLZO) have partially occupied 24d and 96h Li sites, which forms 3D interconnected pathways.21 The high Li content activates the concerted migration mode,22 which enables lithium garnet material to have fast ion transport and thus high ionic conductivity. However, the solid electrolytes in all-solid-state battery are usually polycrystalline, rather than single crystal. Besides the bulk diffusion, ion diffusion along and across the grain boundaries (GBs) are an essential part of the ionic transport. Therefore, a good solid electrolyte should also have decent ion transport at grain boundaries. Sluggish ionic transport at grain boundaries would have an negative impact on the overall ionic conductivity of SEs and the battery performance.26,27 Grain boundaries have different structures and chemistry from bulk, and thus different properties. Since the discovery of superionic cubic phase Li-garnet materials, researchers have been investigating the grain boundary ion transport in this material. Early experiments in 2007 reported that for Li7La3Zr2O12, the resistance contributed from grain boundary is comparable to that from bulk under room temperature.142 In 2010 another work reported that by increasing the calcination temperature, well-sintered Li7La3Zr2O12 sample 77 shows negligible grain boundary resistance.57 Later Al dopant becomes widely used to stabilize the cubic phase Li-garnet under room temperature. In 2014, experimental results of resolved grain boundary and bulk resistance in Al-doped c-LLZO were reported, where grain boundary shows larger resistance and higher activation energy compared to bulk.143 A following work in 2015 reported that using proper hot-pressing temperature can increase the sample density and decrease the grain boundary resistance in Al-doped c-LLZO.141 While another work also in 2015 reported that using solid state reaction, the synthesized Al- doped c-LLZO have more ionic conductive grain boundary compared to bulk.144 Moreover, the grain boundaries in solid electrolytes are closely related to lithium dendrite growth, which is currently a critical problem for all-solid-state lithium batteries. Recent experimental works show that lithium dendrite preferentially grows inside the grain boundary of solid electrolyte, which eventually leads to short circuit and cell failure.28,29 The properties of grain boundaries, such as the ion diffusion property, may help us understand the dendrite growth in the grain boundaries of solid electrolytes. However, despite its significance, current understanding on the structure, composition and more importantly, the diffusion properties of grain boundaries in solid electrolyte materials, is still limited. This is mainly due to the complex nature of grain boundaries, which poses difficulty for both experimental characterization and computational modeling. In this chapter, we aim to bridge these knowledge gap. Using molecular dynamics simulation, we investigate the ionic diffusion properties of grain boundaries in lithium garnet. We study the distinct ion diffusion behavior at grain boundaries different from the 78 facile ion transport in the bulk. We investigate the relationship between the diffusion properties and the structure properties of the grain boundaries. We also reveal why the conventional space-charge theory at grain boundary does not apply to super-ionic conductors like lithium garnet. 5.1 Computational Methods 1. Interatomic potential We performed classic molecular dynamics simulation to study the ionic diffusion in Li garnet. The sample we constructed is Al doped cubic phase Li garnet (below denoted as c-LLZO) based on experimental reported structure.21 We used composition Li6.25Al0.25La3Zr2O12, in which Al3+ were randomly doped to Li1 (24d) sites to stabilize the cubic phase. The interatomic potential takes a pairwise form and consists of two parts: a long- range term describes the Coulomb interaction plus a short-range term describes the repulsive and van der Waals interactions. The short range interaction term were modeled by the Buckingham potential Ebuck(r) = Ae-r/ρ - C⋅r--6 (r < rc), where r is the atomic pair distance. rc is the cutoff distance and is set to 12 Å. A, C, ρ are fitting parameters and are listed in Table 5.1, most of which are taken from literatures.145,146 All ions except oxygen ions take their formal charge. The core/shell model is applied on O ions to account for the polarization behavior. The shell mass Mshell, shell charge Y and the coupling core-shell harmonic spring constant k are also listed in Table 5.1. 79 Table 5.1. Interatomic potential parameters in Al-LLZO system. Buckingham potential parameters Core-shell model parameters for O Ionic pair A (eV) ρ (Å) C(eV·Å-6) Y (e) k (eV·Å-2) Mshell(a.u.) Li - O 1430.0 0.260 4.0 -2.76 30.2 1.6 La - O 4579.23 0.3044 0.0 Zr - O 1385.02 0.350 0.0 Al - O 1738.0 0.289 0.0 O – O 22764.3 0.149 27.63 Fig. C1 shows the lattice parameters of Al-doped LLZO (Li6.25Al0.25La3Zr2O12) and undoped LLZO (Li7La3Zr2O12) in quenching simulation. Our potential successfully captures the spontaneous cubic-tetragonal phase transition behavior in undoped LLZO, while the Al-doped LLZO is stabilized at cubic phase during quenching. The calculated lattice parameters and thermal expansion coefficient agree very well with experimental results (Fig. C1).21 The room temperature lattice parameter difference is only 0.33%. Fig. 5.1 shows the calculated the diffusion properties using the aforementioned classic interatomic potential. We observed very typical Arrhenius behavior and the calculated activation energy is 0.23 eV, which is very close to the value from experiments147 and from first principles calculations.22,24 (Fig. 5.1) The room temperature conductivity is roughly one magnitude higher than experimental values. In terms of lattice parameter, activation energy and the capture of phase change, our potential shows good performance compared to previous reported classical potentials, especially in the activation energy. 80 Figure 5.1. Arrhenius plot of ionic conductivity from this work and previous literature reports, including experimental reports (EXP),147 ab initio molecular dynamics simulation (AIMD),22,24 and classic molecular dynamics simulation (CMD).148-150 The activation energy listed in the legend are either from literature if available, or fitted if unavailable from literature. 2. Grain boundary sample construction and molecular dynamics simulation details Since garnet is a group of material share similar framework structure,26 we assumed they probably share similar low-energy surface orientation. Garnet-type minerals generally have <110> and <112> as the low energy surface orientation.151 Therefore, here we also selected <110> and <112> as candidate orientations to 81 construct twist grain boundaries. To satisfy the periodic boundary condition. We construct S3 <110> and S5 <112> twist grain boundaries. The grain boundaries are perpendicular to z axis, and the twist angles are 70.53° and 78.46° respectively around the z axis. The simulation box size (under 500 K) are 77.8 Å ´ 110.1 Å ´ 147.6 Å for <110> grain boundary and 91.7 Å ´ 56.2 Å ´ 127.8 Å for <112> grain boundary, respectively. All simulation are performed with LAMMPS package.152 Samples were first heated up to 1200 K and equilibrated for 0.1 ns, and then quenched to target temperature. The self-diffusion simulations were performed under NVE ensemble. 3. Local diffusivity calculation We tracked the atom trajectories and use the mean-squared-displacement (MSD) to calculate the diffusivity. Conventionally, the diffusivity, as a statistical quantity, is calculated by averaging MSD on all diffusive ions. However, in our grain boundary model, the diffusivity is also a local quantity, which has a non-uniform distribution along z direction due to the existence of grain boundaries. To capture the local diffusivity distribution, we calculate local diffusivity based on the locally averaged squared displacement. We first sliced the sample into 40 layers perpendicular to z direction, and then grouped the atom displacement into different layers based on the average z position during that period of lag time. We then calculate the MSD for each layer with different lag time. In this way, we can capture the local diffusivity distribution with a reasonable special resolution. 82 5.2 Diffusion and Structure Properties of Grain Boundaries in Lithium Garnet (a) (b) (c) (d) (e) (f) (g) Z coordinates [Å] 83 Average displacement [Å] Occupancy Norm. ρLi Ea [eV] Diffusivity [cm 2/s] Snapshot Figure 5.2. a) constructed slab structure of twist <110> S3 grain boundaries. The twist angle is 70.53° and the slab thickness is 2 x 73.80 Å. The labeled axes correspond to the conventional cell of lithium garnet. The navy polyhedral are ZrO6 polyhedral. b) Local ionic diffusivity in the grain boundary slab (blue line) and reference bulk diffusivity (red dashed line) at 500 K. c) Local ionic diffusivity along the grain boundary plane (xy) direction (orange line) and across the grain boundary plane (z) direction (green line). d) Local activation energy in the grain boundary slab (blue line) and reference bulk activation energy (red dashed line). e) The temporally averaged Li density in the grain boundary slab (blue line) and the bulk reference (red dashed line). Both are normalized to per formula. f) The temporally averaged site occupancies of Li1(24d) and merged Li2 (48g) sites. g) The averaged displacement between the initial position and average position of framework atoms La, Zr and O. O positions are based on the positions of O cores. We picked <110> S3 twist grain boundary as a representative example to study the grain boundary diffusion in lithium garnet. Fig. 5.2a snapshot shows the constructed structure used in our modeling. Two <110> S3 twist grain boundaries locate at z = 0 Å and z = 73.80 Å, respectively (red dashed line). Fig. 5.2b shows the calculated local ionic diffusivity versus z coordinates at 500 K (the blue line). In the center part of the slabs, the diffusivity is very close to the bulk diffusivity 10-6 cm2/s (the red dashed line). However, the diffusivity shows a sharp decrease near the two grain boundaries. In the center region of ±1 nm near the grain boundary planes, the diffusivity drops to 10-7 cm2/s, which is about one order of magnitude lower than that of the center part of the slab. This indicates that the grain boundaries impede the facile ion diffusion in lithium garnet. Fig. 5.2c shows the calculated the directional ionic diffusivity along the grain boundary (xy) and across the grain boundary 84 (z) direction at 500 K. Diffusivity along both directions show similar behavior, which drops to an about one-order-of-magnitude lower in the ±1 nm region near the grain boundaries. This indicates that near the grain boundary core region, diffusion along all directions slows down. The trend observed at 500 K (Fig. 5.2 b-c) also applies to other temperatures. We fitted the diffusivity using Arrhenius equation to get the local activation energy Ea (Fig. 5.2d). The local activation energy at grain boundary core increases to 0.36 eV, while for regions that are more than 1 nm away from the grain boundaries, the activation energy is still close to the bulk value 0.23 eV. From our calculation, we found that grain boundaries slow down the ionic diffusion, in both along and across the grain boundary directions. The affected region is localized, only in an approximately 2 nm-thick region beside the grain boundary plane. At the core of grain boundaries, the activation energy of ionic diffusion increases to 0.36 eV. The calculated ionic conductivity at 300 K (extrapolated from Arrhenius relationship) is only 0.03 mS/cm, which is nearly two orders of magnitude lower than the bulk. Both in plane and out of plane diffusion are slowed down, which may because the diffusion channels of lithium garnet are 3-D interconnected. For region that are more than 1 nm away from the grain boundary plane, the ionic diffusion appears to be unaffected. Both the diffusivity and activation energy are very close to bulk value, suggesting that these regions still exhibit bulk-like fast ion diffusion. Structure analysis was performed to understand the origin of slower diffusion near grain boundary. Fig. 5.2e shows the distribution of temporally averaged local Li content ρLi normalized to per formula. Compared to the bulk value (6.25 per formula, red dashed line), 85 there is a slight decrease of lithium content near the two grain boundaries, as they drop to approximately 5.2 and 6 Li per formula, respectively. The two grain boundaries show some variance, which may because the two grain boundaries are not completely symmetric. We also performed site analysis to calculate the site occupancy on Li1(24d) and Li2 (original 96h sites, but merged neighboring pairs into 48g for occupancy calculation). We used 0.9 Å as a cutoff radius to judge whether a site is occupied or not. As shown in Fig. 5.2f, the occupancy of both Li sites show decrease near the grain boundaries, especially for Li1 sites. This suggested that the Li ions near the grain boundary are not at their original lattice sites. The mismatched lattice at these high-angle grain boundaries strongly disturbs the original interconnected channels of Li ions, which provides fast ion diffusion. As a consequence, the grain boundaries exhibit slower ion diffusion. In comparison, we also quantified the framework (La, Zr and O atoms) change near the grain boundary. Fig. 5.2g shows the averaged root square displacement between their initial position and average position of these framework atoms. (O positions is calculated based on O core positions) In the region near grain boundaries, the averaged displacement for La, Zr, and O atoms are only around 0.2 Å, 0.5 Å, and 0.4 Å away from their original lattice sites, respectively. This suggested that there is no severe framework structure change near the grain boundaries. We also performed same analysis for <112> twist grain boundaries (Appendix C), and similar trends were observed. Therefore, our simulation reveals the grain boundary impedes the ion diffusion in approximately 2 nm thick nearby region. Structure analysis shows that framework atoms (La, Zr and O) change is limited, while the original Li sites 86 show occupancy decrease. The local Li reconfiguration is more likely to be the structure origin of slower diffusion near grain boundaries. 5.3 Discussion Our simulations reveal the Li diffusion mechanism at grain boundaries in cubic phase lithium garnet material. Grain boundaries slow down the ionic diffusion in the nearby region, and the slower diffusion is observed in an approximately 2 nm thick region near the grain boundary plane. At room temperature, the diffusivity in the grain boundary core is about two orders of magnitude lower than bulk, and the activation energy increases to 0.36 eV. Outside of this 2 nm-thick region, the ionic diffusion seems unaffected by the grain boundary and still show bulk-like behavior. Our structure analysis shows that the framework (La, Zr, and O atoms) shows limited change. However, the Li sublattice is severely distorted near the grain boundaries. The occupancy of Li sites show decrease near the grain boundaries, especially for Li1(24d) sites. This suggested that the Li atoms near the grain boundary moved away from their original lattice sites. Since the grain boundaries in our simulation are high angle grain boundaries, the strongly distorted Li sublattice probably leads to the absence of interconnect fast 3-D channel near the grain boundaries. Therefore, the facile ion transport in bulk lithium garnet are impeded at the grain boundaries, which results in slower grain boundary ion diffusion. This mechanism of slow grain boundary diffusion may be a general phenomenon in most lithium super ionic conducting materials. In these super ionic conductors, the crystalline bulk phases usually have highly optimized structure characteristics for fast ion 87 conduction,22,153 while the grain boundary structures usually exhibit a deviation from the ideal bulk. Therefore, the grain boundary usually has slower diffusion compared to the bulk. Our results also show there is no significant space-charge layer effect. Neither major Li accumulation nor major depletion was observed. The ionic conductivity change mostly originates from the increase of activation energy induced by structure change, instead of carrier concentration change. In fact, based on space-charge layer theory, the Debye Length can be calculated as 𝑘 𝑇𝜀𝜀 𝐿 B =#,‚ƒ, = „2𝑒u𝑐 Li where T is the temperature, 𝜀= is the relative permittivity, kB is the Boltzmann constant, cLi is the concentration of charge carriers, and 𝜀= is the absolute permittivity. As an estimation, we consider all the Li ions as effective charge carriers, and use an approximate relative permittivity e = 10.154 The Debye length at 300 K would be only 0.18 Å, even much shorter than the interatomic distance. Note that the space charge theory is based on continuum assumptions, under which both the carrier concentration and the electrostatic potential have continuous special distribution. In contrast, the calculated Debye length in lithium garnet falls into atomistic scale, at which both carrier concentration and the electrostatic potential are ill-defined. This indicates that the conventional space charge layer theory is not applicable to the super-ionic conductor material like lithium garnet. The lack of space charge layer effect may also be generally expected for super ionic conductor materials, most of which have highly disordered ion sublattice and very high carrier concentration. In super ionic conductor materials, the change of carrier concentration near grain boundaries is usually not the limiting factor of ion diffusion. Instead, the 88 mismatched lattice near the grain boundaries, especially high angle grain boundaries, may lead to high migration barrier, and thus resulting in slower diffusion. Therefore, grain boundaries in lithium garnet exhibit different ion diffusion mechanism from bulk. The abrupt change of ion transport properties at grain boundaries may affect the field distribution inside the solid electrolyte and may have impact on Li dendrite growth, which deserve further detailed studies in the future. 89 Chapter 6. Strain Induced Order-disorder Transition in Lithium Garnet Mechanical strain is an important factor that can significantly affect the ion diffusion in solid electrolytes and battery performance.30,31. During the assembling of all-solid-state battery, externally applying mechanical strain is a commonly used method to improve the physical contact and to decrease the interfacial resistance.155 Besides, electrode materials, such as Li metal anode and LiFePO4 cathode, can undergo considerable volume change during the battery cycling.156,157 This would also induce a strain on the solid electrolytes that are in intimate contact with the electrodes. Mechanical strain can have a critical impact on the ionic diffusion in solid electrolytes. It was previously reported that in oxygen ionic conductor and lithium ionic conductor materials, strain can induce orders of magnitude change in diffusivity.158,159 Nevertheless, for lithium garnet, currently there is very little information about the effect of mechanical strain on its ion diffusion properties. In this chapter, we aim to understand the strain effect on lithium garnet. By constructing atomistic model and using molecular dynamics simulations, we investigate the effect of mechanical strain on the ionic diffusion in lithium garnet material. We explore how the strain could affect the ionic diffusion in lithium garnet. We also reveal the coupling between mechanical strain and Li sublattice ordering, and demonstrate the strain can induce order-disorder transition in Li sublattice of lithium garnet, which has a significant effect on the diffusion. These understanding can be helpful to the further development of all-solid-state Li-ion battery. 90 6.1 Effect of Mechanic Deformation on the Ionic Diffusion in lithium garnet We study how mechanical deformation on lithium garnet affects its ionic diffusion properties. Tensile and compressive uniaxial strain were applied, respectively, and the engineering normal strain is 1% and 3%. The lateral strain is controlled by Poisson ratio, which was set to 0.26.136,160 Figure 6.1. Arrhenius relationship of lithium garnet under mechanical deformation. Fig. 6.1 shows the calculated diffusivity under different strain. In all cases, mechanical deformation decreases the ionic diffusivity. 1% tensile or compressive strain only slight slows the diffusion. The calculated activation energy are 0.21 eV (1% tensile, blue line in Fig. 6.1) and 0.22 eV (1% compressive, orange line in Fig. 6.1), respectively, which are very close to the bulk value (red dashed line in Fig. 6.1). Increasing tensile strain to 3% further decreases the diffusivity, while the calculated activation energy is 0.24 eV, which still does not show any significant change from the undeformed bulk. However, 3% 91 compressive strain leads to a significant decrease in diffusivity (purple line in Fig. 6.1) and an increase of the activation energy to 0.31 eV. 6.2 Strain-induced Order-Disorder Transition in Lithium Garnet Undoped lithium garnet Li7La3Zr2O12 has two polymorphs, high temperature cubic phase c-LLZO and low temperature tetragonal phase t-LLZO. c-LLZO has a disordered Li sublattice, in which Li ions partially occupy 24d and 96h sites, while t-LLZO has an ordered Li sublattice, in which Li ions fully occupy 8a, 16f, and 32g sites. The cubic phase has a higher ionic conductivity thanks to its disordered Li sublattice. To obtain the more ionic conductive cubic phase lithium garnet, dopants like Al are commonly used to stabilize the cubic phase under room temperature. Before applying deformation, our Al-doped LLZO sample stays at the cubic phase with disordered Li sublattice. Applying a 3% compressive strain would change its lattice parameters close to the undoped tetragonal phase of lithium garnet. The a / c ratio of the sample under 3% compressive engineering strain is 1.039, while the a / c ratio of tetragonal phase undoped LLZO reported from experiments is 1.036.21 This suggest that the applied compressive strain may leads Li sublattice undergo an order-disorder transition, and thus leads to slow ionic diffusion. We calculated the virtual XRD pattern of Li sublattice in the c-LLZO sample to investigate the Li sublattice ordering change. The XRD pattern is calculated based on virtual diffraction analysis,161 using a wavelength of 1.542 Å (Cu Kα as the virtual radiation source). The XRD pattern is determined by both the lattice parameters and the sublattice ordering. The compressive strain loading leads to tetragonal lattice parameters with a / c = 92 1.039. We calculated the XRD patterns based on an ordered Li sublattice, where Li fully occupy 8a, 16f, and 32g sites (Fig. 6.2a), and a disordered Li sublattice, where Li partially occupy 24d and 96h sites (Fig. 6.2b). We then compared them with the pattern calculated from the compressed sample (Fig. 6.2c). The pattern from the Li sublattice in the compressed sample (Fig. 6.2c) matches well with the pattern from the one with an ordered Li sublattice (Fig. 6.2a). This confirms that the 3% compressive strain can induce an disorder-order transition in Al doped c-LLZO. The ordered Li sublattice then leads to the ionic diffusion slowing down. Ordered: 8a, 16f, 32g a = 1.039 c Disordered: 24d, 96h a = 1.039 c c-LLZO + compression a = 1.039 c Figure 6.2. Calculated virtual XRD patterns of Li sublattice in a tetragonal lattice. The lattice parameters is fixed to a / c = 1.039, corresponds to a 3% compressive strain loading. The sublattice orderings are a) ordered Li sublattice (fully occupied 8a, 16f, and 32g sites) b) disordered Li sublattice (partially occupied 24d and 96h sites) c) Al doped c-LLZO under 3% compressive strain loading and after equilibrium. The pattern in c) matches the pattern in a), confirming the Li sublattice in the deformed c-LLZO changes to an ordered one. 93 The coupling between the lattice parameters and sublattice ordering is probably the origin of this strain-induced order-disorder transition in lithium garnet.23 To further explore the strain-induced order-disorder transition in lithium garnet, we also studied the strain effect on undoped LLZO (Li7La3Zr2O12). Ordered: 8a, 16f, 32g a = c Disordered: 24d, 96h a = c t-LLZO + tension a = c Figure 6.3. Calculated virtual XRD patterns of Li sublattice in a cubic lattice. The lattice parameters are fixed to a / c = 1. The sublattice orderings are a) ordered Li sublattice (fully occupied 8a, 16f, and 32g sites) b) disordered Li sublattice (partially occupied 24d and 96h sites) c) from undoped t-LLZO under a tensile strain loading to cubic lattice parameters. The pattern in c) matches the pattern in b), confirming the Li sublattice in deformed t-LLZO changes to a disordered one. The undoped LLZO naturally stays at tetragonal phase with ordered Li sublattice below the transition temperature. We applied a tensile strain loading and deform the sample of undoped t-LLZO, forcing it to have cubic lattice parameters, i.e. a = b = c . The lattice parameters are extrapolated from the cubic phase lattice parameters of undoped LLZO 94 above the transition temperature (Fig. C1). Then we used the same method to calculated the virtual XRD pattern. Fig. 6.3 shows the XRD pattern of Li sublattice in the deformed t-LLZO sample with a cubic lattice. By comparing with the pattern from ordered Li sublattice (fully occupied 8a, 16f, and 32g sites, Fig. 6.3a), and the pattern from disordered Li sublattice (partially occupied 24d and 96h sites, Fig. 6.3b), we found that the pattern from t-LLZO sample under tensile loading (Fig. 6.3c) matches the one with disordered Li sublattice (Fig. 6.3b). Therefore, we confirmed that strain can also induce an order-disorder transition in undoped t-LLZO. 6.3 Discussion Our study revealed the effect of mechanical deformation on the ion diffusion in lithium garnet. Both tensile and compressive strain would decrease the ion diffusivity in c- LLZO. At moderate strain, the decrease in ion diffusion is minor in c-LLZO, and the activation energy does not show any significant change. A compressive deformation on c-LLZO changes its lattice into a tetragonal lattice. The lattice parameters are coupled with the Li sublattice ordering. Therefore, the compressive deformation can transform the original disordered Li sublattice into ordered sublattice similar to that in t-LLZO. The ordered Li sublattice results in a significant decrease in conductivity. The order-disorder transition induced by mechanical deformation can also be used to facilitate Li-ion diffusion. The undoped LLZO stays at tetragonal phase below the transition temperature, with an ordered Li sublattice and low ionic conductivity. By 95 applying a tensile mechanical deformation on t-LLZO and changing the lattice into a cubic lattice, a Li sublattice undergoes an order-disorder transition. The Li sublattice transform into a disordered sublattice, like that in c-LLZO. This sublattice change leads to a c-LLZO type ionic diffusion, which have significant higher ionic conductivity compared to the undeformed t-LLZO. 96 Chapter 7. Conclusion and Future Work 7.1 Conclusions The major conclusions of this dissertation are summarized below. 1) Most solid electrolytes have limited electrochemical stability window, especially the sulfide based solid electrolytes. Conventional CV experiments would overestimate the intrinsic electrochemical window of solid electrolytes. 2) The poor stability at the solid electrolyte–electrode interface indicates a strong thermodynamic driving force of decomposition, which leads to the spontaneous formation of interphase layers. 3) The interphase layers have a significant impact on the battery performance. MIEC interphase layer cannot provide passivation, which can lead to continuous decomposition and high interfacial resistance, while SEI-type interphase layer can provide the essential passivation and protect the solid electrolytes from further reduction. 4) For interface with poor stability, applying coating layer materials is a practical engineering strategy to improve the interface stability and to decrease the interfacial resistance. 5) Material stability against lithium metal is governed by cation and anion chemistry. Metal cations in oxides, sulfides, and fluorides are usually reduced by lithium metal, leading to poor stability against lithium metal. 6) Nitride anion chemistry shows unique stability against Li metal. In general, metal nitrides have a significantly lower reduction potential than oxides, sulfides, and fluorides. 97 7) Many nitrides that have intrinsic lithium metal stability can be used as buffer layers to protect materials from lithium metal reduction. Enriching nitrogen content at the interface with lithium metal can potential leads to the formation of a stable, passivating SEI, transforming a lithium-unstable interface into a lithium-stable interface. 8) Grain boundary slows down the ionic diffusion in lithium super-ionic conducting materials. 9) Strain can induce order-disorder transition in lithium garnet materials, and this transition has a significant effect on the ionic diffusion. 7.2 Future Work To improve the performance and reliability of all-solid-state Li-ion battery, interface stability is a critical factor that need to be considered. This dissertation revealed the intrinsic electrochemical window of solid electrolyte is narrow, and the formed interphase plays a crucial role in interfacial stability. However, there are more issues regarding the stability and interface issues in all-solid-state Li-ion batteries. A few possible future directions are listed below. 1) Despite the high ionic conductivity and easy processing, sulfide-based solid electrolytes are generally quite sensitive to moisture, posing difficulty for manufacturing process. The moisture sensitivity can also be assessed by computations. Such knowledge would be useful to improve the moisture resistivity of sulfide-based solid electrolytes. 98 2) For the grain boundary in solid electrolytes, I only looked into a few selected orientation grain boundaries within one typical solid electrolyte (lithium garnet). It might be worthwhile to look into grain boundaries with other orientations and terminations, or in other solid electrolyte materials, to get a more comprehensive understanding on the grain boundary diffusion in solid electrolytes. 3) In this dissertation, I observed the strain-induced order-disorder transition in lithium garnet. Such coupling between strain and order-disorder transition may not be only limited to lithium garnet. Other materials with order-disorder transition may also have similar behavior. It may be worth to investigate the strain effect on other super ionic conductor materials. 99 Publications, Patents, and Presentations Publications 1) Y. Zhu, X. He, Y. Mo, Origin of outstanding stability in the lithium solid electrolyte materials: insights from thermodynamic analyses based on first principles calculations. ACS Applied Materials & Interfaces, 7, 23685-23693, (2015) 2) Y. Zhu, X. He, Y. Mo, First principles study on electrochemical and chemical stability of solid electrolyte–electrode interfaces in all-solid-state Li-ion batteries. Journal of Materials Chemistry A, 4, 3253-3266, (2016) (Featured front cover, selected as the "2016 Journal of Materials Chemistry A HOT Papers” and editorial “most highly cited original research articles in 2016”) 3) Y. Zhu, X. He, Y. Mo, Strategies based on nitride materials chemistry to stabilize Li metal anode. Advanced Science, 1600517, (2017) 4) F. Han§, Y. Zhu§, X. He, Y. Mo, C. Wang, Electrochemical stability of Li10GeP2S12 and Li7La3Zr2O12 solid electrolytes. Advanced Energy Materials, 6, 1501590, (2016) (§ co first-authors) 5) X. He, Y. Zhu, Y. Mo, Origin of fast ion diffusion in super-ionic conductors. Nature Communication, 6, (2017) 6) K. H. Park, Q. Bai, D. H. Kim, D. Y. Oh, Y. Zhu, Y. Mo, Y. S. Jung, Design Strategies, Practical Considerations, and New Solution Processes of Sulfide Solid Electrolytes for All-Solid-State Batteries, Advanced Energy Materials, 8, 1800035, (2018) 100 7) X. Fan, E. Hu, J. Li, Y. Zhu, et al., High energy-density and reversibility of iron fluoride cathode enabled via an intercalation-extrusion reaction, Nature Communication, 9, 2324 (2018) 8) X. He, Y. Zhu, A. Epstein, Y. Mo, Statistical variances of diffusional properties from ab initio molecular dynamics simulations. npj Computational Materials. 4 (1), (2018) 9) K. He, S. Zhang, J. Li, X. Yu, Q. Meng, Y. Zhu, et al., Visualizing non-equilibrium lithiation of spinel oxide via in situ transmission electron microscopy, Nature Communication, 7, 11441 (2016) 10) W. Luo, Y. Gong, Y. Zhu, et al., Transition from super-lithiophobicity to super- lithiophilicity of garnet solid-state electrolyte, Journal of the American Chemical Society, 138 (37), 12258, (2016) 11) W. Luo, Y. Gong, Y. Zhu, et al., Reducing interfacial resistance between garnet- structured solid-state electrolyte and Li metal anode by a germanium Layer, Advanced Materials, 29 (22), 1606042, (2017) 12) K. He, F. Lin, Y. Zhu, et al., Sodiation kinetics of metal oxide conversion electrodes: a comparative study with lithiation. Nano Letters, 15, 5755-5763 (2015) 13) J. Li, K. He, Q. Meng, X. Li, Y. Zhu, et al., Kinetic phase evolution of spinel cobalt oxide during lithiation. ACS Nano, 10, 9577 (2016) 14) K. K. Fu, Y. Gong, B. Liu, Y. Zhu, et al., Toward garnet electrolyte–based Li metal batteries: An ultrathin, highly effective, artificial solid-state electrolyte/metallic Li interface. Science Advance, 3(4), 1601659 (2017) 101 15) K. K. Fu, Y. Gong, S. Xu, Y. Zhu, et al., Stabilizing the Garnet Solid- Electrolyte/Polysulfide Interface in Li-S Batteries. Chemistry of Materials, 29, 8037 (2017) 16) Q. Bai, Y. Zhu, X. He, E. Wachsman, Y. Mo, First principles hybrid functional study of small polarons in doped SrCeO3 perovskite: towards computation design of materials with tailored polaron. Ionics. 15, 5755-5763 (2017) 17) Q. Bai, X. He, Y. Zhu, Y. Mo, First Principles Study of Oxyhydride H-Ion Conductors: Toward Facile Anion Conduction in Oxide-Based Materials. ACS Applied Energy Materials.15, 5755-5763 (2018) Patents 1) C. Wang, Y. Mo, Y. Zhu, F. Han, X. Zhu, J. Yue, Novel Strategies to Stabilize Li Metal Anode Application Number: US 62/637,584. Presentations 1) Y. Zhu, X. He, Y. Mo., Atomistic Modeling of Interface Transport in All-Solid- State Li-ion batteries, Materials Research Society, Phoenix, Arizona (2018) 2) Y. Zhu, X. He, Y. Mo., Novel Strategies for Lithium Metal Anode Protection Based on Nitride Materials Chemistry: Insight from First Principles Calculations, 232nd Electrochemical Society Meeting, National Harbor, Maryland (2017) 3) Y. Zhu, X. He, Y. Mo., Electrochemical and Chemical Stability of Solid Electrolyte–Electrode Interfaces: A First Principles Computation Study, 232nd Electrochemical Society Meeting, National Harbor, Maryland (2017) 102 4) Y. Zhu, X. He, Y. Mo. “Novel Strategies for Lithium Metal Anode Protection based on Nitride Materials Chemistry”, American Chemical Society, Washington, DC. (2017) 5) Y. Zhu, X. He, Y. Mo. “Interface Engineering Guided by First Principles Computation – Towards Enabling Li metal In All-Solid-State Li-ion Batteries”, Materials Research Society, Boston (2016) 6) Y. Zhu, X. He, Y. Mo., First principles study of solid electrolyte–electrode interfaces in all-solid-state Li-ion batteries, The Society of Engineering Science, 53rd Annual Technical Meeting, Maryland (2016) 7) Y. Zhu, X. He, Y. Mo., “First Principles Study of Electrochemical and Chemical Stability of the Solid Electrolyte–Electrode Interfaces in All-Solid-State Li-Ion Batteries”, American Physics Society, Baltimore (2016) 8) Y. Mo, Y. Zhu., “Problems at the Electrolyte–Electrode Interfaces in All-Solid-State Li-Ion Batteries: Insight from First-Principles Computation”, Materials Research Society, Boston (2015) 103 Appendices Appendix A 1. Materials systems The crystal structures and the phase equilibria of the commonly used solid electrolyte materials investigated in this study are summarized in Table A1. The energies of most materials in this study were obtained from the MP database.34 The materials that were not included in MP database were calculated on the basis of the experimentally determined structures. For the structures with disordered site occupancies, we sampled 50-60 different configurations using the same method in the previous studies13,162 and chose the configuration with the lowest energy as the ground state structure. For example, the ground state structure of Li0.33La0.56TiO3 (LLTO) material was ordered from the experimental structure determined in ref. 163. The Li3PS4 phase based on the Pmn21 𝛾 phase164 was determined to have the lowest energy among its polymorphs and was added into the phase diagrams including Li-P-S compositions. For the materials with a non-zero energy above hull, we evaluated the electrochemical window of the material by reducing the energy above hull to zero. These solid electrolyte materials may be entropically stabilized by the disordering of Li and other ions as in the LGPS.13 For those materials, such as Li3.25Ge0.25P0.75S4, Li1.5Al0.5Ge1.5(PO4)3 and Li7P2S8I, that have no experimentally determined structures from the ICSD database, the materials entries were constructed as the phase equilibria at the given composition with a zero energy above hull. There is no experimentally determined crystal structure available for the LiPON material systems, which are typically amorphous with a range of compositions. We added 104 the crystalline structure of Li2PO2N reported in Ref. 165 as a representative of the LiPON materials systems. We also used Li2.88PO3.73N0.14 and Li2.98PO3.3N0.46 as typical representative LiPON phase. Table A1. Crystal structures and phase equilibria of solid electrolyte materials Structure E above hull Composition Acronym Phase equilibria used (meV/atom) Li10GeP2S12 LGPS mp-696138166 21 Li3PS4, Li4GeS4 Li3.25Ge0.25P0.75S4 - - Li3PS4, Li4GeS4 Li3PS4 164 0 Li3PS4 Li4GeS4 mp-30249167 0 Li4GeS4 Li7P3S11 mp-6417034 22 Li3PS4, P2S5 Li PS Cl 1686 5 83 Li3PS4, Li2S, LiCl Li7P2S8I - - Li3PS4, LiI Li2PO2N LiPON 165 0 Li2PO2N Li2O, La2O3, Li La 1697 3Zr2O12 LLZO 7 Li6Zr2O7 TiO2, Li4Ti5O12, Li0.33La0.56TiO3 LLTO 163 68 La2Ti2O7 LiTi2(PO4)3, Li3PO4, Li1.3Ti1.7Al0.3(PO4)3 LATP 170 29 AlPO4 AlPO4, Ge5O(PO4)6, Li1.5Al0.5Ge1.5(PO4)3 LAGP - - GeO2, Li4P2O7 Li3.5Zn0.25GeO4 LISICON 171 34 Li2ZnGeO4, Li4GeO4 105 2. Calculated plateaus for the lithiation and delithiation of solid electrolyte materials. Table A2. Calculated plateaus for the lithiation and delithiation of solid electrolyte materials. (a) Li10GeP2S12 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.28 -0.28 23.75 Li2S, Li15Ge4, Li3P 0.45 -0.45 22.25 Li9Ge4, Li2S, Li3P 0.56 -0.56 21 Li3P, Li2S, LiGe 0.87 -0.87 20 Li3P, Li2S, Ge 0.93 -0.93 16 Li2S, LiP, Ge 1.17 -1.17 14.86 Li3P7, Li2S, Ge 1.30 -1.30 14.29 Ge, Li2S, LiP7 1.62 -1.62 14 Li2S, P, Ge 1.71 -1.71 10 Li4GeS4, P, Li2S 0 Li10GeP2S12 2.14 -2.14 -4 GeS2, Li3PS4, S 2.31 -2.31 -10 P2S5, GeS2, S 106 (b) Li3.25Ge0.25P0.75S4 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.28 -0.28 23.75 Li2S, Li15Ge4, Li3P 0.45 -0.45 22.25 Li9Ge4, Li2S, Li3P 0.56 -0.56 21 Li3P, Li2S, LiGe 0.87 -0.87 20 Li3P, Li2S, Ge 0.93 -0.93 16 Li2S, LiP, Ge 1.17 -1.17 14.86 Li3P7, Li2S, Ge 1.30 -1.30 14.29 Ge, Li2S, LiP7 1.62 -1.62 14 Li2S, P, Ge 1.71 -1.71 10 Li4GeS4, P, Li2S 0 Li3.25Ge0.25P0.75S4 2.14 -2.14 -4 GeS2, Li3PS4, S 2.31 -2.31 -10 P2S5, GeS2, S 107 (c) Li4GeS4 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.28 -0.28 7.75 Li2S, Li15Ge4 0.45 -0.45 6.25 Li9Ge4, Li2S 0.56 -0.56 5 Li2S, LiGe 1.62 -1.62 4 Li2S, Ge 0 Li4GeS4 2.14 -2.14 -4 GeS2, S (d) Li6PS5Cl Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.87 -0.87 8 Li3P, Li2S, LiCl 0.93 -0.93 6 LiP, Li2S, LiCl 1.17 -1.17 5.43 Li3P7, Li2S, LiCl 1.30 -1.30 5.14 LiP7, Li2S, LiCl 1.71 -1.71 5 P, Li2S, LiCl 0 Li6PS5Cl 2.01 -2.01 -2 Li3PS4, S, LiCl 2.31 -2.31 -5 P2S5, S, LiCl 2.88 -2.88 -6 P2S5, S, PCl3 108 (e) Li3PS4 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.87 -0.87 8 Li3P, Li2S 0.93 -0.93 6 LiP, Li2S 1.17 -1.17 5.43 Li3P7, Li2S 1.30 -1.30 5.14 LiP7, Li2S 1.71 -1.71 5 P, Li2S 0 Li3PS4 2.31 -2.31 -3 P2S5, S (f) Li7P2S8I Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.87 -0.87 16 Li3P, Li2S, LiI 0.93 -0.93 12 LiP, Li2S, LiI 1.17 -1.17 10.86 Li3P7, Li2S, LiI 1.30 -1.30 10.29 LiP7, Li2S, LiI 1.71 -1.71 10 P, Li2S, LiI 0 Li7P2S8I 2.31 -2.31 -6 P2S5, S, LiI 2.47 -2.47 -7 P2S5, S, I 109 (g) Li7P3S11 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.87 -0.87 24 Li3P, Li2S 0.93 -0.93 18 LiP, Li2S 1.17 -1.17 16.29 Li3P7, Li2S 1.30 -1.30 15.43 LiP7, Li2S 1.71 -1.71 15 P, Li2S 2.03 -2.03 1.25 Li3PS4, P 2.15 -2.15 1.08 Li3PS4, P4S3 2.24 -2.24 0.67 Li3PS4, P4S7 2.28 -2.28 0.29 Li3PS4, P4S9 0 Li7P3S11 2.31 -2.31 -7 P2S5, S 110 (h) LiPON Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.01 -0.01 8 Li3P, Li3N, Li2O 0.61 -0.61 6 Li3P, Li7PN4, Li2O 0.68 -0.68 4 Li3P, LiPN2, Li2O 0 Li2PO2N 2.63 -2.63 -0.86 P3N5, Li4P2O7, N2 2.75 -2.75 -0.92 P4ON6, Li4P2O7, N2 2.77 -2.77 -1.36 P4ON6, LiPO3, N2 3.05 -3.05 -1.50 PON, LiPO3, N2 3.63 -3.63 -2 PON, P2O5, N2 (i) LLZO Potential 𝜙 ref to 𝜇:; ref to Li ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.004 0.004 8 Zr, La2O3, Li2O 0.048 -0.048 6.67 Zr3O, La2O3, Li2O 0 Li7La3Zr2O12 2.91 -2.91 -0.50 Li2O2, La2O3, Li6Zr2O7 3.17 -3.17 -3.50 Li2O2, La2O3, La2Zr2O7 3.30 -3.30 -7 O2, La2O3, La2Zr2O7 111 (j) LLTO Potential 𝜙 ref to 𝜇:; ref to Li ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.018 -0.018 3.67 Ti6O, La2O3, Li2O 0.12 -0.12 3.33 Ti3O, La2O3, Li2O 0.29 -0.29 1.67 Ti3O, La2O3, Li4TiO4 0.37 -0.37 1.57 Ti2O, La2O3, Li4TiO4 0.46 -0.46 0.92 LiTiO2, La2O3, Li4TiO4 0.50 -0.50 0.64 LiTiO2, La2TiO5, Li4TiO4 1.12 -1.12 0.56 LiTiO2, La2TiO5, Li2TiO3 1.14 -1.14 0.28 LiTi2O4, La2TiO5, Li2TiO3 1.63 -1.63 0.14 LiTi2O4, La2Ti2O7, Li2TiO3 Li7/6Ti11/6O4, La2Ti2O7, 1.749 -1.749 0.09 Li2TiO3 Li7/6Ti11/6O4, La2Ti2O7, 1.752 -1.752 0.01 Li4Ti5O12 0 Li0.33La0.56TiO3 3.71 -3.71 -0.33 O2, La2Ti2O7, TiO2 112 (k) LAGP Potential 𝜙 ref to 𝜇:; ref to Li ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.06 -0.06 38.25 Li2O, Li15Ge4, Li3P, Li9Al4 0.28 -0.28 35.63 Li2O, Li15Ge4, Li3P, Li5AlO4 0.45 -0.45 33.38 Li2O, Li9Ge4, Li3P, Li5AlO4 0.56 -0.56 31.50 Li2O, LiGe, Li3P, Li5AlO4 0.69 -0.69 30 Li2O, Ge, Li3P, Li5AlO4 0.75 -0.75 10 Li3PO4, Ge, Li3P, Li5AlO4 0.87 -0.87 8 Li3PO4, Ge, Li3P, LiAlO2 0.93 -0.93 7.50 Li3PO4, Ge, LiP, LiAlO2 1.17 -1.17 7.36 Li3PO4, Ge, Li3P7, LiAlO2 1.30 -1.30 7.29 Li3PO4, Ge, LiP7, LiAlO2 1.49 -1.49 7.25 Li3PO4, Ge, P, LiAlO2 1.94 -1.94 7 Li3PO4, Ge, P, LiAl5O8 2.64 -2.64 6 Li3PO4, Ge, AlPO4 2.70 -2.70 3.50 Ge, GeO2, Li4P2O7, AlPO4 0 Li1.5Al0.5Ge1.5(PO4)3 Ge5O(PO4)6, O2, Li4P2O7, 4.27 -4.27 -0.10 AlPO4 Ge5O(PO4)6, O2, GeP2O7, 4.31 -4.31 -1.50 AlPO4 113 (l) LATP Potential 𝜙 ref to 𝜇:; ref to Li ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.026 -0.026 30.30 Ti3P, TiAl, Li3P, Li2O 0.082 -0.082 29.60 Ti2P, TiAl, Li3P, Li2O 0.088 -0.088 29.38 Ti2P, TiAl2, Li3P, Li2O 0.23 -0.23 28.25 Ti2P, Li5AlO4, Li3P, Li2O 0.69 -0.69 24 Li4TiO4, Li5AlO4, Li3P, Li2O 0.75 -0.75 16 Li4TiO4, Li5AlO4, Li3P, Li3PO4 0.79 -0.79 14.80 Li4TiO4, LiAlO2, Li3P, Li3PO4 0.86 -0.86 8.33 Li4TiO4, LiAlO2, TiP, Li3PO4 1.49 -1.49 7.33 Li2TiO3, LiAlO2, TiP, Li3PO4 1.55 -1.55 6.52 LiTi2O4, LiAlO2, TiP, Li3PO4 1.71 -1.71 6.35 LiTi2O4, LiAl5O8, TiP, Li3PO4 1.84 -1.84 5.82 TiO2, LiAl5O8, TiP, Li3PO4 1.94 -1.94 5.43 TiO2, LiAl5O8, TiP2, Li3PO4 1.98 -1.98 4.76 TiO2, AlPO4, TiP2, Li3PO4 2.12 -2.12 4.25 TiO2, AlPO4, P, Li3PO4 2.17 -2.17 2.13 LiTiPO5, AlPO4, P, Li3PO4 0 Li1.3Ti1.7Al0.3P3O12 4.21 -4.21 -0.15 LiTi2(PO4)3, AlPO4, Li4P2O7, O2 4.33 -4.33 -0.60 LiTi2(PO4)3, AlPO4, TiP2O7, O2 114 4.66 -4.66 -1.30 Ti5P4O20, AlPO4, TiP2O7, O2 (m) LISICON Potential 𝜙 ref to 𝜇:; ref to Li ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.18 -0.18 8.50 Li2O, Li15Ge4, LiZn 0.28 -0.28 7.85 Li2O, Li15Ge4, Li8Zn2Ge3 0.42 -0.42 6.91 Li2O, Li9Ge4, Li8Zn2Ge3 0.45 -0.45 6.69 Li2O, Li9Ge4, Li2ZnGe 0.56 -0.56 5.75 Li2O, LiGe, Li2ZnGe 0.61 -0.61 5 Li2O, Ge, Li2ZnGe 0.73 -0.73 4.75 Li2O, Ge, LiZnGe 1.02 -1.02 4.50 Li2O, Ge, Zn 1.44 -1.44 0.50 Li4GeO4, Zn 0 Li3.5Zn0.25GeO4 3.39 -3.39 -1.50 Li2ZnGeO4, Li2GeO3, O2 3.84 -3.84 -2.40 Li2ZnGeO4, Li4Ge5O12, O2 3.85 -3.85 -2.79 Li2ZnGeO4, Li2Ge7O15, O2 LiZn1.5Ge3O8, Li2Ge7O15, 3.89 -3.89 -3.19 O2 3.98 -3.98 -3.33 LiZn1.5Ge3O8, GeO2, O2 4.00 -4.00 -3.50 Zn2GeO4, GeO2, O2 115 3. Calculated plateaus for the lithiation and delithiation of coating layer materials. Table A3. Calculated plateaus for the lithiation and delithiation of coating layer materials. (a) LiTaO3 Potential 𝜙 ref to 𝜇:; ref to Li metal ∆𝑛:; per Phase equilibria Li/Li+(V) (eV) formula 0.35 -0.35 5 Ta, Li2O 0.54 -0.54 2 Ta, Li5TaO5 1.18 -1.18 1.25 Ta, Li3TaO4 0 LiTaO3 3.95 -3.95 -0.67 LiTa3O8, O2 4.12 -4.12 -1 Ta2O5, O2 (b) LiNbO3 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.55 -0.55 5 Nb, Li2O 0.62 -0.62 2 LiNbO2, Li2O 0.94 -0.94 1.2 LiNbO2, Li8Nb2O9 1.74 -1.74 1 LiNbO2, Li3NbO4 0 LiNbO3 3.88 -3.88 -0.67 LiNb3O8, O2 3.92 -3.92 -1 Nb2O5, O2 116 (c) Li4Ti5O12 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.018 -0.018 18.33 Ti6O, Li2O 0.12 -0.12 16.67 Ti3O, Li2O 0.29 -0.29 7.27 Ti3O, Li4TiO4 0.37 -0.37 6.86 Ti2O, Li4TiO4 0.50 -0.50 4 LiTiO2, Li4TiO4 1.12 -1.12 3 LiTiO2, Li2TiO3 1.63 -1.63 1.5 LiTi2O4, Li2TiO3 Li7/6Ti11/6O4, 1.75 -1.75 1 Li2TiO3 0 Li4Ti5O12 3.71 -3.71 -4 O2, TiO2 117 (d) Li2SiO3 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per Phase equilibria Li/Li+(V) formula 0.10 -0.10 8.2 Li21Si5, Li2O 0.23 -0.23 7.25 Li13Si4, Li2O 0.26 -0.26 3.62 Li13Si4, Li8SiO6 0.26 -0.26 1.81 Li13Si4, Li4SiO4 0.37 -0.37 1.43 Li12Si7, Li4SiO4 0.39 -0.39 1.25 LiSi, Li4SiO4 0.76 -0.76 1 Si, Li4SiO4 0 Li2SiO3 3.74 -3.74 -1 Li2Si2O5, O2 3.78 -3.78 -2 SiO2, O2 (e) Li3PO4 Potential 𝜙 ref to 𝜇:; ref to Li (eV) ∆𝑛:; per formula Phase equilibria Li/Li+(V) 0.69 -0.69 8 Li3P, Li2O 0 Li3PO4 4.21 -4.21 -1 Li4P2O7, O2 4.33 -4.33 -2 LiPO3, O2 4.99 -4.99 -3 P2O5, O2 118 Appendix B 1. List of Compounds. The lithium ternary compounds Li-M-X (X = N, O, S, F) investigated in Fig. 4.1 are summarized in Table 4.3. The structures and energies of most compounds were obtained from the Materials Project (MP) database. We added 26 additional lithium ternary compounds (Table 4.4), of which 21 were predicted using substitution prediction algorithm and 5 were from Inorganic Crystal Structure Database (ICSD). Table B1. Lithium ternary compounds Li-M-X (X = N, O, S, F) Cation Ternary Compounds Li-M-X Cation Ternary Compounds Li-M-X Er3+ LiErO2, LiErS2 Tm3+ LiTmO2 Ho3+ LiHoO2, LiHoS2 Y3+ LiYO2, LiYS2, LiYF4 Gd3+ LiGdO2 Sm3+ LiSmS2 Ca2+ LiCaN Be2+ LiBeN, Li2BeF4 Sr2+ LiSrN, Li4SrN2 Mg2+ LiMgN Li2HfN2, Li2HfO3, Li6Hf2O7, Sc3+ Li3ScN2, LiScO2, LiScS2 Hf4+ Li2HfF6 Li2ZrN2, Li6Zr2O7, Li2ZrO3, Zr4+ Li2ZrS3, Li4ZrF8, Li2ZrF6, Na+ LiNaS Li3Zr4F19 K+ LiKS Ce4+ Li2CeN2, Li2CeO3, LiCeF5 119 Li5SiN3, Li2SiN2, LiSi2N3 Li3AlN2, LiAl5O8, LiAlO2, Li2SiO3, Li2Si2O5, Li8SiO6, Al3+ Li5AlO4, LiAlS2, Li5AlS4, Si4+ Li4SiO4, Li8SiS6, Li2SiS3, Li3AlF6 Li2SiF6 LiZnN, Li10Zn4O9, Li6ZnO4, Zn2+ In3+ LiInO2, LiInF4 Li2ZnF4 Li4TaN3, Li7TaN4, LiTaN2, Li2TiN2, Li5TiN3, Li4TiO4, Ta5+ LiTa3O8, Li5TaO5, LiTaO3, Ti4+ Li2TiO3, Li4Ti5O12, Li2TiS3, Li3TaO4, Li3TaS4, LiTaF6 Li2TiF6 Li3BN2, Li2B4O7, LiBO2, Li2SnN2, Li8SnO6, Li2SnO3, Sn4+ B3+ Li3BO3, Li3B7O12, Li3BS3, Li2SnS3, Li2SnF6 Li5B7S13, LiBF4 Li3GaN2 LiGaO2, Li5GaO4, Ga3+ Ni2+ Li2NiF4 LiGa5O8, LiGaS2, Li3GaF6 Li7NbN4, LiNbN2, LiNbO3, Li3FeN2, LiFeO2, Li5FeO4, Nb5+ LiNb3O8, Li3NbO4, Li8Nb2O9 Fe3+ Li3FeF6 Li3NbS4, LiNbF6 Li5GeN3, Li2GeN2, LiGe2N3, Li2Ge7O15, Li4GeO4, Ge4+ W6+ Li6WN4, Li2WO4, Li4WO5 Li4Ge5O12, Li2GeO3 Li2GeS3, Li4GeS4, Li2GeF6 120 LiPN2, Li7PN4, Li4P2O7, P5+ Mo6+ Li6MoN4, Li2MoO4, Li4MoO5 LiPO3, Li3PO4, Li3PS4, LiPF6 Li7VN4, LiVN2, Li3VO4, LiMn2N3, Li2MnN2, Li2MnO3, V5+ Mn4+ LiVO3, Li3VS4, LiVF6 Li4Mn5O12, Li2MnF6 Co4+ Li8CoO6 Table B2. Additional lithium ternary compounds. Form. E Space Compound Structure source (eV/atom) group LiAlS2 -1.614 P41212 S-O substitution from LiAlO2 (ICSD-23815) Li5AlS4 -1.583 Pcab S-O substitution from Li5AlO4 (ICSD-42697) Li8SiS6 -1.496 P63cm S-O substitution from Li8SiO6 (ICSD-65176) Li2SiS3 -1.385 Ccm21 S-O substitution from Li2SiO3 (ICSD-853) Li2SiF6 -3.423 P321 ICSD-425923 Li2GeN2 -0.609 P-3m1 Ge-Zr substitution from Li2ZrN2 (ICSD-78790) LiGe2N3 -0.516 Cmc21 Ge-Si substitution from LiSi2N3 (ICSD-34118) Li2GeS3 -1.218 Ccm21 S-O substitution from Li2GeO3 (ICSD-100403) Li2SnN2 -0.403 P-3m1 Sn-Zr substitution from Li2ZrN2 (ICSD-78790) Li2SnS3 -1.211 C2/c S-O substitution from Li2SnO3 (ICSD-21032) Li2SnF6 -3.043 P-31m Sn-Zr substitution from Li2ZrF6 (ICSD-2644) Li2TiN2 -1.194 P-3m1 Ti-Zr substitution from Li2ZrN2 (ICSD-78790) Li2TiS3 -1.668 C2/c S-O substitution from Li2TiO3 (ICSD-15150) 121 Li2ZrS3 -1.800 C2/c S-O substitution from Li2ZrO3 (ICSD-94893) Li-Na, Nb-Ta substitution from NaTaN2 (ICSD- LiNbN2 -1.035 R-3m 67347) LiNbF6 -3.414 R-3 ICSD-165202 LiTaN2 -1.233 R-3m Li-Na substitution from NaTaN2 (ICSD-67347) Li3TaS4 -1.526 C2/c S-O substitution from Li3TaO4 (ICSD-281301) LiTaF6 -3.598 R-3 ICSD-165205 Li2ZnF4 -2.941 Fd-3m Zn-Ni substitution from Li2NiF4 (ICSD-202859) Li2HfF6 -3.788 P-31m ICSD-251074 Li2HfN2 -1.353 P-3m1 Hf-Zr substitution from Li2ZrN2 (ICSD-78790) LiCeF5 -3.830 I41/a ICSD-426211 Li-Mg, V-Si substitution from MgSiN2 (ICSD- LiVN2 -1.024 Pna21 23501) LiMn2N3 -0.496 Cmc21 Mn-Si substitution from LiSi2N3 (ICSD-78790) Li2MnN2 -0.663 Pbca Mn-Si substitution from Li2SiN2 (ICSD-420126) 2. Lithiation/delithiation reaction of materials Table B3. Lithiation and delithiation reactions of typical Al-abiding ternary nitride, oxide, sulfide and fluoride. 122 a) Li3AlN2 Potential ϕ (V) µLi (eV) ΔnLi Phase equilibria ref. to Li/Li+ ref. to Li metal per formula -0.316 0.316 5.25 Li9Al4, Li3N / / 0 Li3AlN2 0.789 -0.789 -2.667 AlN, LiN3 1.668 -1.668 -3 AlN, N2 b) LiAlO2 Potential ϕ (V) µLi (eV) ΔnLi Phase equilibria ref. to Li/Li+ ref. to Li metal per formula 0.056 -0.056 5.25 Li9Al4, Li2O 0.071 -0.071 2.625 Li9Al4, Li5AlO4 0.167 -0.167 2.25 Li3Al2, Li5AlO4 / / 0 LiAlO2 3.697 -3.697 -0.8 LiAl5O8, O2 4.101 -4.101 -1 Al2O3, O2 123 c) LiAlS2 Potential ϕ (V) µLi (eV) ΔnLi Phase equilibria ref. to Li/Li+ ref. to Li metal per formula 0.071 -0.071 5.25 Li9Al4, Li2S 0.190 -0.190 4.5 Li3Al2, Li2S 0.359 -0.359 4 LiAl, Li2S 0.963 -0.963 3 Al, Li2S 0.973 -0.973 1.5 Li5AlS4, Al / / 0 LiAlS2 2.812 -2.812 -1 Al2S3, S d) Li3AlF6 Potential ϕ (V) µLi (eV) ΔnLi Phase equilibria ref. to Li/Li+ ref. to Li metal per formula 0.071 -0.071 5.25 Li9Al4, LiF 0.190 -0.190 4.5 Li3Al2, LiF 0.359 -0.359 4 LiAl, LiF 1.057 -1.057 3 LiF, Al / / 0 Li3AlF6 6.477 -6.477 -3 AlF3, F2 124 Table B4. Cathodic limits and lithiation reactions for all binary and ternary fluorides, sulfides, oxides and nitrides plotted in Fig. 4.1. The reaction energy ED normalized to per Li inserted is for the lithiation of selected compound to from the phase equilibria with Li metal. For materials that are stable against Li metal (i.e., with negative cathodic limit), the “Phase equilibria at Li” column lists the fully lithiated products after lithiation with overpotential. 125 Cathodic Phase Phase equilibria ED /eV Compound limit /V equilibria at cathodic limit norm. to Li ref. to Li/Li+ at Li ErN -0.699 Li3N, Er Li3N, Er 0.699 ErF3 0.371 Er, LiF LiF, Er -0.371 Er2S3 0.991 LiErS2, ErS Er, Li2S -0.284 LiErS2 0.484 Li2S, ErS Er, Li2S -0.200 Er2O3 -0.239 LiErO2, Er Er, Li2O 0.288 LiErO2 -0.304 Er, Li2O Er, Li2O 0.304 TmN -0.728 Li3N, Tm Li3N, Tm 0.728 TmF3 0.837 Tm, LiF LiF, Tm -0.837 Tm2S3 0.660 TmS, Li2S Li2S, Tm -0.276 Tm2O3 -0.238 LiTmO2, Tm Tm, Li2O 0.316 LiTmO2 -0.342 Li2O, Tm Tm, Li2O 0.342 HoN -0.675 Li3N, Ho Li3N, Ho 0.675 HoF3 0.378 LiF, Ho LiF, Ho -0.378 Ho2S3 0.997 LiHoS2, HoS Ho, Li2S -0.280 LiHoS2 0.541 Li2S, HoS Ho, Li2S -0.204 Ho2O3 -0.237 LiHoO2, Ho Ho, Li2O 0.270 LiHoO2 -0.280 Ho, Li2O Ho, Li2O 0.280 YN -0.639 Y, Li3N Li3N, Y 0.639 YF3 0.625 LiYF4, Y Y, LiF -0.429 126 LiYF4 0.364 Y, LiF LiF, Y -0.364 Y2S3 1.167 Y5S7, LiYS2 Y, Li2S -0.278 LiYS2 0.600 Li2S, YS Y, Li2S -0.209 Y2O3 -0.204 LiYO2, Y Y, Li2O 0.215 LiYO2 -0.219 Y, Li2O Y, Li2O 0.219 LaN -0.398 Li3N, La Li3N, La 0.398 LaF3 0.339 La, LiF LiF, La -0.339 La2S3 0.846 LaS, Li2S La, Li2S -0.174 La2O3 -0.139 La, Li2O La, Li2O 0.139 GdN -0.559 Li3N, Gd Li3N, Gd 0.559 GdF3 0.856 LiF, Gd LiF, Gd -0.856 Gd2S3 0.935 GdS, Li2S Li2S, Gd -0.281 Gd2O3 -0.132 LiGdO2, Gd Li2O, Gd 0.149 LiGdO2 -0.154 Gd, Li2O Li2O, Gd 0.154 SmN -0.500 Li3N, Sm Li3N, Sm 0.500 SmF3 0.404 LiF, Sm Sm, LiF -0.404 Sm2S3 0.906 LiSmS2, SmS Sm, Li2S -0.283 LiSmS2 0.858 Li2S, SmS Sm, Li2S -0.275 Sm2O3 -0.126 Sm, Li2O Li2O, Sm 0.126 Ca3N2 0.263 Ca2N, LiCaN Li3N, Li2Ca 0.075 LiCaN -0.111 Li3N, Li2Ca Li2Ca, Li3N 0.111 CaF2 0.009 LiF, Li2Ca Li2Ca, LiF -0.009 127 CaS -0.061 Li2Ca, Li2S Li2S, Li2Ca 0.061 CaO -0.096 Li2Ca, Li2O Li2Ca, Li2O 0.096 NdN -0.418 Li3N, Nd Li3N, Nd 0.418 NdF3 0.441 Nd, LiF LiF, Nd -0.441 Nd2S3 0.900 Li2S, NdS Li2S, Nd -0.288 Nd2O3 -0.059 Nd, Li2O Nd, Li2O 0.059 Be3N2 -0.069 LiBeN, Be Be, Li3N 0.414 LiBeN -0.587 Be, Li3N Li3N, Be 0.587 BeF2 1.119 Li2BeF4, Be Be, LiF -0.999 Li2BeF4 0.879 Be, LiF Be, LiF -0.879 BeS 0.943 Li2S, Be Be, Li2S -0.943 BeO -0.015 Be, Li2O Be, Li2O 0.015 SrLiN 0.237 SrLi4N2, Sr2N Li3N, Sr3Li2 0.011 SrLi4N2 -0.054 Sr3Li2, Li3N Li3N, Sr3Li2 0.054 SrF2 -0.013 Sr3Li2, LiF Sr3Li2, LiF 0.013 SrS -0.105 Sr3Li2, Li2S Sr3Li2, Li2S 0.105 SrO 0.013 Li2O, Sr Li2O, Sr3Li2 -0.011 Mg3N2 0.064 LiMg2, LiMgN Li3N, Li5Mg 0.022 LiMgN -0.047 Li3N, Li5Mg Li3N, Li5Mg 0.047 MgF2 0.585 Mg, LiF Li5Mg, LiF -0.186 MgS 0.575 Mg, Li2S Li5Mg, Li2S -0.184 MgO 0.065 LiMg2, Li2O Li2O, Li5Mg -0.030 128 ScN -0.689 Li3ScN2, Sc Sc, Li3N 0.790 Li3ScN2 -0.892 Li3N, Sc Li3N, Sc 0.892 ScF3 0.600 Sc, LiF Sc, LiF -0.600 Sc2S3 1.252 LiScS2, ScS Sc, Li2S -0.410 LiScS2 0.664 ScS, Li2S Sc, Li2S -0.312 Sc2O3 0.148 LiScO2, Sc Sc, Li2O 0.215 LiScO2 -0.336 Sc, Li2O Li2O, Sc 0.336 Hf3N4 0.667 HfN, Li2HfN2 Li3N, Hf 0.486 Li2HfN2 -0.771 Li3N, Hf Li3N, Hf 0.771 HfF4 1.481 Li2HfF6, Hf LiF, Hf -1.173 Li2HfF6 1.019 Hf, LiF LiF, Hf -1.019 HfS2 0.942 Li2S, Hf2S Hf, Li2S -0.821 HfO2 0.457 Hf, Li2HfO3 Hf, Li2O -0.077 Li6Hf2O7 -0.153 Li2O, Hf Li2O, Hf 0.153 Li2HfO3 0.122 Li6Hf2O7, Hf Li2O, Hf 0.113 Zr3N4 0.956 Li2ZrN2, ZrN Li3N, Zr 0.374 Li2ZrN2 -0.603 Zr2N, Li3N Zr, Li3N 0.651 ZrF4 1.619 Li3Zr4F19, Zr LiF, Zr -1.331 Li4ZrF8 1.209 LiF, Zr LiF, Zr -1.209 Li2ZrF6 1.239 Li4ZrF8, Zr LiF, Zr -1.216 Li3Zr4F19 1.510 Li2ZrF6, Zr Zr, LiF -1.277 ZrS2 1.705 Li(ZrS2)2 Li2S, Zr -0.870 129 Li2ZrS3 1.227 Li(ZrS2)2, Li2S Zr, Li2S -0.810 ZrO2 0.583 Li2ZrO3, Zr3O Zr, Li2O -0.232 Li6Zr2O7 0.048 Li2O, Zr3O Zr, Li2O -0.041 Li2ZrO3 0.339 Li6Zr2O7, Zr3O Zr, Li2O -0.079 NaF 0.445 LiF, Na Na, LiF -0.445 Na2S 0.417 Na, NaLiS Na, Li2S -0.399 NaLiS 0.381 Na, Li2S Na, Li2S -0.381 Na2O 0.925 Li2O, Na Na, Li2O -0.925 KF 0.449 LiF, K LiF, K -0.449 K2S 0.488 KLiS, K K, Li2S -0.410 KLiS 0.333 Li2S, K Li2S, K -0.333 K2O 1.209 Li2O, K K, Li2O -1.209 Li2CeN2 -0.451 CeN, Li3N Ce, Li3N 0.495 CeF4 3.228 LiCeF5, CeF3 Ce, LiF -1.257 LiCeF5 3.158 CeF3, LiF Ce, LiF -1.248 CeS2 1.955 Ce2S3, Li2S Li2S, Ce -0.734 CeO2 1.225 Ce11O20, Li2CeO3 Li2O, Ce -0.147 Li2CeO3 1.174 Ce11O20, Li2O Ce, Li2O -0.144 AlN -0.004 Li3AlN2, Li9Al4 Li3N, Li9Al4 0.160 Li3AlN2 -0.316 Li3N, Li9Al4 Li3N, Li9Al4 0.316 AlF3 1.289 Li3AlF6, Al Li9Al4, LiF -0.767 Li3AlF6 1.057 Al, LiF Li9Al4, LiF -0.701 130 Al2S3 1.597 Al, LiAlS2 Li9Al4, Li2S -0.740 LiAlS2 0.973 Li5AlS4, Al Li9Al4, Li2S -0.650 Li5AlS4 0.963 Al, Li2S Li9Al4, Li2S -0.647 Al2O3 1.233 LiAl5O8, Al Li9Al4, Li2O -0.233 LiAlO2 0.167 Li5AlO4, Li3Al2 Li9Al4, Li2O -0.105 Li5AlO4 0.056 Li2O, Li9Al4 Li9Al4, Li2O -0.056 LiAl5O8 0.802 LiAlO2, Al Li9Al4, Li2O -0.195 Si3N4 1.127 LiSi2N3, Si Li21Si5, Li3N -0.067 LiSi2N3 0.245 Li2SiN2, Li13Si4 Li3N, Li21Si5 0.011 Li2SiN2 0.022 Li5SiN3, Li21Si5 Li3N, Li21Si5 0.091 Li5SiN3 -0.148 Li3N, Li21Si5 Li3N, Li21Si5 0.148 SiF4 2.298 Si, Li2SiF6 Li21Si5, LiF -1.112 Li2SiF6 1.839 Si, LiF LiF, Li21Si5 -1.038 SiS2 1.633 Li2SiS3, Si Li21Si5, Li2S -0.872 Li2SiS3 1.455 Li8SiS6, Si Li21Si5, Li2S -0.840 Li8SiS6 1.411 Si, Li2S Li21Si5, Li2S -0.829 SiO2 1.320 Li2Si2O5, Si Li21Si5, Li2O -0.454 Li2SiO3 0.761 Li4SiO4, Si Li21Si5, Li2O -0.295 Li2Si2O5 1.265 Li2SiO3, Si Li21Si5, Li2O -0.372 Li4SiO4 0.258 Li13Si4, Li8SiO6 Li21Si5, Li2O -0.222 Li8SiO6 0.227 Li13Si4, Li2O Li2O, Li21Si5 -0.213 Zn3N2 1.423 Zn, LiN3 LiZn, Li3N -0.532 131 LiZnN 0.380 Li3N, LiZn3 LiZn, Li3N -0.365 ZnF2 2.441 Zn, Li2ZnF4 LiZn, LiF -1.764 Li2ZnF4 2.428 Zn, LiF LiZn, LiF -1.760 ZnS 1.200 Zn, Li2S LiZn, Li2S -0.941 ZnO 1.364 Li10Zn4O9, Zn Li2O, LiZn -1.014 Li6ZnO4 1.179 Li2O, Zn LiZn, Li2O -0.927 Li10Zn4O9 1.318 Li6ZnO4, Zn LiZn, Li2O -0.968 InN 1.037 LiN3, In Li3N, Li13In3 -0.376 InF3 2.625 In, LiInF4 LiF, Li13In3 -1.195 LiInF4 2.517 LiF, In LiF, Li13In3 -1.184 In2S3 1.634 InS, Li2S Li13In3, Li2S -0.785 In2O3 1.579 LiInO2, In Li13In3, Li2O -0.734 LiInO2 1.363 In, Li2O Li13In3, Li2O -0.712 EuN -0.051 Li3N, Eu Li3N, Eu 0.051 EuF3 3.389 EuF2, LiF Eu, LiF -1.278 Eu2O3 1.655 EuO, Li2O Li2O, Eu -0.531 Ta3N5 1.371 Li2Ta3N5 Ta, Li3N 0.073 Li4TaN3 -0.016 Li7TaN4, Ta Li3N, Ta 0.583 Li7TaN4 -0.771 Li3N, Ta Li3N, Ta 0.771 LiTaN2 0.877 Li2Ta3N5, Li4TaN3 Li3N, Ta 0.250 TaF5 2.895 Ta, LiTaF6 LiF, Ta -2.045 LiTaF6 1.876 LiF, Ta LiF, Ta -1.876 132 Li3TaS4 1.853 Li2S, Li2(TaS2)3 Li2S, Ta -1.302 Ta2O5 1.694 LiTa3O8, Ta Li2O, Ta -0.758 LiTaO3 1.175 Li3TaO4, Ta Li2O, Ta -0.583 Li3TaO4 0.543 Li5TaO5, Ta Ta, Li2O -0.386 LiTa3O8 1.591 LiTaO3, Ta Ta, Li2O -0.695 Li5TaO5 0.347 Ta, Li2O Li2O, Ta -0.347 Li5TiN3 -0.828 Li3N, Ti2N Li3N, Ti 0.829 Li2TiN2 0.712 TiN, Li5TiN3 Li3N, Ti 0.572 TiF4 2.928 TiF3, Li2TiF6 Ti, LiF -1.725 Li2TiF6 1.896 TiF3, LiF Ti, LiF -1.553 TiS2 1.939 Li(TiS2)3 Ti, Li2S -1.050 Li2TiS3 1.717 Li2S, LiTiS2 Ti, Li2S -1.009 TiO2 1.712 LiTi2O4 Ti, Li2O -0.468 Li2TiO3 0.500 LiTiO2, Li4TiO4 Ti, Li2O -0.167 Li4TiO4 0.123 Ti3O, Li2O Ti, Li2O -0.095 Li4Ti5O12 1.749 Li7Ti11O24, Li2TiO3 Ti, Li2O -0.347 Li17Sn4, Li2SnN2 0.473 Li3N, Li13Sn5 -0.424 Li3N SnF4 4.319 SnF3, Li2SnF6 LiF, Li17Sn4 -1.651 Li2SnF6 2.911 SnF2, LiF LiF, Li17Sn4 -1.528 SnS2 1.945 Li2SnS3, SnS Li2S, Li17Sn4 -1.058 Li2SnS3 1.822 SnS, Li2S Li17Sn4, Li2S -1.043 133 Li17Sn4, SnO2 1.986 Li2SnO3, Sn5O6 -0.955 Li2O Li17Sn4, Li8SnO6 1.243 Li2O, Sn -0.824 Li2O Li17Sn4, Li2SnO3 1.362 Li8SnO6, Sn -0.853 Li2O BN 0.062 Li3BN2, LiB Li3N, LiB 0.195 Li3BN2 -0.453 Li3N, LiB LiB, Li3N 0.453 BF3 2.933 B, LiBF4 LiF, LiB -1.720 LiBF4 1.937 B, LiF LiB, LiF -1.534 B2S3 1.870 B, Li5B7S13 Li2S, LiB -1.319 Li3BS3 1.589 B, Li2S LiB, Li2S -1.272 Li5B7S13 1.615 Li3BS3, B Li2S, LiB -1.280 B2O3 2.067 B6O, Li3B7O12 LiB, Li2O -0.652 Li3BO3 0.284 Li7B18O, Li2O Li2O, LiB -0.249 Li2B4O7 1.291 LiBO2, B6O Li2O, LiB -0.508 Li3B7O12 1.294 Li2B4O7, B6O LiB, Li2O -0.519 LiBO2 1.014 B6O, Li3BO3 Li2O, LiB -0.436 GaN 0.484 Li3GaN2, LiGa Li2Ga, Li3N -0.281 Li3GaN2 0.125 Li3N, Li2Ga Li2Ga, Li3N -0.125 GaF3 2.581 Li3GaF6, Ga Li2Ga, LiF -1.637 Li3GaF6 2.277 Ga, LiF Li2Ga, LiF -1.546 134 Ga2S3 2.099 LiGaS2, GaS Li2Ga, Li2S -1.053 LiGaS2 1.307 Ga, Li2S Li2Ga, Li2S -0.964 Ga2O3 2.131 LiGa5O8, Ga Li2Ga, Li2O -0.902 LiGaO2 1.046 Li5GaO4, Ga Li2Ga, Li2O -0.752 LiGa5O8 1.892 Ga, LiGaO2 Li2Ga, Li2O -0.865 Li5GaO4 0.862 Ga, Li2O Li2Ga, Li2O -0.697 NiF2 3.200 Ni, Li2NiF4 Ni, LiF -3.165 Li2NiF4 3.130 Ni, LiF Ni, LiF -3.130 NiO 2.172 Li2O, Ni Li2O, Ni -2.172 Li7NbN4 -0.639 Li3N, Nb Nb, Li3N 0.639 LiNbN2 0.787 NbN, Li7NbN4 Li3N, Nb 0.092 NbF5 3.169 Nb2F5, LiNbF6 Nb, LiF -2.309 LiNbF6 2.204 Nb2F5, LiF Nb, LiF -2.171 Li3NbS4 2.083 Li2S, Li5(NbS2)7 Nb, Li2S -1.378 Nb2O5 2.286 LiNb3O8, Nb12O29 Li2O, Nb -0.970 LiNbO3 1.745 Li3NbO4, LiNbO2 Li2O, Nb -0.814 LiNb3O8 2.225 Nb12O29, LiNbO3 Nb, Li2O -0.917 Li8Nb2O9 0.623 LiNbO2, Li2O Nb, Li2O -0.577 Li3NbO4 0.937 Li8Nb2O9, LiNbO2 Nb, Li2O -0.603 FeN 0.630 Fe3N, Li3FeN2 Li3N, Fe -0.405 Li3FeN2 0.188 Li3N, Fe Fe, Li3N -0.188 FeF3 3.655 LiFe2F6 Fe, LiF -2.663 135 Li3FeF6 3.151 FeF2, LiF Fe, LiF -2.553 Fe2O3 2.299 Fe3O4, LiFeO2 Fe, Li2O -1.535 Li5FeO4 1.282 Li2O, Fe Fe, Li2O -1.282 LiFeO2 1.540 Li2FeO2 Fe, Li2O -1.373 Li3N, Ge3N4 1.728 LiGe2N3, Ge -0.438 Li15Ge4 Li15Ge4, Li5GeN3 0.198 Li15Ge4, Li3N -0.198 Li3N Li3N, LiGe2N3 0.735 Li2GeN2, Ge -0.356 Li15Ge4 Li15Ge4, Li2GeN2 0.500 Li5GeN3, LiGe -0.282 Li3N GeF4 3.773 Ge5F12, Li2GeF6 LiF, Li15Ge4 -1.735 Li2GeF6 2.710 LiF, Ge Li15Ge4, LiF -1.598 GeS2 1.913 GeS, Li2GeS3 Li2S, Li15Ge4 -1.107 Li4GeS4 1.622 Li2S, Ge Li2S, Li15Ge4 -1.037 Li2GeS3 1.865 Li4GeS4, Ge Li15Ge4, Li2S -1.068 Li15Ge4, GeO2 2.408 Li2Ge7O15, Ge -0.992 Li2O Li2O, Li4GeO4 1.016 Ge, Li2O -0.724 Li15Ge4 136 Li2O, Li2Ge7O15 2.275 Li4Ge5O12, Ge -0.960 Li15Ge4 Li2O, Li4Ge5O12 2.259 Li2GeO3, Ge -0.911 Li15Ge4 Li2O, Li2GeO3 1.579 Ge, Li4GeO4 -0.797 Li15Ge4 CuF2 3.638 LiF, Cu LiCu3, LiF -3.136 CuS 1.821 Cu7S4, Li2S LiCu3, Li2S -1.569 CuO 2.411 LiCuO LiCu3, Li2O -1.863 WN2 2.056 Li5W7N12, N2 Li3N, W -0.353 Li6WN4 -0.465 Li3N, W Li3N, W 0.465 WF6 3.593 LiWF6 W, LiF -3.067 WO3 2.564 Li2WO4, W18O49 W, Li2O -1.650 Li2WO4 2.097 Li4WO5, W Li2O, W -1.394 Li4WO5 1.219 W, Li2O W, Li2O -1.219 YbS2 2.154 YbS, LiS4 Yb, Li2S -0.947 YbO2 2.853 YbO, Li2O Yb, Li2O -1.155 P3N5 2.286 LiPN2, P Li3P, Li3N -0.552 LiPN2 0.610 Li7PN4, Li3P Li3P, Li3N -0.307 Li7PN4 0.005 Li3P, Li3N Li3P, Li3N -0.005 PF5 3.302 P, LiPF6 Li3P, LiF -2.118 LiPF6 2.735 P, LiF Li3P, LiF -2.059 137 P2S5 2.273 Li3PS4, P4S9 Li3P, Li2S -1.532 Li3PS4 1.719 P, Li2S Li3P, Li2S -1.423 P2O5 3.274 LiPO3, P Li3P, Li2O -1.215 Li4P2O7 2.307 Li3PO4, P Li3P, Li2O -0.826 Li3PO4 0.688 Li2O, Li3P Li3P, Li2O -0.688 LiPO3 2.479 P, Li4P2O7 Li3P, Li2O -0.979 Li6MoN4 -0.362 Li3N, Mo Li3N, Mo 0.362 MoF6 4.335 Li2MoF6 LiF, Mo -3.257 MoO3 3.350 Mo8O23, Li2MoO4 Mo, Li2O -1.757 Li2MoO4 1.842 Li2MoO3, Li4MoO5 Li2O, Mo -1.420 Li4MoO5 1.333 Mo, Li2O Mo, Li2O -1.333 Li7VN4 -0.718 Li3N, V Li3N, V 0.718 LiVN2 0.841 VN, Li7VN4 Li3N, V 0.083 VF5 5.104 VF4, LiVF6 V, LiF -2.968 LiVF6 4.570 Li2VF6 V, LiF -2.849 Li3VS4 1.795 V3S4, Li2S V, Li2S -1.433 V2O5 3.675 LiVO3, V3O7 V, Li2O -1.493 Li3VO4 1.126 Li2O, LiVO2 Li2O, V -0.947 LiVO3 3.229 LiV2O5, Li3VO4 V, Li2O -1.287 Li2MnN2 0.552 Li7MnN4, Mn2N Li3N, Mn -0.092 MnF4 6.030 MnF3, Li2MnF6 LiF, Mn -3.353 Li2MnF6 4.557 LiF, Li2MnF5 LiF, Mn -3.065 138 MnS2 1.830 Li(MnS2)2 Mn, Li2S -1.598 MnO2 3.711 Mn2O3, Li5Mn7O16 Li2O, Mn -1.748 Li2MnO3 1.732 Li2O, LiMnO2 Mn, Li2O -1.381 Li5Mn7O16, Li4Mn5O12 3.417 Mn, Li2O -1.565 Li2MnO3 CoF4 6.635 LiCoF4 LiF, Co -4.150 CoS2 1.984 Co3S4, Li2S Co, Li2S -1.750 CoO2 4.096 Li(CoO2)2 Co, Li2O -2.253 Li8CoO6 2.605 Li5CoO4, Li2O Co, Li2O -1.943 3. Comparison between DFT calculated and experimental values To evaluate the error of the DFT calculation results, we compared the formation enthalpy of binary nitrides and the potential (referenced to Li/Li+) of lithium conversion reaction of binary nitrides based on the following reaction MxNy + 3y Li → y Li3N+ x M. These results are summarized and compared in Table S5. We found that the calculated conversion reaction potential is general in agreement with the experimental value with an error of less than 0.1 V. 139 Table B5. Experimental and calculated formation enthalpy and conversion reaction potential of binary nitrides. ΔHexp. * ΔHcal. ϕexp. ϕcal. (eV/atom) (eV/atom) (V) (V) Li3N -0.428 -0.460 - - TaN -1.309 -1.366 -0.303 -0.297 Ca3N2 -0.912 -0.946 -0.190 -0.175 GaN -0.569 -0.666 0.191 0.169 VN -1.131 -1.164 -0.184 -0.162 NbN -1.227 -1.252 -0.248 -0.221 Be3N2 -1.224 -1.233 -0.450 -0.414 HfN -1.939 -1.940 -0.722 -0.680 Cr2N -0.434 -0.523 0.136 0.091 CeN -1.694 -1.684 -0.559 -0.509 Ta2N -0.942 -0.934 -0.372 -0.321 Zn3N2 -0.047 -0.033 0.531 0.586 TiN -1.754 -1.906 -0.600 -0.657 UN -1.529 -1.501 -0.449 -0.387 BN -1.309 -1.472 -0.303 -0.368 LaN -1.553 -1.517 -0.465 -0.398 ZrN -1.911 -1.870 -0.704 -0.633 Si3N4 -1.104 -1.312 -0.074 -0.152 AlN -1.652 -1.595 -0.531 -0.450 140 Mg3N2 -0.958 -0.909 -0.228 -0.144 CrN -0.608 -0.527 0.165 0.262 Th3N4 -1.935 -1.829 -0.559 -0.453 PuN -1.553 -1.438 -0.465 -0.345 EuN -1.129 -0.996 -0.183 -0.051 YN -1.553 -1.879 -0.465 -0.639 ThN -1.965 -1.761 -0.740 -0.561 ScN -1.628 -2.106 -0.515 -0.790 InN -0.717 -0.105 0.092 0.543 * Chase, M. W. NIST-JANAF thermochemical tables 4th ed.; National Institute of Standards and Technology: Gaithersburg, 1998. 141 4. Li-M-O-N grand potential phase diagrams Figure B1. Grand potential phase diagrams in equilibrium with Li metal. The composition ranges that form passivating interphase and that form non-passivating interphase are colored green and red, respectively. (a) O (b) O Li2O Li2O Li3N LiB Li3N Zr N Li Zr3BN2 B N Li2ZrN2 ZrN 3N2 Zr (c) O (d) O Li2O Li2O Ti6O Li16(TaN4)2O Li3N Ta Li3N Ti N Li7TaN4 Li4TaN3 Ta N Li5TiN3 TiN Ti2N Ti (e) O (f) O Li2O Li2O Li3N Li21Si5 Li3N V N Li5SiN3 Si N Li7VN4 V2N V (g) O (h) O Li2O Li2O Li3N W Li3N Li5Mg N Li6WN4 W N LiMgN Mg 142 (i) O (j) O Li2O Li2O Li3N LiCl Li3N LiBr N Li4NCl Cl N Li10BrN3 Li5Br2N Br (k) O (l) O Li2O Li2O Li3N LiI Li3N Sn N Li7IN2 I N Sn (m) O (n) O Li2O Li2O Li3N Li2Ga Li3N LiZn N Ga N Zn 143 Appendix C 1. Lattice parameters based on the classic interatomic potential Figure C1. 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