University of Southern California Dissertations and Theses

Understanding the role of electrode design in determining the electrochemical performance of high-energy/high-power lithium-ion and lithium-sulfur batteries

(USC Thesis Other)

Understanding the role of electrode design in determining the electrochemical performance of high-energy/high-power lithium-ion and lithium-sulfur batteries

PDF

Download a page range

Download transcript

Copy asset link

Request this asset

Transcript (if available)

Content Copyright 2022 Rodrigo Elizalde Segovia Understanding the Role of Electrode Design in Determining the Electrochemical Performance of High-Energy/High-Power Lithium-Ion and Lithium-Sulfur Batteries by Rodrigo Elizalde Segovia A Dissertation Presented to the FACULTY OF THE USC GRADUATE SCHOOL UNIVERSITY OF SOUTHERN CALIFORNIA in Partial Fulfilment of the Requirements for the Degree DOCTOR OF PHILOSOPHY (CHEMISTRY) December 2022 ii The dependence in modern society and the energy stored in fossil fuel is not sustainable. We have to learn to harness the energy that comes to us from the sun either in the form of wind or in the form of radiant energy. And we have to be able to convert it into electric power which we know how to do. But as diverse… you can transport electric power over wires over some distances but you have to have a collection site. But you have to be able to store that energy because it comes in at time scales that are very different from the time scales of demand. So that is one of the reasons you work on batteries because they store electric power. John B. Goodenough —Nobel Prize in Chemistry, 2019— iii Dedication To my family. Bertha, Antonio and Diego. All my achievements are because of you. I am standing here thankful, peaceful, and humble because of your infinite love and support that guided me through this long and winding journey. iv Acknowledgments • My deepest thanks to my advisor Professor Sri R. Narayan. I am a really grateful for all your mentorship, support, and guidance during my Ph.D. Your experience and knowledge allowed me to think always creatively and going beyond to what is expected. You believed in me. I will always be a proud member of your lab. Thanks. • Professor G. K. Surya Prakash, thanks for all your mentorship during my research, and for opening the opportunity to do my Ph.D at USC. I am proud, grateful of being part of the Loker Hydrocarbon Research Institute founded by Professor George Olah and you. • Thank you, Professor Barry C. Thompson for all the great and insightful research collaboration towards new chemistries and materials. Your support and guidance as a committee member of my PhD dissertation has been fundamental and extremely valuable. • I cordially thank Dr. Jayakanth Ravichandran, Dr. Kathrine Shing and Dr. Megan Fieser for providing me useful suggestions towards successful completion of my doctoral dissertation. • My sincere gratitude to Professor Francisco Valero-Cuevas, you gave us the opportunity to reach this dream, I will never forget what you did for all the young Mexican students. • Thank you, Dr. Ahamed Irshad for your patience, training, and great collaboration that guided me during my PhD studies. • To Dr. Pratyusha Das for the fundamental and successful research collaboration in the last years. Thank you. • To all Narayan’s lab past and present members: Dr. Derek Moy, Dr. Advaith Murali, Dr. Billal Zayat, Sara Bollstetter, Germany Diaz De la Cruz, Ryuichi Fukuda, Dr. Sundar Rajan, Dr. Debanjan Mitra, Dr. Bo Yang, Dr. Buddhinie Jayathilake, and Dr. Dan Fang. • Special thanks to the Synthetic Control Across Length-scales for Advancing Rechargeables (SCALAR) a US Department of Energy, Energy Frontier Research Center (EFRC). • Thanks to Dr. Robert Aniszfeld, Magnolia Benitez, Anahit Martirosyan, Michele Dea, David Hunter, Jessie May, and to the chemistry department at USC. • To Dr. Octavio Marin Pardo, we fought together during this endless PhD journey. We survived the pandemic, we succeed together, and will continue having this great friendship. • Special thanks to all the Mexican community at USC. José Anaya, Dr. Víctor Martínez, Santiago Zapata, Arturo Cajal, Sofia Plata, Xanath Ispizua, Darío Urbina, Dr. Enrique Arguelles, Dr. Leonardo Nava, and Ivan Trujillo. • My friends at USC. Dr. Jinghan He, Sungah Kim, Justin Overhulse, Cindy Tseng, Austin Mencke, and Dr. Mahta Barekatain. • My dear mom, dad and brother, thanks for enriching my life with your infinite love. Ineffable to describe how grateful I am, I feel fulfilled. • Carla Sofia Lara Peralta, you are my inspiration, my motivation, that luminary in my life. v Table of Contents Epigraph……………………………………………………………………………………………………………………………………ii Dedication…………………………………………………………………………………………………………………………………iii Acknowledgements ......................................................................................................................... iv List of Tables .................................................................................................................................... x List of Figures and Schemes ............................................................................................................ xi Abstract .......................................................................................................................................... xii Preface .......................................................................................................................................... xiii Chapter 1- Introduction and background ....................................................................................... 1 Lithium-ion batteries .................................................................................................................. 1 Positive electrodes for lithium-ion batteries .......................................................................... 4 Negative electrodes for lithium-ion batteries ........................................................................ 5 Electrolytes ............................................................................................................................. 6 Problems and Significance .......................................................................................................... 7 Grasping the theoretical limits on lithium-ion batteries ........................................................ 7 Lithium-ion cell evaluation ..................................................................................................... 9 Formation and Equilibration ................................................................................................. 10 Optimization of a lithium-ion cell ......................................................................................... 11 Overcoming the trade-off between energy and power density........................................... 12 Lithium-Sulfur batteries ............................................................................................................ 13 Li-S batteries charge/discharge mechanism ......................................................................... 15 The polysulfide shuttle .......................................................................................................... 18 Challenges and opportunities in Li-S batteries ..................................................................... 19 Solid-State electrolytes ......................................................................................................... 20 Hypotheses tested in the thesis. ............................................................................................... 22 Lithium-ion batteries ............................................................................................................ 22 Approach ............................................................................................................................... 22 Lithium-sulfur batteries: ....................................................................................................... 23 Objectives and Hypotheses................................................................................................... 23 Innovation ................................................................................................................................. 23 Importance of Electrochemical Techniques ............................................................................. 24 References ................................................................................................................................ 24 Chapter 2 – Experimental methods .............................................................................................. 29 vi Abstract ..................................................................................................................................... 29 Methods for characterizing PProDOT-Hx 2 as cathode binder (Chapter 3) ............................... 29 Polymer Synthesis ................................................................................................................. 29 GIWAXS and Ellipsometry of PProDOT-Hx 2 .......................................................................... 31 Performance as an NCA Cathode Binder .............................................................................. 33 Ionic and Electronic Conductivity of PProDOT-Hx 2 ............................................................... 34 Fabrication and Electrochemical Characterization of PProDOT-Hx 2 Thin Films ................... 36 Electron paramagnetic resonance ........................................................................................ 37 Methods for High-Energy/High-Rate π-Conjugated Polymers as Binders for Lithium Metal Batteries (Chapter 4) ................................................................................................................. 39 Monomer and Polymer Synthesis ......................................................................................... 39 Polymer physical Characterization ....................................................................................... 41 Conductivity measurements ................................................................................................. 44 Cathode Preparation ............................................................................................................. 45 Lithium Anode Preparation ................................................................................................... 46 Electrolyte ............................................................................................................................. 46 Electrode Characterization ................................................................................................... 46 Cell assembly ......................................................................................................................... 47 Electrochemical Testing and Analysis ................................................................................... 47 Methods for Enhancing the Ionic Conductivity of Poly(3,4-propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries (Chapter 5) ................................................................................................................................ 49 Monomer and Polymer Synthesis ......................................................................................... 49 Molecular Characterization .................................................................................................. 51 Monomer NMR ..................................................................................................................... 53 Polymer NMR ........................................................................................................................ 55 Additional NMR (Polymers soluble in battery electrolyte) ................................................... 57 Interdigitated Electrode Preparation .................................................................................... 59 Electrochemical Testing ........................................................................................................ 60 Conductivity Measurement .................................................................................................. 60 Electrode preparation ........................................................................................................... 60 Cell assembly ......................................................................................................................... 61 Cell Testing ............................................................................................................................ 62 Electron Paramagnetic Resonance (EPR) .............................................................................. 62 GIWAXS ................................................................................................................................. 65 Swelling Studies .................................................................................................................... 65 GPC Traces of the Polymers: ................................................................................................. 67 Methods for Li xNaNb 13O 33 as an anode material for lithium-ion Batteries (Chapter 6) .......... 70 Electrode preparation and cell assembly ............................................................................. 70 Electrochemical testing ......................................................................................................... 71 Methods for Ketjen black-based Sulfur Cathodes in Li-S (Chapter 7) ...................................... 71 Electrode preparation, cell assembly and testing ................................................................ 71 vii Methods for Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature (Chapter 8) ................................................ 72 Preparation of the Composite sulfur Electrode. ................................................................... 72 Preparation of the bilayer solid electrolyte .......................................................................... 73 Coin cell fabrication .............................................................................................................. 74 Electrochemical characterization ......................................................................................... 75 Scanning electron microscopy (SEM).................................................................................... 75 Methods for The Role of Functionalized Conducting Polymer Binders in Addressing the Technical Challenges of Lithium-Sulfur Batteries (Chapter 9) .................................................. 75 Experimental ......................................................................................................................... 75 Polymer Synthesis ................................................................................................................. 76 Electrochemical Characterization of Polymers ..................................................................... 80 Physical Characterization ...................................................................................................... 81 Full Cell Assembly, Shuttle Current Measurement, Impedance Spectroscopy, and Electrical Performance Testing ............................................................................................................. 83 Monomer and Polymer NMR ................................................................................................ 85 References ................................................................................................................................ 87 Chapter 3 - High-Rate Lithium-Ion Cell with Dihexyl-Substituted Poly (3, 4- Propylenedioxythiophene) Mixed Conductive Polymer As Electrode Binder 1 ............................. 89 Abstract ..................................................................................................................................... 89 Introduction .............................................................................................................................. 90 Electrochemical Properties of PProDOT-Hx 2 ............................................................................ 93 Electron Paramagnetic Resonance (EPR) .............................................................................. 98 Electronic and Ionic Conductivity of PProDOT-Hx 2 ............................................................. 100 Electronic Conductivity as a Function of Degree of Doping ............................................... 102 Ionic conductivity ................................................................................................................ 104 Morphology Changes of PProDOT-Hx 2 with Electrochemical Doping ................................ 105 Performance as a Cathode Binder .......................................................................................... 110 NCA electrodes without binder .......................................................................................... 112 NCA composite electrodes with a higher loading of 11 mg/cm 2 ........................................ 113 Porosity and impedance analysis of the NCA electrodes. .................................................. 116 Conclusions ............................................................................................................................. 120 References .............................................................................................................................. 122 Chapter 4 - High-Energy/High-Rate π-Conjugated Polymers as Binders for Lithium Metal Batteries 1 ..................................................................................................................................... 129 Abstract ................................................................................................................................... 129 Introduction ............................................................................................................................ 129 Cathodes with new π-conjugated polymers for a high-energy lithium battery ..................... 132 Understanding the impact of CPs on the cell’s internal processes. ....................................... 140 Energy-dispersive X-ray spectroscopy (EDS) mapping. ...................................................... 158 Determination of the effective Porosity (Imbibition/Archimedes’ Method) ..................... 163 viii Phase Equilibria and Polarization via Differential capacity analysis ....................................... 166 Rate Capability under mass-efficient and non-limited conditions ......................................... 168 Conclusions ............................................................................................................................. 171 References .............................................................................................................................. 172 Chapter 5 - Enhancing the Ionic Conductivity of Poly(3,4-propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries 1 ....... 177 Abstract ................................................................................................................................... 177 Introduction ............................................................................................................................ 178 Synthesis of (Hex:OE) PProDOT Random Copolymers using Direct Arylation Polymerization (DArP) ...................................................................................................................................... 183 Electrochemical Properties of (Hex:OE) PProDOT Random Copolymers ............................... 184 Electron Paramagnetic Resonance (EPR). ........................................................................... 188 Morphology Changes of (Hex:OE) PProDOTs with Electrochemical Doping ...................... 195 Solvent Swelling of (Hex:OE) PProDOTs Using Battery Electrolytes ................................... 201 Electrochemical cycling of NCA cathodes incorporating (Hex:OE) PProDOTs as binders. ..... 202 Impact of the CP binders on Rate Capability ...................................................................... 204 Warburg coefficient analysis .................................................................................................. 204 Effect of low binder content and carbon content. ............................................................. 215 Conclusions ............................................................................................................................. 220 References .............................................................................................................................. 222 Chapter 6 - NaNb 13O 33 as an anode material for lithium-ion Batteries...................................... 229 Abstract ................................................................................................................................... 229 Incremental capacity analysis ................................................................................................. 229 Rate Capability and incremental capacity analysis. ................................................................ 232 Long-term cycling and EIS characterization. ........................................................................... 236 Full-Cell configuration: NaNb 13O 33||NCA and NaNb 13O 33||NCA-PProDOT-Hx 2 .................... 248 Rate Capability of NCA-PProDOT-Hx 2||NaNb 13O 33 (Full-cell) ............................................ 250 Pouch cell design. ................................................................................................................ 252 Chapter 7 - Understanding the Polarization Behavior of Ketjen black-based Lithium-Sulfur Battery Cathodes ........................................................................................................................ 257 Abstract ................................................................................................................................... 257 Introduction ............................................................................................................................ 258 Results and Discussion ............................................................................................................ 261 Low-rate performance of sulfur electrodes with KB-300 and KB-600 ............................... 261 Discharge rate capability test ............................................................................................. 263 EIS studies on sulfur electrode prepared using KB-600 ...................................................... 265 Physical significance of resistance elements (R0, R1 and R2) ............................................ 267 Variation in R0, R1 and R2 during 1 st discharge .................................................................. 269 Effect of melt infusion of sulfur on the internal resistance ................................................ 275 Effect of mixing with Super-P  ........................................................................................... 278 ix Rate capability and Long-term cycling test ......................................................................... 282 Conclusion ............................................................................................................................... 283 References .............................................................................................................................. 284 Chapter 8 - Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature .............................................................................. 287 Abstract ................................................................................................................................... 287 The concept of a bilayer solid electrolyte and the composite intercalating sulfur electrode 289 Results and Discussion ............................................................................................................ 293 Bilayer electrolyte ............................................................................................................... 293 Discharge profile and impedance response of CEBE Li-S cell ............................................. 294 Effect of Galvanostatic cycling. ........................................................................................... 298 Beneficial Effects of Alumina Addition to the Polymer Electrolyte Layer. ......................... 300 Conclusions ............................................................................................................................. 302 A mass-efficient CEBE Li-S cell (TAG proposal) ....................................................................... 304 Abstract ............................................................................................................................... 304 Cell optimization ................................................................................................................. 304 Shuttle current measurement for the CEBE Li-S cell .......................................................... 305 EIS characterization of the mass-efficient CEBE Li-S cell .................................................... 306 Galvanostatic cycling of the mass-efficient CEBE Li-S cell. ................................................. 308 References .............................................................................................................................. 311 Chapter 9 – The Role of Functionalized Conducting Polymer Binders in Addressing the Technical Challenges of Lithium-Sulfur Batteries ....................................................................................... 315 Abstract ................................................................................................................................... 315 Introduction ............................................................................................................................ 315 Results and Discussion ............................................................................................................ 319 Electrochemical Cycling of Polymer Thin films ................................................................... 319 Electronic and Ionic Conductivity of Polymer Thin films. ................................................... 321 Structural Characterization ................................................................................................. 324 Interaction of the polymers with polysulfides. ................................................................... 327 Polysulfide Shuttle Current Measurement ......................................................................... 329 Electrochemical Impedance Spectroscopy ......................................................................... 332 Full Cell Electrochemical Cycling ......................................................................................... 334 Conclusions ............................................................................................................................. 339 References .............................................................................................................................. 340 References .................................................................................................................................. 343 Appendix - Full list of Figures ...................................................................................................... 344 x List of Tables Table 4.1 Cell parameters/performance for the high energy NCA-lithium metal cells with π-conjugated polymers. Table 4.2. Fitting parameters and their corresponding relative error for the 1 st cycle of the Li- NCA- π-conjugated polymers cells at approximately 4.0 V vs Li + /Li. Table 4.3 Fitting parameters and their corresponding relative error for the 26 th cycle of the Li- NCA- π-conjugated polymers cells at approximately 4.0 V vs Li + /Li. Table 4.4 Effective Porosity by Imbibition/Archimedes’ Method for the NCA- π-conjugated polymers electrodes. Porosity was determined using the average value of triplicated measurements using isopropyl alcohol as the fluid. Table 6.1 Fitting parameters and their corresponding relative error for the 1 st cycle of the Li xNaNb 13O 33||Li cell as a function of potential vs Li + /Li. Table 6.2 Fitting parameters and their corresponding relative error for the 600 th cycle of the Li xNaNb 13O 33||Li cell as a function of potential vs Li + /Li. Table 6.3 Electrodes and electrolyte composition utilized in the NCA|| NaNb 13O 33 Full-cell. Table 7.1. Representative impedance fit parameters for 20 % and 72 % DOD Table 8.1. Energy density at the cell-level for the CEBE Li-S cell. Table 8.2 Cell parameters and energy performance comparison of the CEBE Li-S cell baseline from 2020 against the mass-efficient 2022 TAG cell. xi List of Figures and Schemes Consult Appendix for the full list of Figures and Schemes xii Abstract In the last couple of decades lithium-ion batteries have fulfilled most of the energy storage demands for mobile applications. However, there is a continuous increase in energy and power demands that current lithium-ion batteries will not be able to meet due to their theoretical limits for energy storage. Therefore, there is a necessity for understanding the pathways and designs to increase the capabilities of practical high-energy and high-power lithium batteries. This dissertation presents a systematic, comprehensive, and rationalized electrochemical study and electrode design for building next generation lithium-ion and lithium-sulfur batteries. We introduce the use new conducting polymers as cathode binders/additives, new niobium-oxide based intercalating materials with fast lithium-ion diffusivity and structural stability, and a novel composite solid-state electrolyte for lithium-sulfur batteries. Electrochemical and physicochemical characterization techniques such as Electrochemical Impedance Spectroscopy (EIS), Incremental Capacity Analysis (ICA), Cycling Voltammetry (CV), Electron Paramagnetic Resonance (EPR), and spectroscopy techniques have been utilized to achieve an in-depth understanding. We have found that using various conducting polymers as cathode binders significantly enhance the energy/power capability (320 Wh kg -1 at C/2), and cycle-life of a lithium- ion battery due to their high mixed conductivity and protection against the rapid growth of the solid electrolyte interphase. We found that NaNb 13O 33 as an anode material has an impressive rate capability (80 mAh g -1 at 20C) and a remarkable capacity retention (80% after 600 cycles). Finally, we introduce a unique solid-state lithium sulfur cell based on a flexible bilayer made of a lithium-ion intercalation material and a polymer electrolyte that operates at room temperature to yield as much as 85% of the theoretical capacity. xiii Preface Experimental Acknowledgment The experimental work and analysis presented in this thesis is a conjoined effort form multiple authors in academic institutions across southern California. The following list, acknowledges the name of the experimentalist and their corresponding role: All conducting polymers organic synthesis and physical characterization by: Pratyusha Das (University of Southern California, USC) Measurement of polymer conductivity by: Billal Zayat (University of Southern California, USC) Grazing incidence wide-angle x-ray scattering (GIWAXS) by: Charlene Z. Salamat (University of California, Los Angeles, UCLA) Electron Paramagnetic Resonance (EPR) by: Gordon T. Pace, Dakota Rawlings, Dongwook Lee, Rebecca C. Vincent (University of California at Santa Barbara, UCSB) Synthesis of NaNb 13O 33 lithium intercalating active material by: Ashlea Patterson, Kira Wyckoff, Arava Zohar (University of California at Santa Barbara, UCSB) Synthesis scale-up of PProDOT-Hx 2 conducting polymer by: Liwei Ye and Alexander Schmitt (University of Southern California, USC) Characterization of PProDOT-Hx 2 films: Qiulong Wei (University of California, Los Angeles, UCLA) Sulfur-Ketjen Black polarization experiments by: Ahamed Irshad (University of Southern California, USC) xiv Most of the data and analysis presented in this thesis has been published in indexed peer reviewed chemistry and electrochemistry journals. The following bibliography lists the publications: • Elizalde-Segovia, R.; Das, P.; Zayat, B.; Irshad, A.; Thompson, B. C.; Narayanan, S. R. Understanding the Role of π-Conjugated Polymers as Binders in Enabling Designs for High- Energy/High-Rate Lithium Metal Batteries. J. Electrochem. Soc. 2021, 168 (11), 110541. https://doi.org/10.1149/1945-7111/ac3850. • Das, P.; Elizalde-Segovia, R.; Zayat, B.; Salamat, C. Z.; Pace, G.; Zhai, K.; Vincent, R. C.; Dunn, B. S.; Segalman, R. A.; Tolbert, S. H.; Narayan, S. R.; Thompson, B. C. Enhancing the Ionic Conductivity of Poly(3,4-Propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries. Chem. Mater. 2022, 34 (6), 2672–2686. https://doi.org/10.1021/acs.chemmater.1c03971. • Elizalde-Segovia, R.; Irshad, A.; Zayat, B.; Narayanan, S. R. Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-Layer Solid Electrolyte Operable at Room Temperature. J. Electrochem. Soc. 2020, 167 (14), 140529. https://doi.org/10.1149/1945- 7111/abc4c0. • Das, P.; Zayat, B.; Wei, Q.; Salamat, C. Z.; Magdău, I.-B.; Elizalde-Segovia, R.; Rawlings, D.; Lee, D.; Pace, G.; Irshad, A.; Ye, L.; Schmitt, A.; Segalman, R. A.; Miller, T. F.; Tolbert, S. H.; Dunn, B. S.; Narayan, S. R.; Thompson, B. C. Dihexyl-Substituted Poly(3,4- Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. 2020, 32 (21), 9176–9189. https://doi.org/10.1021/acs.chemmater.0c02601. • Narayan, S. R.; Elizalde-Segovia, R.; Jayathilake, B.; Irshad Maniryanganam, A. Long-Life Lithium-Sulfur Battery Using a Novel Flexible Bi-Layer Solid State Electrolyte. WO/2020/243066, December 3, 2020. • Zayat, B.; Elizalde-Segovia, R.; Das, P.; Salamat, C. Z.; Irshad, A.; Tolbert, S. H.; Thompson, B. C.; Narayanan, S. R. The Role of Functionalized Conducting Polymer Binders in Improving Power Density and Cycle Life of Lithium-Sulfur Batteries. Journal of The Electrochemical Society 2022. • Patterson, A. R.; Elizalde-Segovia, R.; Wyckoff, K.; Narayan, S. R.; Seshadri, R.; Griffith, K. Host Structure Analysis and Lithium Electrochemistry of NaNb13O33. Chemistry of Materials 2022. 1 Chapter 1- Introduction and background Lithium-ion batteries The global demand for energy has increased exponentially in the last couple of centuries due to the growth of the global population, use of land, and technological development. 1–4 To satisfy the energy demands different energy sources have been created such as the combustion engine, nuclear plants, hydroelectric plants, solar panels, wind turbines, etc. However, the time scales of energy production and energy demand fluctuate leading to the crucial need for energy storage. 5– 9 Currently, energy-storage is required in portable devices, electric cars, back-up power for homes and businesses, large-scale community and city level power management, and grid-connected power sources. 10 There are various methods for storing energy such as mechanical devices, fossil fuels, thermal devices, electric capacitors, chemical fuels, and electrochemical systems or batteries. Batteries are devices that can convert chemical energy into electrical energy. Two of the main advantages of utilizing batteries is that they have a minimal contribution to the carbon emissions and are not limited by the Carnot cycle efficiency. The pioneering scientific works that lead to invention of batteries are credited to Alessandro Volta and Luigi Galvani in the eighteenth century. 11 Volta’s invention ignited the research towards new batteries and the progress of the electrochemical science. The isolation of elemental lithium was carried by W.T Brande and Sir Humphrey Davy in 1821 through the electrolysis of lithium oxides. 12,13 In 1913 the electrochemical properties of lithium were firstly studied by Lewis 14 , and it was rapidly recognized that lithium could serve as an anode (negative electrode) due to its excellent properties such as low reduction potential (- 2 3.04 V vs SHE), a high specific charge capacity (3860 mAh g -1 ), and low gravimetric density (0.534 g cm -3 ). The first primary lithium battery was reported by Harris in 1958 utilizing non-aqueous electrolytes. 15 During the 1970s lithium metal batteries were commercialized utilizing lithium insertion compounds such as TiS 2 and MoS 2 by Exxon and Moli Energy, respectively. Nevertheless, this early lithium metal batteries had low voltages and several safety concerns due to the reactivity of the lithium electrode. In 1980, Goodenough et al. introduced high capacity metal oxides such as LiCoO 2 that could intercalate an store lithium ions . 16–18 However, it was not until 1985 that Akira Yoshino from the Asahi Kasei Co in Japan conceived and patented the Lithium-ion battery by replacing the lithium metal electrode with a carbonaceous material and pairing it with Goodenough’s LiCoO 2. 18,19 In 1991 Sony Co. was the first in commercialize the secondary lithium-ion battery. 20,21 Since the 1990s lithium-ion batteries have fulfilled most of the energy needs in portable devices and mobile electrical applications due to its high-energy density, good performance, relative thermal stability, and absence of memory effect. 17,22 The development of lithium-ion batteries is one of the greatest human triumphs in science and technology, being the subject of the 2019 Nobel Prize in Chemistry awarded to John B. Goodenough, Akira Yoshino, and M. Stanley Whittingham in recognition of their pioneering work and fundamental research in lithium-ion batteries. 23–26 It is important to notice that lithium-ion batteries do not have elemental or metallic lithium, there is only migration, insertion, and withdrawal of Li + ions. The process of insertion and removal of lithium ions into an electrode material or host is defined as intercalation and deintercalation respectively. The electrochemical properties and mechanisms of a lithium-ion battery depend on the electrodes and electrolytes that are utilized. The electrochemical reactions occurring at each 3 electrode and the total cell reaction for a conventional LiCoO 2 ||graphite lithium-ion battery are the following: Positive electrode: 𝐿𝑖𝐶𝑜 𝑂 2 ⇌ 𝐿𝑖 1−𝑥 𝐶𝑜 𝑂 2 + 𝑥𝐿𝑖 + 𝑥 𝑒 − ( 1.1) Negative electrode: 6𝐶 + 𝑥 𝐿𝑖 + + 𝑥 𝑒 − ⇌ 𝐿𝑖 𝑥 𝐶 6 ( 1.2) Total cell reaction: 6𝐶 + 𝐿𝑖𝐶𝑜 𝑂 2 ⇌ 𝐿𝑖 𝑥 𝐶 6 + 𝐿𝑖 1−𝑥 𝐶𝑜 𝑂 2 ( 1.3) The positive electrode is the electrode that has the higher reduction potential when compared with the negative electrode. Similarly, the negative electrode is defined as the electrode with the lower reduction potential when compared with the positive electrode. During discharge the positive electrode is referred as a cathode and the negative electrode is referred as the anode because a reduction and oxidation reaction occur respectively at each electrode. However, during charge the positive electrode undergoes an oxidation and the negative electrode through a reduction changing their designations into anode and cathode, respectively. To avoid confusion, in this text we refer to the cathode as the positive electrode and anode as the negative electrode. Equation 1.1 and 2.2 shows the reversible reaction that LiCoO 2 and Carbon (graphite) undergoes during intercalation/deintercalation. During deintercalation or charging an electron is released and Co 3+ is oxidized back into Co 4+ . For the graphite (negative electrode) a Li + is intercalated into the interlayer and the carbon matrix gains an electron. During intercalation or discharging an electron is gained in the positive electrode and lost from the negative electrode 4 where Co 4+ is reduced back into Co 3+ and graphite returns to its uncharged state releasing a lithium ion. Figure 1.1 is a schematic illustration of a LiCoO 2 ||graphite lithium-ion battery. Figure 1.1. Schematic illustration of a LiCoO 2||graphite lithium-ion battery. Positive electrodes for lithium-ion batteries Positive electrode materials or cathodes typically have a layered or spinel structure that can intercalate lithium ions. The ideal cathode should have an excellent reversibility and energy. Energy is directly related to capacity and voltage, and reversibility is related to chemical stability associated with the structural changes and reactivity with the electrolyte. The power that the cathode can deliver is related to the current density that it can support. Utilizing high current densities could induce polarization losses in the cell. To reduce the losses in capacity and the polarization effects, a cathode with high electronic-ionic conductivity and lithium ion diffusivity is needed. The chemical formula of most of the positive electrode intercalating compounds is the following: Li xM yX z where M is a transition metal ion M n+ and X is typically an oxide or a sulfide. 5 To maximize the energy and reversibility, the cathode material should have a low formula weight for achieving a higher specific energy, M should have a high redox potential, Li should have a high x value, and the changes in x should not significantly affect the structure of the compound. Layered cathodes typically have a two-dimensional Li + path diffusion while spinel have three- dimensional paths. LiCoO 2, LiNiO 2 LiCrO 2, Li 2MoO 3 are example of layered cathode materials while manganese-based compounds are spinel-type such as LiMn 2O 4 and Li2Mn 2O 4. In terms of sustainability either cathode or anodes should be environmentally friendly, non-toxic, abundant and inexpensive. Negative electrodes for lithium-ion batteries There is a variety of negative electrodes or anodes including silicon, silicon oxides (SiO X), carbon graphitic compounds, amorphous carbons, hollow carbons, tin-based compounds, metalloid- metal alloys, and metal oxides such as LiTi 2O 3. Intercalation is the main mechanism of lithium insertion for carbonaceous electrodes and metal oxides while alloying occurs in silicon and tin compounds. For maximizing the energy of the cell, the lithiated anode should have a redox potential closer to the potential of metallic lithium and a high capacity. Furthermore, the operational potential window during charge and discharge should not vary significantly. Nevertheless, one the most critical features of the anodes is the formation of the solid electrolyte interface (SEI) in the electrode surface. The SEI film is a Li + -conducting and electronically- insulating mosaic structured layer composed of the degradation products from the electrolyte and salts such as lithium carbonate, lithium oxide and lithium fluoride. Therefore, the ideal negative electrode should have a suitable surface area and compatibility with the electrolyte that can lead to the formation of a homogenous SEI. 6 Electrolytes Electrolytes can be either liquid, solid-state or gels. Non-aqueous liquid lithium-ion electrolytes are mainly carbonate-based and in some cases ether-based. However, fluorinated-based electrolytes have been utilized for extreme low temperature applications and/or high- voltages. 27,28 The most used carbonate-based solvents in lithium-ion batteries are ethylene carbonate (EC), dimethyl carbonate (DMC), propylene carbonate (PC), diethyl carbonate (DEC), and methyl carbonate (EMC). 29 A combination of two or more solvents is usually necessary to enhance the ionic conductivity and/or stability of the electrolyte. On the other hand, the ideal lithium electrolyte salt should have a high ionic conductivity, a high mobility in the solvent, thermal and electrochemical stability, low cost, non-toxic, and the anion should be stable to oxidation and should passivate the aluminum current collector from anodic dissolution. 30 Some of the most utilized lithium salts are the following: Lithium hexafluorophosphate (LiPF 6) lithium Perchlorate (LiClO 4), lithium Tetrafluoroborate (LiBF 4), lithium Bis(trifluoromethanesulfonyl)imide (LiTFSI). To summarize, liquid electrolytes should have good ion solvation, electrode compatibility, ion conductivity, potential window stability, chemical stability, thermal stability, and phase stability. In terms of safety, the electrolyte should not be flammable, non-toxic, and low volatile. Solid electrolytes could be either inorganic, polymer or composite. The main advantage of utilizing solid-state electrolytes is their safety due to their non-flammability. 31 However, several challenges are present in terms of stability, operating temperature, and ionic conductivity. The ideal solid-sate electrolyte should have a high ionic conductivity but low electronic conductivity 7 to prevent any short-circuit in the cell. In addition, thermal, mechanical, chemical and electrochemical stability are required. Problems and Significance Grasping the theoretical limits on lithium-ion batteries Lithium-ion technology has reached technical and commercial maturity. 32 In the last decades, scientists and engineers have made significant efforts for optimizing, characterizing, and understanding the variables that affect the performance of lithium-ion batteries. 32,33 The gap between the theoretical capacity and the observed reversible capacity is getting narrow, and new paths for building next generation lithium-ion batteries are required. 34 These “beyond” lithium- ion batteries improvements can be done at three different but interrelated levels; at the active material level, where the development consists in finding and synthesizing new intercalating compounds with a high mixed conductivity, lithium-ion diffusivity, theoretical capacity, high- voltage, and structural stability. Furthermore, the new materials should be relatively inexpensive, abundant, and non-toxic. 35 At the electrode level, the optimization consists in increasing electrode porosity, decreasing its tortuosity, enhancing its electroactive surface area, and optimizing its formulation. In contrast with many other secondary batteries, the lithium-ion electrodes are not solely composed of the active material. They are composed of a mixture of active material, conductive carbon, and a polymeric binder leading to a porous electrode. A vast variety of carbon materials and binders have been explored and tested. Material selection and finding the right electrode formulation constitutes one the greatest challenges in battery research. Moreover, the choice of electrode fabrication methods such as slurry mixing, coating, drying, and calendaring have a tremendous impact in the actual cell’s performance. At the cell- 8 level, the negative to positive capacity ratio (N/P) and the electrolyte to capacity ratio (E/C) are the crucial parameters to be considered due their impact in the cell’s energy, power, and cycle- life. 36 The N/P value the ratio of the amount of negative material to positive material. Ideally, a N/P value close to 1 is desired to maximize the specific energy of the cell. However, the N/P ratio should not be lower than 1 because it could result in lithium plating at the anode’s surface. The E/C is a measure of the amount of electrolyte that is utilized in the cell. Nevertheless, it is important to consider that certain amount of electrolyte will be consumed during the SEI formation so an excess of electrolyte is needed to fill or “wet” all the electrode pores that will allow a proper Li + migration/diffusion. At the cell-level, the optimization consists of lowering the mass of all the inactive components and maximizing the amount of active material in the cell. A high-energy cell can be achieved by building thicker electrodes with higher active material mass fraction, lower amounts of electrolyte, and by utilizing a minimum amount of the negative electrode (N/P ≈ 1). 36,37 However, lowering the electrolyte amount and reducing the carbon content in the electrode could result in a negative outcome that could affect the cell’s cycle life and rate capability due to the electrolyte consumption during cycling and the lack of ionic transport between the particles. Figure 1.2 shows a schematic illustration that summarizes the parameters affecting the lithium-ion performance at each level. 9 Figure 1.2. Schematic illustration of the strategies that can be pursued to increases the energy and power density at the active material level, electrode level, and cell level. Lithium-ion cell evaluation Battery cell design requires the evaluation of the following fundamental parameters: energy density, reversibility, and power output. These properties are tightly interrelated and improving one of them will have a direct positive impact over the others. It is important to mention that the mass loading of the active material is crucial design parameter because all the other cell properties depend on it. At the cell level, the utilization/energy density level is measured by its specific capacity. The specific capacity is commonly obtained from the galvanostatic plots or charge-discharge curves (GCD). The utilization of materials is given by the gravimetric and volumetric energy density. The utilization is highly dependent on the active material specific capacity and amount that is accessible for redox reaction at a certain current density. Cells with a high polarization resistance have a lower utilization at higher currents. Polarization resistance 10 depends mainly on the overpotential and the reaction rate of each electrode. Rate capability is mainly limited by the electronic-ionic conductivity and lithium-ion diffusivity in the electrodes. The reversibility of the cell is typically evaluated by some parameters as: cycle life, coulombic efficiency, and round-trip efficiency. Cycle life is simply the number of cycles that the battery can perform without losing significant amount of capacity. The coulombic efficiency of the cell is the ratio of the capacity realized during discharge to the charge uptake during charging. Higher coulombic efficiency will lead to a lower capacity fade after each cycling. Coulombic efficiency depends on the accessible active material for charge/discharge and its resistance. The round-trip efficiency, denoted as the ratio of the energy released during discharge and the energy expended in the recharge process. Round-trip efficiency also depends on the overpotential for both discharge and charge reactions and can be calculated from the galvanostatic plots by dividing the discharge voltage by the charge voltage at each value of cell voltage. Formation and Equilibration There are other parameters that have an effect on the response of the cell such as SEI formation and the equilibration time. Equilibration occurs after assembling the battery. A certain time is required for the electrolyte to penetrate all the electrode pores, and this time is designated as the “equilibration time”. Equilibration is highly dependent on the porosity and tortuosity of the electrode, the electrode thickness, the amount of electrolyte, the electrolyte viscosity, and the cell size. Cycling a cell that is not equilibrated will yield a lower utilization. SEI formation or simply the process of formation is a crucial step in the battery assembly. This step ensures a homogenous and well distributed SEI layer on the electrodes surface. Formation is carried after equilibration and it involves cycling the cell at low rates (C/5 or C/10) for 2 or 5 11 cycles, followed by leaving the cell in a charged state. The byproduct gases made during formation are allowed to leave the cell. After equilibration, formation. and degassing the cell is sealed and ready to be tested. Optimization of a lithium-ion cell As discussed in the previous section there are several parameters that affect the cell performance. However, there could be inter-dependence among the variables, and the contributions from the variable is not equal. Optimizing only one of the parameters under constant conditions is known as changing one variable at a time (OVT). The consequences of optimizing by OVT are: the excessive consumed time, the lack of extrapolation to other systems, and the risk of optimizing in a local minimum or maximum. It is evident, that the optimization of the lithium-ion battery relies on studying the critical variables of the system. For lithium-ion batteries the system is the cell; the parameters could be the active material particle size, the electrode porosity, the electrode tortuosity, the electrolyte donor-number etc. The responses could be the specific capacity, gravimetric energy, columbic efficiency, cycle life, etc. To identify the affecting parameter first we need to define the response of interest, then the parameters can be obtained by a statistical screening process. There have been several methodologies for the screening and system optimization. Design of Experiments (DoE) is a systematic experimental methodology for finding the relationships between the studied parameters that have an effect in the response of a process. DoE is based on statistical methods that analyze the contribution of the affecting variables. The motivation of using DoE comes from the inefficiency of using the method of OVT. Some of the main advantages of employing DoE is that it contemplates the 12 interactions among the variables, and that it requires less experiments saving time and resources for the research. Overcoming the trade-off between energy and power density. The main core of this dissertation is to understand the role of electrode and cell design parameters including new materials for building high-power and high-energy lithium-ion and lithium-sulfur batteries. One of the greatest challenges in lithium-ion batteries is the trade-off between energy and power density. 37 Typically, enhancing the energy density of the cell is reached by increasing the electrode thickness, however, a thicker cathode has a higher electrode tortuosity that hinders the cell’s rate capability and utilization. 38 In contrast, thinner electrodes allow the full utilization of the active material leading to an excellent rate capability at the expense of the cell’s energy density. Increasing the mixed conductivity of the electrode would in principle allow a better lithium-ion diffusion through the pores and a better interconnectivity of the intercalating particles narrowing the gap between energy and power in thicker electrodes. Therefore, to achieve a high-energy/high-power lithium-ion battery it is critical to increase the electrode electronic and ionic conductivity, maximize the amount of active material (areal mass loading), and minimize the quantity of the inactive components in the cell. Figure 1.3 illustrates a decision flowchart that summarizes the previous analysis for reaching a high-energy/high- power cell. 13 Figure 1.3. Evaluation flowchart for building a high-energy and high-power lithium battery. Lithium-Sulfur batteries To overcome the energy density limits of the lithium-ion battery, new high-energy electrochemical devices are needed. The lithium-sulfur (Li-S) battery stands-out as the next- generation battery candidate for high-energy storage. 39,40 Li-S batteries have several advantages and four times higher specific energy than Li-ion batteries with a high specific capacity of 1675 mAh g -1 , a high energy density of 2500 Wh kg -1 , a relative low cost, and environmental friendliness. 41,42 Figure 1.4 shows a bar chart comparison of the specific energy and capacity between the Li-ion and Li-S batteries with their corresponding active materials. 14 Figure 1.4 Specific capacity and energy density comparison among the of Li-ion and Li-S active materials and resulting cell. Li-S batteries have a metallic lithium negative electrode and a positive sulfur electrode. Specifically, the sulfur electrode is composed by a mixture of sulfur, carbon, and a polymeric binder. 43 On the other hand, metallic lithium forms the negative electrode and the amount of lithium usually exceeds greatly the maximum capacity of the sulfur electrode. Typically, Li-S battery electrolytes are composed by a mixture of ether-based solvents such as dioxolane (DOL) and dimethoxyethane (DME) with a lithium salt such as LiTFSI. 44 In contrast with other rechargeable batteries, Li-S is the only battery where the positive electrode dissolves and precipitates during the cell’s operation. These electrode phase changes start during the discharge when sulfur gets reduced and forms electrolyte-soluble high-order polysulfides that continue to be reduced into insoluble low-order polysulfides that lead to the formation of the solid Li 2S. During charge Li 2S gets oxidized into the soluble polysulfides and finally into sulfur. Figure 1.5 shows a schematic illustration of the Li-S battery and its general charge/discharge mechanism. 15 Figure 1.5 Schematic representation of a Li-S battery showing a simplification of the electrochemical processes that occur during the cell’s discharge/charge. Li-S batteries charge/discharge mechanism In liquid Li-S batteries the discharge process involves a solid-liquid-solid phase change. Multiple discharge/charge mechanisms have been proposed. 45–48 Specifically, at the beginning of discharge, metallic lithium gets oxidized forming Li + and migrates towards the sulfur electrode. Sulfur is reduced by the electrons released from the lithium electrode and forms high-order soluble polysulfides such as Li 2S 8, Li 2S 6 and Li 2S 4. As the reduction continues the soluble polysulfides precipitate into low-order insoluble polysulfides to finally form lithium sulfide Li 2S at the end of discharge. 49 During charging the system goes back from Li 2S into soluble polysulfides and then into insoluble sulfur. This discharge process involves sequential electron transfer steps involving 16 electrons transferred to every S 8 molecule. Three-fourths of the capacity comes from the precipitation of the insoluble sulfides. The discharge and charge curve of the Li-S battery 16 exhibits two plateaus; the first one between 2.3 and 2.4 V vs Li + /Li is related to the dissolution of sulfur, and the second plateau at approximately 2.05 V vs Li + /Li that corresponds to the precipitation of the low order polysulfides. Figure 1.6 illustrates the Li-S charge/discharge profiles divided in four regions and the region’s approximate capacity contribution. Different reaction mechanisms have been proposed for the discharge and charge process, the present dissertation assumes the same mechanism described by Moy et al. 46 Figure 1.6 Charge and discharge profiles for a Li-S battery. 4 regions are present; I) correspond to the dissolution of elemental sulfur into the high order polysulfides, II) soluble polysulfides reduction, III) reduction of the soluble polysulfides to the non-soluble polysulfides, and IV) Li 2S 2 to Li 2S. The discharge/charge mechanism is divided in four main steps: 46 Step I: Elemental Sulfur is reduced into Li 2S 8 that is soluble in the liquid electrolyte. In this step there is a phase change from solid to liquid. For this electrochemical reaction there is an exchange of 2 electrons. 𝑆 8 + 2𝐿𝑖 + + 2𝑒 − → 𝐿𝑖 2 𝑆 8 ( 1.4) 17 Step II: Li 2S 8 gets reduced to form Li 2S 6 that is also soluble in the electrolyte. No phase change is present on this step. The net electron exchange for this electrochemical reaction is 2. 3𝐿𝑖 2 𝑆 8 + 2𝐿𝑖 + + 2𝑒 − → 4𝐿𝑖 2 𝑆 6 ( 1.5) Step III: In the third step Li 2S 6 gets reduced to form Li 2S 4 that gets reduced further to produce Li 2S 2 and Li 2S that are insoluble in the electrolyte. The net electron exchange for this electrochemical reaction is 10. 2𝐿𝑖 2 𝑆 6 + 2𝐿𝑖 + + 2𝑒 − → 3𝐿𝑖 2 𝑆 4 ( 1.5) 𝐿𝑖 2 𝑆 4 + 2𝐿𝑖 + + 2𝑒 − → 2𝐿𝑖 2 𝑆 2 ( 1.6) 𝐿𝑖 2 𝑆 4 + 6𝐿𝑖 + + 6𝑒 − → 4𝐿𝑖 2 𝑆 ( 1.7) Step IV: Finally, Li 2S 2 gets completely reduced to form Li 2S, a solid-solid reaction. The net electron exchange for this electrochemical reaction is 2. 𝐿𝑖 2 𝑆 2 + 2𝐿𝑖 + + 2𝑒 − → 2𝐿𝑖 2 𝑆 ( 1.8) On the lithium electrode the cell the reaction that is present is given by 𝐿 𝑖 → 𝐿𝑖 + + 𝑒 − ( 1.9) 18 The polysulfide shuttle The polysulfide shuttle or simply shuttling, is a phenomena that occurs during the discharge and charge of a Li-S battery. 46,50 At the beginning of the discharge process the high-order soluble polysulfides tend to diffuse towards the lithium electrode due to the differences in the electrolyte’s concentration between the two electrodes. When the soluble polysulfides reach the surface of the lithium electrode they get chemically reduced into low-order polysulfides forming a new concentration gradient, then the newly formed polysulfides diffuse and migrate back towards the positive electrode. The process continues until the insoluble polysulfides are formed at the negative electrode passivating the surface of the lithium electrode. The process is called shuttling due to the continuous diffusion of the polysulfides between the electrodes. The shuttling causes the self-discharge of the battery and irreversible capacity losses. Figure 1.7 shows and scheme of the polysulfide process occurring in the Li-S battery. Figure 1.7 Schematic illustration of the polysulfide shuttling occurring in the Li-S battery. 19 Challenges and opportunities in Li-S batteries The practical application of Li-S batteries has not been achieved due to several challenges that affect the overall performance of the cell. 49 On the positive electrode, sulfur and lithium sulfide exhibit significant low electrical conductivity. Moreover, the sulfur electrode undergoes through a severe volume change of almost 80% during operation due to the differences in density between the active species, where Li 2S 8 and Li 2S have a density of 2.07 and 1.66 g cm -3 respectively. Furthermore, the negative electrode exhibits lithium dendrites and “dead” lithium that can lead to an electronic short across the cell and irreversible capacity losses. During charge and discharge the polysulfide shuttle is present affecting the cell’s reversibility. Finally, ether- base liquid electrolytes are flammable, constituting a safety concern. 40,49,51 These challenges affect the overall utilization of the sulfur electrode, the cell’s cycle life, rate capability, capacity retention, and coulombic efficiency. Therefore, there is a necessity to overcome each of these challenges to build a practical Li-S battery. Different strategies have been investigated in the last decades to achieve a practical cell. However, not all of the technical challenges have been addressed. Typically solving one of these challenges aggravates the others and a holistic solution is lacking. 43,52 It is crucial to understand the fundamental processes present in the Li-S battery and find the variables and their contribution in the actual cell’s performance. For example, the impact of the cathode architecture, areal loading of sulfur, and electrode composition are interrelated in determining the cell specific capacity, and rate capability. 53 Understanding the discharge and charge mechanism of Li-S constitutes the first step to rationalize the effect of the electrode 20 architecture (geometry, thickness, surface area), active material loading, electrolyte concentration, and electrolyte compatibility. Solid-State electrolytes Utilizing solid-state electrolytes such as polymeric, ceramic, composite, and gels have gained significant attention for solving the issues affecting Li–S batteries arising from organic liquid electrolytes. 54 Some of the advantages of utilizing solid-state electrolytes is the suppression of the polysulfide shuttle and the cell’s flammability. However, solid-state electrolytes exhibit other disadvantages such as a high interfacial impedance between the layers, lower ionic conductivity, low mechanical strength, and poor processability. 54 Ceramic electrolytes tend to have a higher mechanical strength and ionic conductivity compared to polymers and gels but lower processability and higher interfacial impedance. Briefly, the magnitude of the solid-electrolyte shear modulus determines if lithium dendrite growth will be suppressed or not. 51,55 It is known that polymeric electrolytes still exhibit dendrite formation after prolongated cycling. 56 Nevertheless, other challenges remain such as the cathode volume expansion. Therefore, utilizing composite solid-state electrolytes stands out as the most promising alternative that could combine the best properties of both ceramic and polymeric electrolytes and that can overcome at the same time all the issues in liquid electrolytes. 31 Figure 1.8 shows a scheme summarizing the Li-S electrolyte types and their challenges. 21 Figure 1.8 Challenges and electrolyte classification of lithium-sulfur batteries. 22 Hypotheses tested in the thesis. Lithium-ion batteries • Enhancing the conductivity of the porous electrode with new multi-functional materials such as conducting binders or faster intercalating materials will maximize the energy, power, and reversibility of a lithium-ion cell. • Utilizing composite solid-state electrolytes with high lithium-ion diffusivity, high conductivity, and mechanical strength will enhance the reversibility and cycle-life of a lithium-sulfur cell. Approach • Understand and characterize the electrochemical properties of new conducting polymers that could serve as cathode binders. • Increase the rate capability and capacity of the cell by utilizing π-conjugated conducting polymers with enhanced mixed conductivity as cathode binders. • Design and build a mass-efficient lithium-ion battery in conjunction with the conducting polymers for reaching a high-energy/high-power cell. • Tune the polymer ionic conductivity by introducing oligoether functional groups in the π- conjugated conducting polymers to reach higher rates and electrodes with a lower binder mass fraction. • Understand and characterize the electrochemical properties of new and faster negative electrode materials with high Li + diffusivity and structural stability. 23 • Pair the new anode materials with the cathodes made with conducting polymer binders to build a fast lithium-ion “full” cell. Lithium-sulfur batteries: Objectives and Hypotheses • Understand the effect of carbon and sulfur loading in the electrode design of lithium- sulfur batteries. • Develop a new bilayer composite solid-state electrolyte made of an intercalating material and a polymer electrolyte that can effectively suppress the polysulfide shuttle, stop the dendrite formation, and eliminate the electrolyte flammability. Innovation The data and analysis presented in this dissertation provide a novel comprehensive body of scientific and technological evidence regarding the electrode and cell design for lithium-ion and lithium-sulfur batteries. This dissertation attempts to identify, capture, and understand the variables affecting the cell performance to elucidate the paths for building next-generation lithium batteries. The novelty comes from converging mass-efficient cell designs and new electrode architectures/compositions with novel materials such as conducting polymers, composite solid-electrolytes and faster Li + intercalating compounds. In addition, it demonstrates for the first time that the trade-off between power and energy can be overcome in a lithium-ion battery. Finally, we introduce a new solid-state lithium-sulfur battery that operates at room temperature by utilizing a composite electrode and bilayer ceramic/polymeric solid-electrolyte. 24 Importance of Electrochemical Techniques The cell’s electrochemistry is the most fundamental and core characteristic that dictates the cell’s properties and performance. Characterizing and analyzing the cell’s electrochemical response brings an understanding to the physicochemical phenomena involved. In this work, electrochemical and physicochemical characterization techniques have been utilized in depth such as Electrochemical Impedance Spectroscopy (EIS), Incremental Capacity Analysis (ICA), Cycling Voltammetry (CV), Electron Paramagnetic Resonance (EPR), and spectroscopy techniques. The present dissertation shows in-depth electrochemical cell and electrode characterization, a rationalization of phenomena, and an exhaustive amount of data that shows the benefits of utilizing multi-functional new materials and cell configurations. An electrochemical cell is a complex device affected by several interrelated factors such as the intrinsic material properties, geometrical configurations, transport phenomena at the nano and macro scale, etc. Understanding all the variables and their interrelation constitutes one of the greatest challenges in electrochemistry and battery technology. This work attempts to provide a rational analysis and performance data of various materials and their role in the electrode and cell design for lithium-ion and lithium-sulfur batteries. References 1. Bilgen, S. Structure and environmental impact of global energy consumption. Renewable and Sustainable Energy Reviews 38, 890–902 (2014). 2. Jackson, R. B. et al. Global energy growth is outpacing decarbonization. Environ. Res. Lett. 13, 120401 (2018). 3. Bradshaw, M. J. Global energy dilemmas: a geographical perspective: Global energy dilemmas: a geographical perspective. Geographical Journal 176, 275–290 (2010). 25 4. Ellis, E. C. et al. Used planet: A global history. Proceedings of the National Academy of Sciences 110, 7978–7985 (2013). 5. Chen, T. et al. Applications of Lithium-Ion Batteries in Grid-Scale Energy Storage Systems. Trans. Tianjin Univ. 26, 208–217 (2020). 6. Gür, T. M. Review of electrical energy storage technologies, materials and systems: challenges and prospects for large-scale grid storage. Energy Environ. Sci. 11, 2696–2767 (2018). 7. Hadjipaschalis, I., Poullikkas, A. & Efthimiou, V. Overview of current and future energy storage technologies for electric power applications. Renewable and Sustainable Energy Reviews 13, 1513–1522 (2009). 8. Ibrahim, H., Ilinca, A. & Perron, J. Energy storage systems—Characteristics and comparisons. Renewable and Sustainable Energy Reviews 12, 1221–1250 (2008). 9. Whittingham, M. S. History, Evolution, and Future Status of Energy Storage. Proc. IEEE 100, 1518–1534 (2012). 10. Dunn, B., Kamath, H. & Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices. Science 334, 928–935 (2011). 11. Reddy, M. V., Mauger, A., Julien, C. M., Paolella, A. & Zaghib, K. Brief History of Early Lithium-Battery Development. Materials 13, 1884 (2020). 12. Weeks, M. E. The discovery of the elements. IX. Three alkali metals: Potassium, sodium, and lithium. J. Chem. Educ. 9, 1035 (1932). 13. Brande, W. T. A Manual of Chemistry. vol. 2 (John Murray, 1821), London. 14. Lewis, G. N. & Keyes, F. G. The Potential of the Lithium Electrode. J. Am. Chem. Soc. 35, 340–344 (1913). 15. Harris, W. S. Electrochemical Studies in Cyclic Esters. (University of California Berkeley, 1958), USA. 16. Mizushima, K., Jones, P. C., Wiseman, P. J. & Goodenough, J. B. LixCoO 2 (0<x~l): A New Cathode Material for Batteries of High Energy Density. Materials Research Bulletin 15, 7 (1980). 17. Blomgren, G. E. The Development and Future of Lithium Ion Batteries. J. Electrochem. Soc. 164, A5019–A5025 (2017). 18. Yoshino, A. The Lithium-ion Battery: Two Breakthroughs in Development and Two Reasons for the Nobel Prize. BCSJ 95, 195–197 (2022). 19. Yoshino, A., Nakajima, T. & Sanechika, K. Secondary battery, USP 4688595. 20. Nishi, Y. The Dawn of Lithium-Ion Batteries. Interface magazine 25, 71–74 (2016). 21. Nagaura, T. & Tozawa, K. Progress in batteries and solar cells. vol. 9 (1990). 26 22. Schipper, F. & Aurbach, D. A brief review: Past, present and future of lithium ion batteries. Russ J Electrochem 52, 1095–1121 (2016). 23. Brédas, J.-L. et al. An Electrifying Choice for the 2019 Chemistry Nobel Prize: Goodenough, Whittingham, and Yoshino. Chem. Mater. 31, 8577–8581 (2019). 24. Kamat, P. V. Lithium-Ion Batteries and Beyond: Celebrating the 2019 Nobel Prize in Chemistry – A Virtual Issue. ACS Energy Lett. 4, 2757–2759 (2019). 25. Doeff, M. M. 2019 Nobel Prize in Chemistry Winners. The Electrochemical Society Interface 8–17 (2019). 26. Hu, Y.-S. & Lu, Y. 2019 Nobel Prize for the Li-Ion Batteries and New Opportunities and Challenges in Na-Ion Batteries. ACS Energy Lett. 4, 2689–2690 (2019). 27. Smart, M. C. et al. Improved performance of lithium-ion cells with the use of fluorinated carbonate-based electrolytes. Journal of Power Sources 119–121, 359–367 (2003). 28. Zhang, Z. et al. Fluorinated electrolytes for 5 V lithium-ion battery chemistry. Energy Environ. Sci. 6, 1806 (2013). 29. Fu, J. et al. Electrically Rechargeable Zinc-Air Batteries: Progress, Challenges, and Perspectives. Adv. Mater. 29, 1604685 (2017). 30. Aravindan, V., Gnanaraj, J., Madhavi, S. & Liu, H.-K. Lithium-Ion Conducting Electrolyte Salts for Lithium Batteries. Chem. Eur. J. 17, (2011). 31. Dirican, M., Yan, C., Zhu, P. & Zhang, X. Composite solid electrolytes for all-solid-state lithium batteries. Materials Science and Engineering: R: Reports 136, 27–46 (2019). 32. Berg, E. J., Villevieille, C., Streich, D., Trabesinger, S. & Novák, P. Rechargeable Batteries: Grasping for the Limits of Chemistry. J. Electrochem. Soc. 162, A2468–A2475 (2015). 33. Kim, U.-H. et al. Pushing the limit of layered transition metal oxide cathodes for high- energy density rechargeable Li ion batteries. Energy Environ. Sci. 11, 1271–1279 (2018). 34. Radin, M. D. et al. Narrowing the Gap between Theoretical and Practical Capacities in Li- Ion Layered Oxide Cathode Materials. Adv. Energy Mater. 7, 1602888 (2017). 35. Cao, Y., Li, M., Lu, J., Liu, J. & Amine, K. Bridging the academic and industrial metrics for next-generation practical batteries. Nat. Nanotechnol. 14, 200–207 (2019). 36. Liu, J. et al. Pathways for practical high-energy long-cycling lithium metal batteries. Nat Energy 4, 180–186 (2019). 37. Zheng, H., Li, J., Song, X., Liu, G. & Battaglia, V. S. A comprehensive understanding of electrode thickness effects on the electrochemical performances of Li-ion battery cathodes. Electrochimica Acta 71, 258–265 (2012). 27 38. Du, Z., Wood, D. L., Daniel, C., Kalnaus, S. & Li, J. Understanding limiting factors in thick electrode performance as applied to high energy density Li-ion batteries. J Appl Electrochem 47, 405–415 (2017). 39. Kang, W. et al. A review of recent developments in rechargeable lithium–sulfur batteries. Nanoscale 8, 16541–16588 (2016). 40. Manthiram, A., Fu, Y., Chung, S.-H., Zu, C. & Su, Y.-S. Rechargeable Lithium–Sulfur Batteries. Chem. Rev. 114, 11751–11787 (2014). 41. Zhao, H. et al. A review on anode for lithium-sulfur batteries: Progress and prospects. Chemical Engineering Journal 347, 343–365 (2018). 42. Albertus, P., Babinec, S., Litzelman, S. & Newman, A. Status and challenges in enabling the lithium metal electrode for high-energy and low-cost rechargeable batteries. Nat Energy 3, 16– 21 (2018). 43. Eftekhari, A. & Kim, D.-W. Cathode materials for lithium–sulfur batteries: a practical perspective. J. Mater. Chem. A 5, 17734–17776 (2017). 44. Wang, Y., Sahadeo, E., Rubloff, G., Lin, C.-F. & Lee, S. B. High-capacity lithium sulfur battery and beyond: a review of metal anode protection layers and perspective of solid-state electrolytes. J Mater Sci 54, 3671–3693 (2019). 45. Barchasz, C. et al. Lithium/Sulfur Cell Discharge Mechanism: An Original Approach for Intermediate Species Identification. Anal. Chem. 84, 3973–3980 (2012). 46. Moy, D., Manivannan, A. & Narayanan, S. R. Direct Measurement of Polysulfide Shuttle Current: A Window into Understanding the Performance of Lithium-Sulfur Cells. J. Electrochem. Soc. 162, A1–A7 (2015). 47. Waluś, S., Barchasz, C., Bouchet, R. & Alloin, F. Electrochemical impedance spectroscopy study of lithium–sulfur batteries: Useful technique to reveal the Li/S electrochemical mechanism. Electrochimica Acta 359, 136944 (2020). 48. Wang, D. R. et al. Rate Constants of Electrochemical Reactions in a Lithium-Sulfur Cell Determined by Operando X-ray Absorption Spectroscopy. J. Electrochem. Soc. 165, A3487–A3495 (2018). 49. Fotouhi, A., Auger, D., O’Neill, L., Cleaver, T. & Walus, S. Lithium-Sulfur Battery Technology Readiness and Applications—A Review. Energies 10, 1937 (2017). 50. Kaewruang, S. et al. Strong adsorption of lithium polysulfides on ethylenediamine- functionalized carbon fiber paper interlayer providing excellent capacity retention of lithium- sulfur batteries. Carbon 123, 492–501 (2017). 51. Wang, C. et al. Suppression of Lithium Dendrite Formation by Using LAGP-PEO (LiTFSI) Composite Solid Electrolyte and Lithium Metal Anode Modified by PEO (LiTFSI) in All-Solid-State Lithium Batteries. ACS Appl. Mater. Interfaces 9, 13694–13702 (2017). 28 52. Chen, X., Hou, T., Persson, K. A. & Zhang, Q. Combining theory and experiment in lithium– sulfur batteries: Current progress and future perspectives. Materials Today 22, 142–158 (2019). 53. Peng, H.-J., Huang, J.-Q., Cheng, X.-B. & Zhang, Q. Review on High-Loading and High- Energy Lithium-Sulfur Batteries. Adv. Energy Mater. 7, 1700260 (2017). 54. Hassoun, J. & Scrosati, B. Moving to a Solid-State Configuration: A Valid Approach to Making Lithium-Sulfur Batteries Viable for Practical Applications. Adv. Mater. 22, 5198–5201 (2010). 55. Chen, L. et al. PEO/garnet composite electrolytes for solid-state lithium batteries: From “ceramic-in-polymer” to “polymer-in-ceramic”. Nano Energy 46, 176–184 (2018). 56. Brissot, C., Rosso, M., Chazalviel, J.-N., Baudry, P. & Lascaud, S. In situ study of dendritic growth inlithium/PEO-salt/lithium cells. Electrochimica Acta 43, 1569–1574 (1998). 29 Chapter 2 – Experimental methods Abstract This chapter discusses and details the experimental methods for each chapter. Specifically, it includes the material preparation for binders, material characterization and dopability, conductivity measurements, interactions with solvent, stability to repeated cycling, electrode fabrication, cell assembly, and cell testing. The electrochemical characterization of NaNb 13O 33 electrodes, including cell assembly and performance testing are detailed. Furthermore, this chapter describes the methods for the electrode preparation and cell assembly of the lithium- sulfur studies and the preparation of the bilayer composite solid-state electrolyte. Methods for characterizing PProDOT-Hx2 as cathode binder (Chapter 3) Polymer Synthesis • Acknowledgment: Polymer materials were supplied by Dr. Barry Thompson’s lab through the efforts of Pratyusha Das • Objectives: Synthesize PProDOT-Hx 2 conjugated polymers. • Detailed procedures can be obtained from Das et al. 2020. 1 All small molecules except compound S1 and both the monomers S3 and S4 were prepared following literature procedure. 2 30 Scheme 2.1. Synthesis of Monomers (S3 and S4). Direct Arylation Polymerization (DArP) of PProDOT-Hx 2: PProDOT-Hx 2 Scheme 2.2. Polymerization of PProDOT-Hx 2 using DArP. PProDOT-Hx 2 was obtained as a shiny purple solid in 76% yield which was then dried under vacuum overnight at room temperature. NMR data is consistent with literature reports. 2 Molecular Characterization All small molecules and monomer NMR were recorded at 25°C using CDCl3 on either a Varian Mercury 400 MHz or Varian VNMRS-500 MHz. Polymer NMR was obtained at 25°C using CDCl3 on a Varian VNMR-500 MHz. All NMR data are consistent with literature reports.1 Number average molecular weight (Mn) and polydispersity (Ð) were determined by size exclusion 31 chromatography (SEC) using a Waters Alliance HPLC System with 2690 Separation Module and a Waters 2410 Differential Refractometer (RI) detector, with HPLC grade chloroform (CHCl3) with 0.25% triethylamine (TEA) as eluent at a flow rate of 0.5 mL/min on three Acquity APC XT Columns (45 + 200 + 450 pore sizes). The instrument was calibrated vs. polystyrene standards (200−400,000 g/mol), and data were analysed using Empower 3 software. Polymer samples were dissolved in HPLC grade CHCl3 at a concentration of 3-4 mg/mL, stirred until dissolved, and filtered through a 0.45 μm PTFE filter. Figure 2.1 GPC analysis in Chloroform with calibration to a polystyrene standard: Mn = 19,182 g mol-1; Mw =31,820 g mol-1; PDI = 1.6 GIWAXS and Ellipsometry of PProDOT-Hx 2 • Acknowledgment: Grazing Incidence Wide-Angle X-ray Scattering (GIWAXS) and Ellipsometry characterization was carried out in Dr. Sarah Tolbert’s lab from UCLA through the efforts of Charlene Salamat. • Objectives: To understand the structural changes induced by doping. 32 Neutral films for 2-D GIWAXS measurements were prepared on 1.5 × 1.5 cm silicon substrates. Electrochemically doped films were prepared on 1 cm × 2 cm silicon substrates with <100> orientation, resistivity of 0.001-0.005 Ω/cm, and with 100 nm of aluminum evaporated on the back-side. Measurements were performed at the Stanford Synchrotron Radiation Lightsource on beamline 11-3 using a wavelength of 0.9742 Å with an incidence angle of 0.12°. To collect the diffraction patterns, a helium chamber was used with a detector distance of 250 mm and spot size of ~150 μm. The IgorPro macro, Nika was used to calibrate and reduce the GIWAXS data. All analysis was performed in IgorPro. Each integration pattern is background corrected. Ellipsometry was performed on a PS-1100 instrument from Semilab at room temperature. The instrument used a UV-visible CCD detector adapted to a grating spectrograph to analyze the signal reflected by the sample. The light source was a 75 W Hamamatsu Xenon lamp. Films for spectroscopic ellipsometry were prepared on 1.5 × 1.5 cm ITO substrates. Each film was measured both outside and inside of a home-built custom vial used for solvent swelling. First, the thickness of the film was measured outside of the vial as a control and then inside of the vial to ensure the thickness was the same. The films were prepared less than 1-hour prior to the measurement. The home-made vial enables in-situ swelling experiments with solvents that are incompatible with this elipsometeric porosimeter, such as the propylene carbonate used here. The vial has a height of 30 mm, inner diameter of 18 mm, and outer diameter of 20 mm. There are two holes 3 mm in diameter, 30 mm from the bottom of the vial that allows light to come in and out while only interacting with the film sample. Data analysis was performed using the associated Spectroscopic Ellipsometry Analyzer software. 33 Performance as an NCA Cathode Binder Objectives: Test the performance of PProDOT-Hx 2 as cathode binder. To prepare PProDOT-Hx 2-NCA electrodes, a slurry composed of NCA powders (LiNi 0.8Co 0.15Al 0.05O 2, Quallion Corp., Sylmar, CA), Super P carbon black (Sigma-Aldrich), multiwalled carbon nanotubes (CNTs, OD×ID×L: 10 nm × 4.5 nm × 3-6 μm), and the PProDOT- Hx 2/ODCB solution (20 g L -1 ) respectively were mixed in a weight ratio of 90:3:3:4. The CNT (3 wt %) extends through the entire electrode to provide long-range electronic conductivity. 3 To prepare the PVDF-NCA control electrode, polyvinylidene fluoride (PVDF, Sigma-Aldrich) was dissolved in N-methyl-2-pyrrolidone (NMP, Sigma-Aldrich), and mixed with NCA, Super P carbon black, CNTs in the same ratio. For the electrodes without binder a weight ratio of 94:3:3 was employed for NCA, Super P and CNT respectively utilizing NMP as the solvent. The slurries were coated onto carbon@aluminum foil using doctor blading. The NCA electrodes were dried in a vacuum oven at 120 ℃ for 12 hours, then punched into discs with a diameter of 14 mm. Two different loadings were tested with an approximate mass loading of NCA in the electrode of about ~6 mg cm 2 and 11 mg cm -2 . The electrochemical performance of NCA- PProDOT-Hx 2, NCA-PVDF, and NCA without binder was tested using the same electrolyte in CR2032 coin-type cells with a Li counter electrode, Celgard 2325 separator, and 50 μL of electrolyte. The specific capacity (mAh/g) was based on the weight of the NCA in the electrode, and the capacity contribution from PProDOT-Hx 2 is negligible. Cyclic voltammetry (CV) and galvanostatic charge-discharge (GCD), were measured in a VMP-3 (Bio-Logic) potentiostat/galvanostat. For the GCD testing, different C-rates were used based on 1 C = 160 mA g −1 as previously reported. 3 34 Ionic and Electronic Conductivity of PProDOT-Hx 2 • Acknowledgment: Conductivity measurements were carried by Billal Zayat from USC. • Objectives: To characterize the electronic and ionic conductivity of the polymer as function of doping (potential). Electrode Preparation. A 20 mg/mL PProDOT-Hx 2 polymer solution was prepared by dissolving the polymer powder in 1,2-dichlorobenzene (99%, sigma aldrich). The solution was then heated at 40°C under argon for two hours to insure complete dissolving of the polymer powder in the solvent. 5 μL of the prepared solution was then spin-coated on a gold interdigitated microelectrode on a glass substrate (Metrohm Dropsens) at 1000 RPM for 30 seconds. The prepared electrode was then annealed under vacuum at 110℃ for two hours. Similarly, commercially available P3HT powder (Sigma Aldrich) was used to prepare thin films on gold interdigitated electrodes. Electrochemical Doping. A 3-electrode cell was assembled (Figure 2.2b) under nitrogen to electrochemically dope the polymer and determine its ionic conductivity. The two gold microelectrodes were shorted to form the working electrode while lithium foil was used as the counter and working electrodes. 1 M LiTFSI in a 1:1 by volume mixture of ethylene carbonate and dimethyl carbonate was used as the electrolyte. To electrochemically dope the polymer, a voltage was held at the working electrode for 300 seconds to ensure complete electrochemical doping at that voltage. Cycling voltammetry scans were also performed between 3.0 V and 4.2 V vs Li/Li + at various scan rates to observe the electrochemical doping and de-doping process and identify the neutral-polaron and polaron-bipolaron transitions 35 Figure 2.2. (a) SEM image of the spin-coated PProDOT-Hx 2 on interdigitated gold microelectrodes. The light color shows the gold electrodes and the dark color shows the polymer imbedded between the electrodes. (b) The 2-electrode configuration used to measure the electronic conductivity of the polymer as a function of voltage. (c) Cyclic voltammetry plot of PProDOT-Hx 2 at various scan rates. (d) The 3-electrode configuration used to electrochemically dope the polymer and determine its ionic conductivity. Ionic Conductivity Measurement. Potentiostatic electrochemical impedance spectroscopy of the 3-electrode cell was collected at the end of the potentiostatic hold at the same voltage as that of the electrochemical doping without any interruption. Impedance was collected between 100 mHz and 100 kHz with an excitation of 10 mV. The potentiostatic EIS measurement was repeated at different voltages to determine the change in ionic conductivity as a function of voltage. The ionic resistance and ionic conductivity are represented by R ionic and σ ionic, respectively. The polymer film thickness is represented by h, the length and the number of the digits of the gold 36 electrode are represented by l and N, respectively, and the distance between the gold electrodes is represented by d. 𝜎 𝑖𝑜𝑛𝑖𝑐 = 1 𝑅 𝑖𝑜𝑛 𝑖𝑐 × ℎ 𝑙 × ( 𝑁 − 1)× 𝑑 ( 2.1) Electronic Conductivity Measurement. To measure R e accurately, the use of a different configuration for impedance measurement is necessary. In this configuration, the impedance between the two interdigitated gold electrodes is measured. Electrochemical doping of the polymer was performed in the 3-electrode configuration as before. The cell was then allowed to relax for 100 seconds to reach a steady state potential. The cell setup was then switched from a 3-electrode configuration to a 2-electrode configuration in which the gold microelectrodes are the two electrodes used (Figure 2.2d). EIS measurement was then performed at open circuit potential between 100 kHz and 100 mHz at an excitation of 10 mV. Electronic conductivity, σ e, can be calculated using equation S3 where the parameters are defined as in equation 2.2. Note that d in this configuration is in the numerator since the current flow is between the two gold electrodes. 𝜎 𝑒 = 1 𝑅 𝑒 × 𝑑 𝑙 × ( 𝑁 − 1)× ℎ ( 2.2) Fabrication and Electrochemical Characterization of PProDOT-Hx 2 Thin Films • Acknowledgment: Film fabrication was carried out by Dr. Billal Zayat of Prof. Narayan’s Lab at USC, and by Dr. Bruce Dunn’s lab from UCLA through the efforts of Qiulong Wei. • Objectives: To understand the structural changes induced by doping. A 20 g L -1 solution of PProDOT-Hx 2 was dissolved in 1,2-dichlorobenzene (ODCB, Sigma-Aldrich). To prepare a thin film, 20 μL of the solution was spin-coated at a speed of 4000 rpm for 60 s onto 37 a 1 × 1 cm 2 fluorine doped tin oxide (FTO, 7 Ω/sq, Sigma-Aldrich) coated glass. The resulting film was approximately 50 nm as measured by profilometry. The thin-film PProDOT-Hx 2@FTO was transferred into a vacuum oven at 120 ℃ for 4 hours and then transferred into an Ar-filled glovebox. The electrochemical doping of PProDOT-Hx 2 was tested in a three-electrode cell, thin- film PProDOT-Hx 2@FTO as the working electrode and two lithium foils as counter and reference electrodes. The utilized electrolyte is 1 M LiTFSI in ethylene carbonate (EC)/dimethyl carbonate (DMC) in a 1:1 volume ratio. Electron paramagnetic resonance • Acknowledgment: Electron paramagnetic resonance characterization was carried by Dr. Rachel Segalman’s from UCSB through the efforts of Gordon Pace, Dakota Rawlings and Dongwook Lee. • Objectives: To confirm the nature of charge carriers in the binder at the molecular level using spin counting. 1.0 mg/mL of PProDOT-Hx 2 solution in chloroform was dropcast on ITO-coated PET substrate under nitrogen to form EPR specimens. They were further dried in vacuum at 60 C for 1 hour, in order to prevent mechanical weakening of the PET substrate above its glass transition temperature. After cooling down to room temperature, the PProDOT-Hx 2 films were transferred to Ar-filled glove box for electrochemical characterizations. Each film was employed as a working electrode in a beaker cell, where 2 separate pieces of Li metal served as counter and reference electrodes, respectively. The electrolyte was chosen 1.0 M LiPF 6 in EC:DMC (50:50 v/v). The beaker cell was sealed with parafilm to prevent evaporation of the electrolyte. 38 All electrochemical tests were performed under Ar. The beaker cells were firstly put under open- circuit condition for 9 hours in order to allow the cell to equilibrate, so that charge can be counted precisely during the following electrochemical tests. After the cell equilibration, repeated CV was run at 2.9 V – 4.2 V vs. Li/Li + , with sweeping rate = 1 mV/s. Such low sweeping rate was selected to minimize influence of overpotential, so that charge count can become accurate. When CV curves from different cycles converge, the cell was allowed to rest under open-circuit condition. At this stage, spontaneous self-discharge (de-doping) of the PProDOT-Hx 2 was observed; open- circuit potential of the cell returns to 2.95 V vs. Li/Li + after sufficient time. At the last stage of electrochemistry, the cell was galvanostatically oxidized to a target potential, immediately followed by chronoamperometric oxidation at the same target potential. The working electrode was dismounted from the cell and rinsed with DMC 3 times. Finally, the rinsed PProDOT-Hx 2 film was dried under vacuum at room temperature for 1 hour to avoid possible de-doping at elevated temperatures. Integration of the current in CV measurements gives the total charge (Q) via equation S1 where dV/dt is the voltage sweep rate and I ox represents oxidation current. 4 Q = ∫ I ox ( dV/dt) dV ( 2.3) Following the doping process, samples were inserted into 5 mm-diameter EPR tubes and sealed with plastic caps and parafilm in a nitrogen filled glovebox. They were then transferred outside of the glovebox and EPR spectra were measured within 1 hour. Spin concentrations were determined by comparing the normalized, integrated signal intensity of a sample with the normalized, integrated signal intensity for a series of standards at known spin concentration. 2,2- Diphenyl-picrylhydrazyl (DPPH) was used as a calibrant for determining the spin concentration. 39 The calibrant was mixed with chloroform with known concentrations and cast onto PET substrates to obtain a known quantity of DPPH on the substrate. EPR was then measured on the standard samples as described above. See Figure 2.1 for the calibration curve. Figure 2.3 The normalized, integrated EPR signal intensity for a series of DPPH calibration samples with known spin concentrations. Methods for High-Energy/High-Rate π-Conjugated Polymers as Binders for Lithium Metal Batteries (Chapter 4) Monomer and Polymer Synthesis • Acknowledgment: Polymer materials were supplied by Dr. Barry Thompson’s lab through the efforts of Pratyusha Das • Objectives: Synthetize the conjugated polymers. • Detailed procedures can be consulted in Elizalde-Segovia et al. 2021. 5 All chemical reactions were performed in oven-dried glassware under dry N 2, unless otherwise noted. Monomers S1-S4 (Figure 2.4) were prepared following established literature procedure. 6,7 All polymers, P3HT (P1), PProDOT-Hx 2 (P2), and (Hex:OE)(80:20) (P3) (Figure S1) were prepared using Direct Arylation Polymerization (DArP). 6 40 Synthesis of ProDOT-OE2 (S5): Compound S5 was prepared by modifying a reported literature procedure. 8 Synthesis of P3HT (P1): The polymer P3HT was prepared according to a reported literature procedure. 6 Synthesis of PProDOT-Hx 2 (P2): The polymer PProDOT-Hx 2 was prepared by modifying literature procedure 1 consistent with literature reports. 7 Synthesis of (Hex:OE)(80:20) PProDOT Random Copolymer (P3): The (Hex:OE) PProDOT random copolymer was prepared by following literature procedure. 1 The conducting polymers are sensitive to air oxidation upon extended exposure. All polymers were stored under nitrogen at low temperature and in the dark prior to use. 41 Figure 2.4 Synthesis of Monomer (S5) and Polymers P3HT (P1), PProDOT-Hx 2 (P2) and (Hex:OE)(80:20) PProDOT Random Copolymer (P3). Polymer physical Characterization All small molecule and monomer NMR spectra were recorded at 25 °C using CDCl 3 on either a Varian Mercury 400 MHz or Varian VNMRS-500 MHz. Polymer NMR was obtained on a Varian VNMR-500 MHz at 25 °C using CDCl 3. All spectra were referenced to CHCl 3 (7.26 ppm). Number average molecular weight (Mn) and polydispersity (Ð) were determined by size exclusion chromatography (SEC) using Agilent 1260 Infinity II High Temperature GPC with a Differential Refractive Index (DRI) detector with 80 °C HPLC grade 1,2,4-trichlorobenzene (o-DCB) as eluent at a flow rate of 1 mL/min. The instrument was calibrated vs. polystyrene standards (1050−3,800,000 g/mol), and data were analyzed using Agilent GPC/SEC software. Polymer 42 samples were dissolved in HPLC grade 1,2,4-trichlorobenzene (TCB) at a concentration of 0.5 mg ml −1 , stirred at 80°C until dissolved and filtered through a 0.2 μm PTFE filter. For PProDOT-Hx 2, GPC was measured on a Waters e2695 Separations Module system using Agilent PLgel columns. Refractive index traces were used for molecular weight determination using polystyrene calibration standards (Agilent Technologies, EasiVial) and were collected in 35°C in tetrahydrofuran with a flow rate of 0.3 mL/min. Monomer NMR Figure 2.5. 1 H-NMR of S5 in CDCl 3 at 25 °C and 500 MHz. 43 Polymer NMR Figure 2.6. 1 H-NMR of P1 (P3HT) in CDCl 3 at 25 °C and 500 MHz. Figure 2.7. 1 H-NMR of P2 (PProDOT-Hx 2) in CDCl 3 at 25 °C and 500 MHz 44 Figure 2.8. 1 H-NMR of P3 ((Hex:OE) (80:20) PProDOT Random Copolymer) in CDCl 3 at 25 °C and 500 MHz. Conductivity measurements • Acknowledgment: Conductivity measurements were carried out by Billal Zayat from USC. • Objectives: To characterize the electronic and ionic conductivity of the polymers as function of doping (potential). Interdigitated microelectrodes (IDM) were purchased from Metrohm Dropsens (DRP-G-IDEAU5- U20). Each microelectrode is composed of 250 digits with a digit length of 6760 µm and a gap of 5 µm between the digits. Polymer solutions were prepared by dissolving the polymer powder in 1,2-dichlorobenzene (99%, Sigma-Aldrich). The solution was then heated under argon for two hours at 40°C. 5 μL of the prepared solution was spin-coated on the gold IDM at 1000 RPM for 30 seconds to produce a 50 nm polymer film measured by SEM. 1,9 The prepared electrodes were then annealed under vacuum at 110℃ for two hours, then transferred to an argon glovebox. All electrochemical tests were performed in an argon glovebox at room temperature. The details of 45 the experimental technique for electrochemical doping and the determination of ionic and electronic conductivity has been previously reported by our group 9 . Briefly, to electrochemically dope the polymers, a 3-electrode cell was assembled using the polymer on the IDM as the working electrode, lithium foil as the counter and reference electrodes, and 1 M LiPF 6 in a mixture of 1:1 by volume of ethylene carbonate and diethyl carbonate (EC/DEC) as the electrolyte. EIS of the 3-electrode cell was measured after electrochemical doping at each potential to obtain the ionic conductivity of the polymer films between 100 mHz and 100 kHz at a sinusoidal excitation of ± 10 mV peak-to-peak. To determine electronic conductivity, impedance was measured between the two terminals of the gold electrodes at open circuit potential between 100 mHz and 100 kHz at a sinusoidal excitation of ± 10 mV peak-to-peak. Cathode Preparation Lithium-ion battery cathodes consisting of lithiated nickel cobalt aluminum oxide (LiNi 0.8Co 0.15Al 0.05O 2 (NCA) and π-conjugated polymer were prepared by mixing pristine powders of NCA (NEI corporation), Super P carbon black (MTI), multiwalled carbon nanotubes (CNTs, Cnano, OD×ID×L: 25 nm ×10 nm × 10 μm), and the π-conjugated polymer/1,2-Dichlorobenzene (ODCB) (50 g L -1 ) in a weight ratio of 95:0.5:0.5:4. The cathode powder was gently mixed in a mortar and pestle before being added to the ODCB solution. Control NCA-PVDF electrodes were fabricated by dissolving polyvinylidene fluoride (PVDF, MTI) in N-methyl-2-pyrrolidone (NMP, Sigma-Aldrich) and mixed with NCA, Super P carbon black, and CNTs in the same ratio. The addition of CNTs creates a mesoscale porosity that improves the electrolyte penetration into the electrode structure. 10–12 The resulting slurry was coated onto an aluminum foil using the doctor blade technique. 13 The electrode films were vacuum-dried overnight, roll pressed (12’’ vise- 46 mount slip roll T10727, Grizzly Industrial) and then punched to 14 mm diameter discs. Areal active material mass loading of the electrodes was approximately 14 mg cm -2 . Lithium Anode Preparation For the cells under mass-efficient conditions, metallic lithium chips (MTI, D×T :16 mm × 0,6 mm) were polished, hammered and roll-pressed between two layers of pure nitrile sheets to obtain an extremely thin lithium foil (46 µm thick). The foil was carefully punched into 14 mm diameter lithium disks. The thinning process was repeated multiple times until obtaining the desired weight of lithium based on an N/P ratio of ≈ 3. For the non-limited cells, lithium chips with a weight of ≈ 40 mg were rendered shiny by scratching the surface with a tweezer and used as is. Lithium handling was carried out in an argon-filled glove box (VAC system 60387, NexGen 2P, Vacuum Atmospheres Company) with less than 0.5 ppm of moisture and 0.2 ppm of oxygen. Electrolyte All cells utilized commercial carbonate electrolyte consisting of 1.0 M Lithium hexafluorophosphate in 1:1 in ethylene carbonate (EC) and diethyl carbonate (DEC) from Sigma- Aldrich. Electrode Characterization The microstructure of the electrodes was examined with a scanning electron microscope (Nova NanoSEM 450 Field Emission SEM). Elemental mapping (EDS) was performed using an energy dispersive X-ray spectrometer JSM-7001F-LV. SEM images for each electrode type in the scale of 100 microns and 500 nm. The macro and meso pores at the surface were characterized using ImageJ (open source software NIH) with a Bandpass filter with a maximum of 500 pixels and a 47 minimum of 6 pixels for each image followed by a particle analysis that included the image edges and bare outlines. Imbibition/Archimedes’ method was utilized for estimating the mean effective porosity of the electrodes using the average value of triplicate measurements using isopropyl alcohol as the fluid. The imbibition method to determine the pore volume consists of saturating the dried electrode with a liquid of known density followed by weighing. The Archimedes’ method is used to determine the electrode volume by immersing the electrode in a fluid of known density (subtracting the volume of the current collector). The mean effective porosity is then obtained by dividing electrode volume by the pore volume. Cell assembly CR2032 coin-type cells were fabricated with the NCA-π-conjugated polymer electrode as the working electrode, metallic lithium as the counter electrode, and Celgard 2325 (PP/PE/PP) as the separator. The amount electrolyte and lithium for the mass-efficient cells was calculated according to the N/P and E/C ratio of ≈ 3. N/P was defined as the negative/positive electrode areal capacity, and E/C defined as electrolyte to capacity ratio in g Ah -1 with an electrolyte density of 1.2 g mL -1 . Non-limited cells had an N/P and E/C ratio of approximately 45 (≈ 40 mg of lithium and 500 µm thick) and 17 (≈ 50 µL), respectively. Cells were crimped with a pressure-controlled electric crimping machine (MSK-160E, MTI) between 0.9-1.0 metric tons (T). Cell assembly was performed in an argon-filled glove box (VAC system 60387, NexGen 2P, Vacuum Atmospheres Company) with less than 0.5 ppm of moisture and 0.2 ppm of oxygen. Electrochemical Testing and Analysis Galvanostatic charge-discharge (GCD) cycling, rate capability tests, and electrochemical impedance spectroscopy (EIS) measurements of all coin cells were conducted using a battery test 48 station BCS-815 series potentiostat/galvanostat with EIS (BioLogic) at room temperature (23 °C). All cells were subjected to two formation cycles at C/10 charge-discharge before additional rate- capability and extended cycling tests. For the extended cycling studies, the cells were charged at a constant rate of C/5 and discharged at C/2. C-rate is based on the reversible capacity of NCA, where 1C corresponds to 160 mA g −1 . For the rate capability tests with active areal loading of ≈ 14 mg cm -2 , these cells were charged at C/5 and discharged in increments of C/5, C/2, 1C, 2C, 3C, 4C, and C/5. The charging potential cut-off for all cells was at 4.2 V, while the discharge cut-off was at 2.7 V. The EIS response was measured at open circuit voltage in the frequency range of 0.1 Hz to 100 kHz using a sinusoidal excitation amplitude of ± 5 mV (peak to peak). The impedance response was analyzed using the Zfit software (EC-Lab/BT-Lab -2018 version). The fitting was executed utilizing the non-linear complex downhill Simplex method with a previous randomized step with 5000 iterations and more than 10000 iterations for the method with a weight based on the |z|. Specific capacity (mAh g -1 ) was based on the weight of the NCA active material as the capacity contribution from the π-conjugated polymer was negligible. Specific energy at the cell- level (Wh kg -1 ) was defined as E cell = E int/W cell, where E int is the calculated energy by integrating the area under the curve of a voltage vs capacity plot for each GCD cycle, and W cell , the total weight of cell components that includes the current collector, cathode, electrolyte, separator, and anode, but excludes the cell casing. Thus, the results presented here can be extended to various types of cell configurations by appropriate inclusion of external packaging mass as per the cell size. 49 Methods for Enhancing the Ionic Conductivity of Poly(3,4- propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries (Chapter 5) Monomer and Polymer Synthesis • Acknowledgment: Polymer materials were supplied by Dr. Barry Thompson’s lab through the efforts of Pratyusha Das • Objectives: Synthesize Poly(3,4-propylenedioxythiophenes) with Oligoether Side Chains conjugated polymers. • Detailed procedures can be found in Das & Elizlade-Segovia et al. 2022. 14 Compounds A, B and 2 (Scheme 2.5) were prepared following an established literature procedure. 1 PProDOT-Hx 2 was prepared by following our previously reported procedure. 7 Synthesis of C. Compound C was prepared by modifying literature procedure. 8 Synthesis of D. Compound D was prepared by modifying literature procedure. 8 Synthesis of (Hex:OE) PProDOT Random Copolymers ((65:35), (75:25), (85:15) and (95:5)). The (Hex:OE) PProDOTs were prepared by following our previously reported literature procedure. 15 The polymers were obtained as a shiny purple solid. Synthesis of (Hex:OE) (50:50) PProDOT Alternating Copolymer. The (Hex:OE) (50:50) PProDOT was prepared by following our previously reported literature procedure. 15 Synthesis of (Hex:OE) (25:75) PProDOT Random Copolymer. The (Hex:OE) (25:75) PProDOT was prepared by DArP by modifying literature procedure. 16 50 Synthesis of (Hex:OE) (15:85) PProDOT Random Copolymer. The (Hex:OE) (15:85) PProDOT was prepared by DArP by modifying literature procedure. 16 Synthesis of PProDOT-OE 2. PProDOT-OE 2 was prepared by DArP by modifying literature procedure. 16 Scheme 2.3: Synthesis of PProDOT-Hx 2. Scheme 2.4: Synthesis of Monomers C and D. 51 Scheme 2.5 Synthesis of (Hex:OE) PProDOT Random Copolymers and PProDOT-OE 2. Molecular Characterization NMR spectra for all small molecules and monomers were recorded on either a Varian Mercury 400 MHz or Varian VNMRS-500 MHz at 25 °C using CDCl3. Polymer NMR was obtained on a Varian VNMR-500 MHz at 25 °C using CDCl3. All spectra were referenced to CHCl3 (7.26 ppm). 52 Number average molecular weight (Mn) and polydispersity index (Ð) for all polymers (except (Hex:OE) (50:50) PProDOT and PProDOT-OE2) were determined by size exclusion chromatography (SEC) using Agilent 1260 Infinity II High Temperature GPC with a Differential Refractive Index (DRI) detector with 80 °C HPLC grade 1,2,4-trichlorobenzene (o-DCB) as eluent at a flow rate of 1 mL/min. The instrument was calibrated vs. polystyrene standards (1050−3,800,000 g/mol), and data were analyzed using Agilent GPC/SEC software. Polymer samples were dissolved in HPLC grade 1,2,4-trichlorobenzene (TCB) at a concentration of 0.5 mg ml−1, stirred at 80°C until dissolved and filtered through a 0.2 μm PTFE filter. For PProDOT-OE2 and (Hex:OE) (50:50) PProDOT (depending on GPC instrument availability), GPC was measured on a Waters e2695 Separations Module system using Agilent PLgel columns. Refractive index traces were used for molecular weight determination using polystyrene calibration standards (Agilent Technologies, EasiVial) and were collected in 35°C in tetrahydrofuran with a flow rate of 0.3 mL/min. 53 Monomer NMR Figure 2.9 1 H-NMR of B in CDCl 3 at 25 °C and 500 MHz. Figure 2.10 1 H-NMR of A in CDCl 3 at 25 °C and 500 MHz. 54 Figure 2.11 1 H-NMR of C in CDCl 3 at 25 °C and 500 MHz. Figure 2.12 1 H-NMR of D in CDCl 3 at 25 °C and 500 MHz. 55 Polymer NMR Figure 2.13 1 H-NMR of PProDOT-Hx 2 in CDCl 3 at 25 °C and 500 MHz. Figure 2.14 1 H-NMR of (Hex:OE) (95:5) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 4.7 % as observed from NMR 56 Figure 2.15 1 H-NMR of (Hex:OE) (85:15) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 14.7 % as observed from NMR Figure 2.16 1 H-NMR of (Hex:OE) (75:25) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 24.6 % as observed from NMR 57 Figure 2.17 1 H-NMR of (Hex:OE) (65:35) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 32.6 % as observed from NMR Additional NMR (Polymers soluble in battery electrolyte) Figure 2.18 1 H-NMR of (Hex:OE) (50:50) PProDOT Alternating Copolymer in CDCl 3 at 25 °C and 500 MHz. 58 Figure 2.19 1 H-NMR of (Hex:OE) (25:75) PProDOT Random Copolymer in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 72.5 % as observed from NMR Figure 2.20 1 H-NMR of (Hex:OE) (15:85) PProDOT Random Copolymer in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 80 % as observed from NMR 59 Figure 2.21 1 H-NMR of PProDOT-OE 2 Random Copolymer in CDCl 3 at 25 °C and 500 MHz. Electrochemical Characterization and Cell Fabrication Interdigitated Electrode Preparation • Acknowledgment: Conductivity measurements were carried by Billal Zayat from USC. • Objectives: To characterize the electronic and ionic conductivity of the polymer as function of doping (potential). Interdigitated microelectrodes (IDM) were purchased from Metrohm Dropsens (DRP-G-IDEAU5) and were rinsed with isopropanol and dried under argon before use. Solutions of PProDOT polymers (20 mg mL -1 ) were prepared by dissolving the synthesized polymer powder in 1,2- dichlorobenzene (99%, Sigma-Aldrich). The solutions were then heated at 40°C under argon for two hours to ensure complete dissolution of the polymer in the solvent. 5 μL of the prepared solutions were then spin-coated on the IDM for 30 seconds at 1000 RPM to produce a 50 nm polymer film. The prepared electrodes were then annealed under vacuum at 110℃ for two hours then transferred to an argon glovebox. At least three IDMs were prepared for each polymer sample. Since PProDOTs are a class of conjugated polymers with high thermal stability up to 60 250°C as observed by thermogravimetric analysis (TGA), they are not expected to degrade at the annealing temperature. Electrochemical Testing All electrochemical tests were performed in an argon glovebox at room temperature using a 3- electrode cell with the IDM as the working electrode and lithium foil as the counter and reference electrodes. 1 M Bis(trifluoromethane)sulfonimide lithium salt (LiTFSI) in a mixture of 1:1 by volume of ethylene carbonate and dimethyl carbonate (EC/DMC) was used as the electrolyte. Conductivity Measurement The electronic and ionic conductivity of the PProDOT-OE polymers as a function of potential were determined as previously reported.6 To determine ionic conductivity, the electrode was held at a target potential for 300 seconds followed by a potentiostatic EIS measurement between 100 kHz and 100 mHz) with an excitation of 10 mV. To measure the electronic conductivity, the electrode was first electrochemically doped for 300 seconds using the 3-electrode cell. The cell was then allowed to reach a steady open circuit potential. Potentiostatic EIS data was then collected between the terminals of the two interdigitated gold electrodes at open circuit with an excitation of 10 mV. Electrode preparation (Hex:OE) PProDOTs and control PProDOT-Hx 2 electrodes were prepared by mixing NCA (MTI corporation), Super P carbon black (MTI), multiwalled carbon nanotubes (CNTs, Cnano, OD×ID×L: 25 nm ×10 nm × 10 μm) powders, and conductive polymer/ODCB solution (50 g L-1) in a weight ratio of 90:3:3:4. Powders were gently mixed in a mortar and pestle before being added to the 61 ODCB solution. For the 2% and 1% binder electrodes a weight ratio of 95:1.5:1.5:2 and 96:1.5:1.5:1 was employed respectively. For the low carbon test, the electrode weight ratio was 94.5:0.75:0.75:4. Control NCA-PVDF electrodes were fabricated by dissolving polyvinylidene fluoride (PVDF, MTI) in N-methyl-2-pyrrolidone (NMP, Sigma-Aldrich) and mixed with NCA, Super P carbon black, and CNTs in the same ratios. The resulting slurries were coated onto aluminum foil with the doctor blade technique. The electrodes were vacuum-dried overnight, roll pressed (12’’ vise-mount slip roll T10727, Grizzly Industrial) and then punched to 14 mm diameter discs. The approximate areal loading of the low-loading, true-discharge low-loading, high-loading, cycle-life, low-carbon, 2% binder, and 1% binder electrodes was 5.0 ± 0.8 mg cm -2 , 5.2 ± 0.8 mg cm -2 , 11.1 ± 0.5 mg cm -2 , 3.1 ± 0.4, 4.9 ± 0.5 mg cm -2 , 5.5 ± 0.3 mg cm -2 and 5.3 ± 0.6 mg cm -2 respectively. Cell assembly CR2032 coin-type cells were fabricated with NCA-conductive polymer electrodes as the working electrode, metallic lithium (MTI, D×T :16 mm × 0,6 mm) as counter, and Celgard 2325 (PP/PE/PP) as separator. Conventional carbonate electrolyte 1.0 M Lithium bis(trifluoromethanesulfonyl)imide in 1:1 in ethylene carbonate and diethyl carbonate, (LiTFSI EC/DEC =50/50 (v/v), Sigma-Aldrich) was utilized. The approximate electrolyte amount was ≈ 50 µL. Cells were crimped with a pressure controlled electric crimper machine (MSK-160E, MTI) between 0.9-1.0 metric tons (T). Cell assembly was performed in an argon-filled glove box (VAC system 60387, NexGen 2P, Vacuum Atmospheres Company) with less than 0.5 ppm of moisture and 0.2 ppm of oxygen. 62 Cell Testing Battery test station BCS-815 series potentiostat/galvanostat with electrochemical impedance spectroscopy (EIS) (BioLogic) was employed for the galvanostatic charge-discharge (GCD) cycling, rate capability tests, and EIS measurements of all coin cells. 4 hours equilibration time at open circuit voltage were imposed to all cells before testing. C-rate was defined as the reversible capacity of NCA at 1C that correspond to 160 mA g −1 .Two formation cycles at C/10 charge- discharge were carried before cycling for all cells. True discharge rate capability tests were charged at C/5 and discharged in increments of C/5, C/2, 1C, 2C, 4C, 6C, and C/5. Symmetric rate capability tests were charged-discharged in increments of C/5, C/2, 1C, 2C, 4C, 6C, and C/5. Cycle- life cells were charged-discharged at a constant rate of 1C. The charging potential cut-off for all cells was imposed at 4.2 V vs Li/Li+ while the discharge cut-off was at 2.7 V. The EIS response was measured at open circuit voltage in the frequency range of 0.1 Hz to 100 kHz using a sinusoidal excitation amplitude of ± 5 mV (peak to peak). Specific capacity (mAh g -1 ) was based on the weight of the active material (NCA) being the capacity contribution from the conductive polymer negligible. All testing was done at room temperature (21 °C). Electron Paramagnetic Resonance (EPR) • Acknowledgment: Electron paramagnetic resonance characterization was carried by Dr. Rachel Segalman’s from UCSB through the efforts of Gordon Pace. • Objectives: To confirm the nature of charge carriers in the binder at the molecular level using spin counting. 10 mg/mL of PProDOT solution in chloroform was dropcast on a PTFE substrate under nitrogen to form films c.a. 1.0 mg and 7mm diameter. The films were subsequently dried at 70° C under 63 vacuum in the nitrogen glove box and lifted off the PTFE substrate to form free standing films. The free films were further dried under high vacuum (2 × 10 –8 torr) for 12 hours. The PProDOT films were transferred to an Argon filled glove box for electrochemical doping. Each film was employed as the working electrode in swagelock cell, where Li metal served as the counter and reference electrode. The electrolyte selected for EPR measurements was 1.0 M LiPF 6 in EC:DMC (1:1 v/v). LiPF 6 was selected as the electrolyte as PF 6 – is the most common counterion used in EPR studies of polythiophene derivatives, 7,15,17–19 likely due to favorable charge carrier stability in ex situ studies. For consistency with prior EPR literature, LiPF 6 was selected. However, LiTFSI has been used for subsequent cell testing due to its better reversible redox reaction and greater ionic conductivity compared to LiPF 6, in EC/DMC electrolyte resulting in higher capacity in the battery. All electrochemical tests were performed under Argon. First, cyclic voltammetry was performed between 2.8 V and 4.2 V vs. Li/Li + . To break in the films and achieve reversible voltammograms, 3 cycles at 20 mV/s were performed. Following this was one cycle at 1 mV/s. Such low sweeping rate was selected to minimize influence of overpotential, increasing the accuracy of the charge count. Following the full (forward and reverse) scan, a linear sweep from 2.8V to the desired target potential was performed at 1 mV/s. The linear sweep was compared with the relevant portion of the full scan to ensure overlap of the two curves at the 1 mV/s scan rate, indicating reversibility. Following the linear scan to the target potential, chronoamperometry at this potential was performed to prevent dedoping before the cell could be dissembled (c.a. 10 min). The swagelock cell was then quickly disassemble and the PProDOT film was rinsed in DMC to remove residual lithium salt. Care was taken to load the free films into the EPR tube, as having 64 any conductive substrate (such as a current collector) can convolute the EPR signal, hurting signal to noise. Integration of the current in CV measurements gives the total charge (Q) via equation S1 where dV/dt is the voltage sweep rate and I ox represents oxidation current. 17 Q = ∫ I ox ( dV/dt) dV ( 𝑆 1) Following the doping process, samples were inserted into 5 mm-diameter EPR tubes and sealed with plastic caps and parafilm in the glovebox to maintain an inert sample environment during measurement. They were then transferred outside of the glovebox and EPR spectra were measured as promptly as possible using a Bruker EMXplus EPR Spectrometer. It is worth noting that the spin character of these films is quite stable over time, where Figure S18 shows EPR results for PProDOT-Hx:OE 75:25 measured the day of oxidation and after 1 week. Spin concentrations were determined by comparing the normalized, integrated signal intensity of a sample with the normalized, integrated signal intensity for a series of standards intended for organic radicals using Bruker’s internal spin counting software. Figure 2.22 (a) A representative cyclic voltammogram for (75:25) PProDOT, acquired at 1 mV s-1 in 1 M LiPF 6 in EC:DMC (1:1 v/v). (b) Polaron character of the (75:25) PProDOT measured the same day as electrochemical doping (red squares) and after 1 week (blue circles). 65 GIWAXS • Acknowledgment: grazing incidence wide-angle x-ray scattering (GIWAXS) and Ellipsometry characterization was carried by Dr. Sarah Tolbert’s lab from UCLA through the efforts of Charlene Salamat. • Objectives: To understand the structural changes induced by doping. All films were spin-coated onto silicon substrates out of 20 mg/mL solutions. Spin conditions of 2000 rpm for 55 seconds, 2500 rpm for 5 seconds was used for all films. Neutral films for 2-D GIWAXS measurements were prepared on a 1.5 × 1.5 cm silicon substrates with <100> orientation. The substrates were sonicated in an alconox/water solution, water, isopropanol, and acetone, subsequently. Electrochemically doped films were prepared on 1 cm × 2 cm silicon substrates with <100> orientation, resistivity of 0.001-0.005 Ω/cm, and 100 nm of aluminum evaporated on the back side. GIWAXS measurements were performed at the Stanford Synchrotron Radiation Lightsource on beamline 11-3 using a wavelength of 0.9742 Å with an incidence angle of 0.12°. The IgorPro macro, Nika was used to calibrate the GIWAXS 2-D data, while WAXStools was used to reduce the data. All analysis was performed on IgorPro. Each integration pattern is baseline subtracted. Swelling Studies All films were spin-coated onto 1 cm × 1 cm silicon substrates out of 20 mg/mL solutions. Spin conditions of 2000 rpm for 55 seconds, 2500 rpm for 5 seconds was used for all films. The substrates were sonicated in an alconox/water solution, water, isopropanol, and acetone, subsequently. Upon cleaning and drying, the substrates were massed. Immediately after getting the mass of the substrate, the polymer film was spun onto the substrate and it’s mass was 66 remeasured, now with the mass of the substrate and the mass of the polymer film. Directly after getting this new mass, the film was placed in a glass container with propylene carbonate (PC) and a smaller glass container inside. The film lays on the glass container inside such that it is not in contact with the propylene carbonate liquid. The glass vial is parafilmed. After swelling the film in this container, the film is once again massed. The percent mass change is calculated from the following equation: ( ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑠𝑤𝑜𝑙𝑙𝑒𝑛 𝑓𝑖𝑙𝑚 )− ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑓 𝑖𝑙𝑚 + 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 ) ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑓𝑖𝑙𝑚 + 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 )− ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 ) )× 100% For example, calculation for one trial of 65:35 polymer: Mass of substrate: 168.371 mg Mass of substrate + film: 168.411 mg Mass of substrate + swollen film: 168.450 mg Therefore, the % mass change is: ( 168.450 𝑚𝑔 − 168.411 𝑚𝑔 168.411 − 168.371 𝑚𝑔 )× 100% = 97.5% 67 Figure 2.23 Swelling set-up with glass vial and propylene carbonate inside. GPC Traces of the Polymers: Figure 2.24 GPC trace of PProDOT-Hx 2: M n = 17.4 kDa, Ð = 1.76 68 Figure 2.25 GPC trace of (95:5) PProDOT: M n = 17.6 kDa, Ð = 1.74 Figure 2.26 GPC trace of (85:15) PProDOT: M n = 13.9 kDa, Ð = 1.81 69 Figure 2.27 GPC trace of (75:25) PProDOT: M n = 16.3 kDa, Ð = 1.78 Figure 2.28 GPC trace of (65:35) PProDOT: M n = 13.7 kDa, Ð = 1.49 70 Methods for LixNaNb13O33 as an anode material for lithium-ion Batteries (Chapter 6) • Acknowledgment: Pristine active material NaNb 13O 33 was provided by Dr. Ram Seshadri’s lab by the efforts of Ashlea Patterson. Electrode preparation and cell assembly Objectives: build lithium metal cells utilizing NaNb 13O 33 as the active material. Pristine active material NaNb 13O 33, Super P carbon black (MTI), and polyvinylidene fluoride (PVDF) were mixed in a weight ratio of 80:10:10. The active material and conductive carbon were blended in a mortar and pestle and added to a solution of polyvinylidene fluoride (PVDF, MTI) in N-Methyl-2-pyrrolidone (NMP, Sigma-Aldrich) (40 g L -1 ) producing a slurry. The slurry was cast onto a copper foil using the doctor blade technique and the coating thickness was approximately 40 μm. The resulting film was vacuum-dried at 80 °C overnight and then punched into 14 mm discs with approximate active material mass loadings of 2, 2.2, and 5.5 mg cm -2 . Assembly of CR2032 coin-type cells with NaNb 13O 33 as the cathode, metallic lithium discs (MTI, D×T :16 mm × 0,6 mm) as the anode, and glass fiber grade GF/C (Whatman) as the separator was performed in an argon-filled glove box (VAC system 60387, NexGen 2P, Vacuum Atmospheres Company) with less than 0.6 ppm of moisture and 0.3 ppm of oxygen. The electrolyte was prepared by dissolving 1.0 M Lithium bis(trifluoromethanesulfonyl)imide (LiTFSI, Sigma-Aldrich) in 1:1 in ethylene carbonate (EC, Sigma-Aldrich) and dimethyl carbonate (DMC, Sigma-Aldrich). Cells were crimped with a pressure-controlled electric crimping machine (MSK-160E, MTI) between 0.8-0.9 metric tons (T). 71 Electrochemical testing Long-term cycling, testing for rate capability and galvanostatic charge/discharge (GCD) of the cells was carried in a battery test station BCS-815 series potentiostat/galvanostat equipped with impedance measurement capability (BioLogic). The cut-off voltage during discharge was 1 V for all the cells, while the charge cut-off voltage was at 3 V vs Li + /Li. Rate capability testing had three formation cycles at C/15 prior to charge-discharge at various rates of C/2, 1C, 2C, 6C, 10C, 15C, 20C, 40C, and 60C. C-rate is defined based on the reversible capacity of NaNb 13O 33 and 1C corresponded to 160 mA g −1 . Long-term cycling was carried out and at a rate of 2C following two formation cycles at C/10. Impedance was measured in the frequency range of 0.1 Hz to 100 kHz at a sinusoidal excitation of± 5 mV (peak to peak) at selected states of charge following the relaxation of the open circuit voltage to < 6 mV h -1 . Methods for Ketjen black-based Sulfur Cathodes in Li-S (Chapter 7) Electrode preparation, cell assembly and testing Sulfur (Aldrich, 99.5% purity), Super-P  (Alfa Aesar), Ketjenblack (Akzo Nobel EC 600JD and EC 300J), N-methyl-2-pyrrolidone (NMP, Sigma-Aldrich) and polyvinylidene fluoride (PVDF) binder (MTI) were used as received. The typical cathode slurry consisted of 70 wt % sulfur active material, 20 wt % Ketjenblack carbon and 10 wt % PVDF binder. The PVDF was first dissolved in NMP to which solution, sulfur and carbon were added. The mixture was stirred in a vial overnight to ensure thorough mixing and then coated onto an aluminum foil using the doctor-blade technique. In the case of the hybrid electrode, 10 wt % Super-P  was also added to the electrode slurry with the final composition being 60 wt% sulfur, 20 wt% KB, 10wt % Super-P  and 10 wt % 72 binder. The thickness of the electrode was `~25 μm. The coating was dried at 80 ᵒC overnight in vacuum. The electrode was then punched into 16 mm diameter disks and transferred to an Argon-filled glove box. The sulfur loading in the electrode was ~2 mg cm -2 . Coin cells (size 2032, MTI) were assembled using lithium foil (MTI) anode, three layers of polypropylene separator (Celgard 2500) and electrolyte consisting of 1 M lithium bis(trifluoromethanesulfonyl) imide (LiTFSI, Aldrich) in 1:1 volume ratio of dioxolane/dimethyoxethane (Aldrich). The electrolyte to sulfur ratio was 20 µL/mg-sulfur. All cells were equilibrated for 3 h prior to testing. EIS measurements were carried out at different depths of discharge (DOD) in the frequency range of 100 kHz-0.1 Hz at a sinusoidal excitation of 5 mV. Prior to this measurement, the cells were discharged (or charged) at C/50 rate in steps of 5 % DOD and allowed to relax under open circuit conditions until a steady cell voltage was reached (in typically 30 minutes). The EIS measurements were then carried out at this stable open circuit voltage. The resulting impedance and phase shift data were analyzed with ZSimpWin software using an appropriate equivalent circuit model. The surface morphology and elemental composition were studied by Transmission electron microscopy (TEM JEOL JEM-2100F equipped with EDX unit). Methods for Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature (Chapter 8) Preparation of the Composite sulfur Electrode. Objective: build a solid-state composite electrode with intercalating nano particles as electrolyte. 73 A mixture of sulfur (Aldrich, 99.5% purity) with Ketjen black (KB, Akzo Nobel EC-600JD) in the weight ratio of 72:28 was prepared by melt infusion at 155 °C for 12 hours. A slurry of the sulfur- KB mixture, lithium cobalt oxide powder (Alfa chemistry 99.9% purity, average particle size <100 nm), and polyvinylidene fluoride (PVDF) binder (MTI) in a weight ratio of 60:33:7 was prepared in N-methyl-2-pyrrolidone (NMP) solvent (anhydrous, 99.5%, Aldrich). The resulting slurry was stirred overnight and coated onto an aluminum foil with a doctor blade and partially dried in a vacuum oven at 70 ᵒC for 2 hours. The electrode was cut into disks of 2 cm 2 . Any residual solvent was removed with a blotting wipe before the film started to crack, leaving 10-15 μL of NMP before assembling, (the volume was measured by the difference in weight against a dried electrode and using the density of NMP). We did not let the solvent evaporate completely so as to avoid cracking of the electrode film and to ensure a good contact of the solid/solid interphase by adhesion between the composite electrode and the bilayer electrolyte. Control electrodes were prepared in the same manner where sulfur was melted into Ketjen Black (KB) in a weight ratio of 75:25. Then, a slurry was prepared with the sulfur-KB mixture, Super P carbon, and PVDF binder in a weight ratio of 80:10:10 dissolved in NMP. The slurry was coated using doctor blade and dried in a vacuum oven at 70 ᵒC overnight. The final sulfur loading for all the electrodes ranged from 1.2 to 1.7 mg/cm 2 . Preparation of the bilayer solid electrolyte Objective: fabricate a flexible composite bilayer solid-state electrolyte. Lithiated cobalt oxide (Aldrich, Lithium cobalt(III) oxide molecular formula LiCoO 2, 99.8% trace metals basis) and PVDF in the ratio 95:5 by weight, respectively, was made into a slurry with NMP. The slurry was coated onto an aluminum foil using a doctor blade, different thicknesses 74 were tested ranging from 30 to 80 micrometers, the coating was dried in a vacuum oven at 70 ᵒC overnight. The coated aluminum foil was then compressed in a hot press at 140 ᵒC at a pressure of 0.5 kg/cm 2 . A solution was prepared by dissolving poly(ethylene oxide) (PEO, Aldrich, MW 300,000) and bis(trifluoromethanesulfonyl) imide (LiTFSI, Aldrich) salt in a molar ratio of 8:1, in tetrahydrofuran (HPLC, ≥99.9%). The polymer solution was coated onto the pressed lithiated cobalt oxide layer using a doctor blade to a thickness that ranged between 10 to 30 μm. The resulting bilayer was dried under vacuum at 40 ᵒC overnight and punched into 2 cm 2 disks. The aluminum foil substrate was subsequently peeled off to yield a free-standing bilayer of lithiated cobalt oxide and PEO. The LCO loading was approximately 15.5 mg/cm 2 . The polymer electrolyte layer with alumina was prepared by the same procedure as above except that the solution had alumina powder (Aldrich, nanopowder, <50 nm particle size) that was 10 wt% of mass of PEO and TFSI, all dissolved in tetrahydrofuran (HPLC grade, ≥99.9%). Coin cell fabrication Coin cells of the 2032-type with aluminum-clad cans (MTI) were assembled in an argon-filled glove box (MTI, less than 2 ppm of moisture) using the composite disk electrodes as the positive electrode and lithium foil (MTI) as the counter electrode. The bilayer was interposed between the electrodes where the polymer electrolyte side of bilayer was in contact with the lithium electrode and the LCO containing MCM side faced the composite electrode. No separator or liquid electrolyte were added. 75 Control cells with liquid electrolyte were assembled using lithium foil as the negative electrode, polypropylene (Celgard 2500) as a porous separator, and the control Sulfur-KB-Super P-PVDF electrodes described in the previous section. The liquid electrolyte was 1 M lithium bis(trifluoromethanesulfonyl) imide (LiTFSI, Aldrich) dissolved in a 1:1 mixture (by volume) of dioxolane and dimethoxyethane (Aldrich). The electrolyte amount was 20 μL per mg of sulfur. Electrochemical characterization Galvanostatic charge/discharge cycling of the coin cells was conducted using a battery test station BST8-MA (MTI) and a PAR VMC-4 potentiostat at room temperature (23 °C). Cells were cycled at C/50, C/30, C/20, C/14, and C/8 rates where the 1C corresponded to 1675 mA/g. The charging was cut off was at 2.8 V vs Li/Li + while the discharge was cut off at 1.4 V. Electrochemical impedance spectroscopy (EIS) measurements were carried out using a PAR VMC-4 potentiostat/impedance analyzer. The EIS response was measured in the frequency range of 0.01 Hz to 100 kHz and a sinusoidal excitation amplitude of ±3 mV (peak to peak). The impedance response was analyzed using ZSimpWin software. Scanning electron microscopy (SEM) The bilayer was frozen in liquid nitrogen and then fractured. The cross-section was imaged using scanning electron microscopy (SEM, JEOL 661 SEM). Methods for The Role of Functionalized Conducting Polymer Binders in Addressing the Technical Challenges of Lithium-Sulfur Batteries (Chapter 9) Experimental Synthesis of N2200 and N2200-OE polymers 76 Polymer Synthesis Scheme 2.3. Synthesis of monomer S4 and polymers PNDI (OD:OE) and P(NDI2OD-T2). 77 Synthesis of S1: Compound S1 was prepared by modifying literature procedure.1 To an oven dried 250 mL three-necked round bottom flask equipped with an addition funnel, Triethyleneglycol monomethyl ether (6 g, 36.54 mmol), triphenylphosphine (9.58 g, 36.54 mmol), and phthalimide (5.23 g, 36.54 mmol) in dry diethyl ether (60 mL) were added. A solution of DIAD (7.39 g, 20.09 mmol) in dry diethyl ether (30 mL) was added slowly through the addition funnel at room temperature. The reaction mixture was stirred at room temperature for 24 hours after which the formed precipitate was filtered off and washed thoroughly with diethyl ether. After removal of the solvent in the filtrate by rotary evaporator, the crude product was purified by column chromatography (silica) using hexane/ethyl acetate (1:1) as an eluent to give a viscous and colorless oil (7.8 g, 73%). 1H-NMR (400 MHz, 25°C, CDCl3): δ = 7.87-7.79 (m, 2H), 7.73-7.67 (m, 2H), 3.92-3.86 (t, 2H), 3.76-3.70 (t, 2H), 3.67-3.62 (m, 2H), 3.61-3.54 (m, 4H), 3.49-3.43 (m, 2H) 3.34-3.31 (s, 3H). Synthesis of S2: Compound S2 was prepared by following a similar literature procedure.1 In an oven dried 250 mL three-necked round bottom flask equipped with a condenser, a solution of S1 (7.8 g, 26.59 mmol) in methanol (150 mL) was added and stirred under Nitrogen. Then hydrazine monohydrate (4.06 g, 81.1 mmol) was added slowly through a syringe dropwise. The reaction mixture was heated to reflux at 80°C for 36 hours, after which it was cooled to room temperature and methanol was evaporated. The residue was dissolved in dichloromethane and was washed twice with 10 % potassium hydroxide (KOH) solution (100 mL). The aqueous layer was extracted three times with dichloromethane. The combined organic layers were then washed with brine two times and dried over MgSO4, and the solvents were removed 78 by rotary evaporator to get a colorless oil (2.38 g, 55%) which was used as it is. 1H-NMR (400 MHz, 25°C, CDCl3): δ = 3.67-3.61 (m, 6H), 3.57-3.53 (m, 2H), 3.53-3.48 (t, 2H), 3.38-3.36 (s, 3H), 2.88-2.83 (t, 2H). Synthesis of S4: Compound S4 was prepared by following a similar literature procedure.1 In an oven dried 100 mL round bottom flask equipped with a condenser, compound S3 (1.6 g, 3.76 mmol) was added, followed by the addition of 46 mL of glacial acetic acid. The above mixture was stirred under Nitrogen. Then compound S2 (2.64 g, 16.15 mmol) was added dropwise to the above mixture. The solution was then heated to 120 °C for 2 hours. The clear red solution was then poured in ice and extracted with chloroform. The organic phase was washed with brine and dried over MgSO4 before evaporating under reduced pressure. The crude orange-red residue was purified by column chromatography on silica gel using dry loading technique (on account of poor solubility) with a mixture of methanol/ethyl acetate (1:9), followed by filtration in hot hexanes to obtain an orange powder (0.8 g, 34%). 1H NMR (400 MHz, 25°C, CDCl3): δ = 8.98 (s, 2H), 4.46 (t, 4H,), 3.86 (t, 4H), 3.71-3.69 (m, 4H), 3.62-3.57 (m, 8H), 3.48-3.46 (m, 4H), 3.32 (s, 6H). Consistent with literature report.2 PNDI (OD:OE) (N2200-OE) Synthesis by Stille Polycondensation: PNDI (OD:OE) (80:20) was prepared by modifying literature procedure.1 Monomers S5 and S6 used for Stille polymerization, compound S3 and the polymer, P(NDI2OD-T2) (Mn = 77.4 kDa, Ð = 2.22) were prepared following our previously reported literature procedure.1 In an oven dried 25 mL round bottom flask equipped with a condenser, compound S5 (80 mg, 0.0812 mmol, 0.8 eqv), compound S4 (14.5 mg, 0.0203 mmol, 0.2 eqv) and compound S6 (49.9 mg, 0.1015 mmol, 1 eqv) were dissolved in 0.05 M of anhydrous (Chlorobenzene:DMF) (9:1) (4.1 mL). The solution was 79 degassed with Nitrogen for 15 min, followed by the addition of tris(dibenzylideneacetone)dipalladium (Pd2(dba)3, 3.8 mg, 2 mol%) and tri(o-tolyl) phosphine (P(o-Tol)3, 4.9 mg, 8 mol%). The solution was then stirred under N2 in a pre-heated oil bath at 120°C for 48 hours. After cooling to room temperature, the reaction mixture was dissolved in minimum amount of CHCl3 and precipitated into a chilled 10% NH4OH/MeOH solution with rapid stirring. The polymer product was then purified via Soxhlet extraction using MeOH, hexanes, and CHCl3. The chloroform fraction was then concentrated by evaporation and precipitated into a chilled MeOH with rapid stirring and subsequently filtered. The polymer was obtained as a dark blue solid which was then dried under vacuum overnight at room temperature, affording 90 mg (95%) in approximately quantitative yield with Mn = 16.9 kDa and Ð = 1.86. 1H NMR (500 MHz, 25°C, CDCl3,) δ: 8.82 (s, 0.72 H), 7.34 (s, 1.88 H), 4.44 (broad s, 0.36 H), 4.11 (broad s, 1.46 H), 3.84 (broad s, 0.61 H), 3.71 (broad s, 0.69 H), 3.67-3.52 (m, 1.65 H), 3.32 (broad s, 1 H), 1.99 (broad s, 1.12 H), 1.48 – 1.15 (m, 49.99 H), 0.91 – 0.77 (m, 9.57 H) ppm. All reactions were performed under dry N2 in oven dried glassware, unless otherwise noted. Solvents and inorganic reagents were purchased from commercial sources through VWR and used as received unless otherwise noted. Anhydrous toluene (EMD Millipore) and anhydrous N,N-Dimethyl formamide (DMF) (Sigma-Aldrich) were purchased and used as received. Ether was dried over activated molecular sieves (3Å) prior to use. Chlorobenzene was freshly distilled from CaH2 and was stored over 4Å Sieves. 1,4,5,8-naphthalenetetracarboxylicdianhydride (NTCDA) (BeanTown Chemical), 2-octyl-1-dodecanol (Sigma-Aldrich), triethylene glycol monomethyl ether (TCI), hydrazine monohydrate (Beantown Chemical), phthalimide (Calbiochem), triphenylphosphine (Alfa Aesar), diisopropyl azodicarboxylate (DIAD) (Acros Organics), Pd2(dba)3 80 (Matrix Scientific) and P(o-Tol)3 (Strem Chemicals) were purchased and used as received. Compounds S3, S5, S6 and the polymer P(NDI2OD-T2) or N2200 (Mn = 77.4 kDa, Ð = 2.22) were prepared following our previously reported literature procedure.52 Electrochemical Characterization of Polymers A polymer solution (20 mg mL -1 ) was prepared by dissolving the polymer powder in 1,2- dichlorobenzene (99%, Sigma-Aldrich). The solution was then heated to 40 °C under argon for two hours to ensure complete dissolution of the polymer powder in the solvent. 5 μL of the prepared solution was then spin-coated onto a gold interdigitated microelectrode (IDM) on a glass substrate (Metrohm Dropsens, DRP-G-IDEAU5) at 1000 RPM for 30 seconds. The resulting film was approximately 50 nm as measured by profilometry. The coated electrode was then annealed under vacuum at 110℃ for two hours. The electrochemical testing of the polymer thin films was done in a three-electrode cell, with the interdigitated electrode as the working electrode and lithium foil as counter and reference electrodes. 1 M lithium bis- (trifluoromethanesulfonyl)imide (LiTFSI) in a 1:1 volume ratio of 1,3 dioxolane (anhydrous, Sigma-Aldrich) and 1,2 dimethoxyethane (anhydrous, Sigma-Aldrich) (DOL/DME) was used as the electrolyte. All electrochemical tests on the polymer thin films were performed using a VersaSTAT single-channel potentiostat (AMETEK Scientific Instruments) in an argon glovebox at room temperature. The electronic and ionic conductivities of N2200 and N2200-OE were measured as a function of electrode potential in 1M LiTFSI in DOL/DME using a method that we have previously reported.52 To determine ionic conductivity, the electrode was held at a target potential for 300 seconds followed by a potentiostatic EIS measurement between 100 kHz and 100 mHz with a 81 sinusoidal excitation of 10 mV. To measure the electronic conductivity, the electrode was first electrochemically doped for 300 seconds using the 3-electrode configuration then allowed to reach a steady open circuit potential. Potentiostatic EIS data was then collected between the two terminals of the interdigitated pair of gold electrodes at open circuit with an excitation of 10 mV. Physical Characterization Nuclear Magnetic Resonance (NMR) was used to characterize both the monomers and polymers. All small molecule and monomer NMR spectra were recorded at 25 °C using CDCl3 or DMSO [d6] on either a Varian Mercury 400 MHz or Varian VNMRS-500 MHz. Polymer NMR was obtained on a Varian VNMR-500 MHz at 25 °C using CDCl3. All spectra were referenced either to CHCl3 (7.26 ppm) or DMSO (2.50 ppm). Number average molecular weight (Mn) and polydispersity (Ð) of the polymers were determined by size exclusion chromatography (SEC) using an Agilent 1260 Infinity II high-temperature GPC with a differential refractive index detector with 140 °C HPLC-grade 1,2,4-trichlorobenzene (TCB) as the eluent at a flow rate of 1 mL min-1. The instrument was calibrated vs polystyrene standards (1050−3,800,000 g mol-1), and data was analyzed using Agilent GPC/SEC software. Polymer samples were dissolved in HPLC-grade 1,2,4-trichlorobenzene at a concentration of 0.5 mg mL−1, stirred at 80 °C until dissolved, and filtered through a 0.2 μm PTFE filter. Grazing Incidence Wide Angle X-ray Scattering (GIWAXS) data was collected at the Stanford Synchrotron Radiation Lightsource at beamline 11-3 in reflection geometry. A wavelength of 0.9742 Å with an incidence angle of 0.12° were used. All films were cast using 20 mg mL−1 82 solutions in 1,2-dichlorobenzene (ODCB). Spin conditions of 2000 rpm for 55 seconds, 2500 rpm for 5 seconds was used for all films. The neutral GIWAXS samples were prepared on 1.5 cm × 1.5 cm silicon substrate with <100> orientation. These substrates were sonicated in an alconox/water solution, water, isopropanol, and acetone, sequentially. The electrochemically doped samples were prepared on 1 cm × 2 cm substrates with <100> orientation, resistivity of 0.001-0.005 Ω cm−1, and 100 nm of aluminum thermally evaporated on the back side. IgorPro 8 along with the macros Nika are used to calibrate the GIWAXS 2-D data. WAXStools was used to reduce the data.53 All analysis was performed on IgorPro. Each integrated pattern shown has the baseline subtracted, using a Baseline macro. The swelling properties of the polymers were determined for DOL/DME. All films were cast using 20 mg mL−1 solutions in 1,2-dichlorobenzene and prepared on 1 × 1 cm silicon substrate. Spin conditions of 2000 rpm for 55 seconds, 2500 rpm for 5 seconds was used for all films. Upon cleaning and drying (same procedure as for GIWAXS samples), the mass of the substrates was measured on a microbalance. Thereafter, the polymer was spin-coated onto the substrate and then weighed again, now with polymer film. The sample was then placed in a sealed swelling chamber with 1,3-dioxolane/dimethoxyethane (DOL/DME) (1:1 ratio). The films sat on a container above the solvent level, such that the films did not touch the solvent. After vapor swelling of the films, the films were weighed again. The percent of mass increase was calculated by Eq. 1. ( ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑠𝑤𝑜𝑙𝑙𝑒𝑛 𝑓𝑖𝑙𝑚 𝑜𝑛 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 )− ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑓𝑖𝑙𝑚 + 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 ) ( 𝑚𝑎𝑠𝑠 𝑜𝑓 𝑓𝑖𝑙𝑚 + 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 )− ( 𝑚𝑎𝑠 𝑠 𝑜𝑓 𝑠𝑢𝑏𝑠𝑡𝑟𝑎𝑡𝑒 ) ) × 100% ( 1) 83 To examine the interaction between the polymers and polysulfides, UV-Vis spectra of polysulfide solutions exposed to PVDF, N2200, and N2200-OE were collected over time. Polymer electrodes were prepared by drop-casting a 20 mg mL−1 polymer solution on aluminum foil and then drying under vacuum at 110 ℃ for two hours. The electrodes were cycled in 1M LiTFSI in DOL/DME at 5 mV/s between 3 V and 1.7 V vs Li+/Li for five cycles. The electrodes were then placed in a 1 mM lithium polysulfide solution (LiPS) (average formula Li2S6) prepared by mixing Li2S powder and sulfur powder in DOL/DME solution and stirring under argon for two days at 60 ℃. UV-Vis absorption spectra of the LiPS solution were collected every 60 minutes for three hours using an Agilent Technologies Cary80 spectrometer. Full Cell Assembly, Shuttle Current Measurement, Impedance Spectroscopy, and Electrical Performance Testing Sulfur electrodes based on N2200 and N2200-OE were prepared by mixing elemental sulfur (Aldrich, 99.5% purity), Ketjen Black (KB, Akzo Nobel EC-600JD), Super P (MTI), and the n-dopable conductive polymer in a weight ratio of 70:15:10:5. First, sulfur was melt infused into KB at 155 °C for 12 hours. Then, Super P was added, and the powders were gently mixed in a mortar and pestle. The mixture was added to a solution of the conductive polymer in ODCB solution (40 g L−1) to form a slurry. The resulting slurry was coated onto an aluminum foil utilizing the doctor blade technique, vacuum-dried at 70 °C overnight, and then cut into disks (14 mm diameter). Control electrodes with PVDF binder were prepared by dissolving PVDF (MTI) in N-methyl-2- pyrrolidone (Sigma-Aldrich) (0.04 g mL−1) and cast into a slurry prepared with the same weight ratio of sulfur, carbon and binder as previously described. The approximate areal loading of sulfur 84 in the electrodes with “moderate” and “high” loadings was 1.8 ± 0.3 and 3.4 ± 0.3 mg cm−2, respectively. 2032-type coin-cells were assembled in an argon-filled glove box (VAC system 60387, NexGen 2P, Vacuum Atmospheres Company) with less than 0.5 ppm of moisture and 0.2 ppm of oxygen. Metallic lithium (MTI, D × T:16 mm × 0.6 mm) was utilized as the counter electrode and Celgard 2325 (PP/PE/PP) was used as the separator. 1 M LiTFSI in DOL/DME was the electrolyte, and the electrolyte to sulfur ratio was 12 µL mg−1. The electrolyte did not contain any other additive. Cells were crimped with a pressure-controlled electric crimping machine (MSK-160E, MTI) between 0.9 and 1.0 metric tons. Galvanostatic charge/ discharge (GCD) cycling, rate capability testing, polysulfide shuttle current, and EIS measurements were carried in a battery test station BCS-815 series potentiostat/galvanostat with EIS (BioLogic) at room temperature (21 °C). An equilibration time of 24 hours at OCV was imposed on all the cells before cycling. Two “formation” cycles were carried out at C/15. The C-rate definition is based on the theoretical capacity of sulfur where 1C = 1675 mA g−1. The cell voltage cut-off for charging and discharging were 2.8 V and 1.55 V, respectively. Shuttle current measurements were performed by first discharging the cell at C/20 for 1 hour and then allowing it to rest at open circuit voltage (OCV) with a cutoff at a voltage rate of change of 0.1 mV h−1. Then, the cell was held at the measured OCV and the current passed through the cell was observed for 1 hour during which the current reached a steady-state value. This steady-state current was recorded as the polysulfide shuttle current and was measured throughout the voltage range of cell operation. The EIS response of the full cell was measured at OCV in the frequency range of 0.1 Hz to 100 kHz using a sinusoidal excitation amplitude of ±4 mV 85 (peak to peak). True-discharge rate capability testing, for the moderate loading electrodes, was done by charging the cells at 0.3 mA cm−2 (C/10) and discharging them in increments of 0.2, 0.3, 0.6, 1.5, 3, and 6 mA cm−2 that corresponded to C/15, C/10, C/5, C/2, 1C and 2C, respectively. High loading electrodes were charged at 0.57 mA cm−2 (C/10) and discharged in increments of 0.29, 0.57, 1.14, 2.85, 5.7, and 11.4 mA cm−2 that corresponded to C/15, C/10, C/5, C/2, 1C and 2C respectively. Cycle-life characterization was carried at a constant charge/discharge rate of C/2 for 500 cycles. Specific capacity was calculated based on the weight of sulfur since the capacity contribution from the conductive polymer was negligible. Monomer and Polymer NMR Figure 2.29. 1 H-NMR of S1 in CDCl 3 at 25 °C and 400 MHz. 86 Figure 2.30. 1 H-NMR of S2 in CDCl 3 at 25 °C and 400 MHz. Figure 2.31. 1 H-NMR of S4 in CDCl 3 at 25 °C and 400 MHz. 87 Figure 2.32. 1 H-NMR of PNDI (OD:OE) or N2200-OE in CDCl 3 at 25 °C and 400 MHz. Oligoether (OE) percentage = 20 % (Feed Ratio) Oligoether (OE) percentage = 17.3 %  17 % as observed from NMR References 1. Das, P. et al. Dihexyl-Substituted Poly(3,4-Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. 32, 9176–9189 (2020). 2. Reeves, B. D. et al. Spray Coatable Electrochromic Dioxythiophene Polymers with High Coloration Efficiencies. Macromolecules 37, 7559–7569 (2004). 3. Lai, C.-H. et al. Application of Poly(3-hexylthiophene-2,5-diyl) as a Protective Coating for High Rate Cathode Materials. Chem. Mater. 30, 2589–2599 (2018). 4. Enengl, C. et al. Doping-Induced Absorption Bands in P3HT: Polarons and Bipolarons. ChemPhysChem 17, 3836–3844 (2016). 5. Elizalde-Segovia, R. et al. Understanding the Role of π-Conjugated Polymers as Binders in Enabling Designs for High-Energy/High-Rate Lithium Metal Batteries. J. Electrochem. Soc. 168, 110541 (2021). 6. Gobalasingham, N. S., Pankow, R. M., Ekiz, S. & Thompson, B. C. Evaluating structure–function relationships toward three-component conjugated polymers via direct arylation polymerization (DArP) for Stille-convergent solar cell performance. J. Mater. Chem. A 5, 14101–14113 (2017). 88 7. Reeves, B. D. et al. Spray Coatable Electrochromic Dioxythiophene Polymers with High Coloration Efficiencies. Macromolecules 37, 7559–7569 (2004). 8. Mazaheripour, A., Thomas, E. M., Segalman, R. A. & Chabinyc, M. L. Nonaggregating Doped Polymers Based on Poly(3,4-Propylenedioxythiophene). Macromolecules 52, 2203–2213 (2019). 9. Zayat, B., Das, P., Thompson, B. C. & Narayan, S. R. In Situ Measurement of Ionic and Electronic Conductivities of Conductive Polymers as a Function of Electrochemical Doping in Battery Electrolytes. J. Phys. Chem. C acs.jpcc.0c08934 (2021) doi:10.1021/acs.jpcc.0c08934. 10. Musumeci, A. W., Silva, G. G., Liu, J.-W., Martens, W. N. & Waclawik, E. R. Structure and conductivity of multi-walled carbon nanotube/poly(3-hexylthiophene) composite films. Polymer 48, 1667–1678 (2007). 11. Alexander, M. et al. Electronic structure studies on the n -type doped superconductors R 2 − x M x CuO 4 − δ ( R =Pr,Nd,Sm; M =Ce,Th) and Nd 2 CuO 4 − x F x by electron-energy-loss spectroscopy. Phys. Rev. B 43, 333–343 (1991). 12. Carmona, F. & El Amarti, A. Anisotropic electrical conductivity in heterogeneous solids with cylindrical conducting inclusions. Phys. Rev. B 35, 3284–3290 (1987). 13. Berni, A., Mennig, M. & Schmidt, H. Doctor Blade. in Sol-Gel Technologies for Glass Producers and Users 89–92 (Springer, 2014). 14. Das, P. et al. Enhancing the Ionic Conductivity of Poly(3,4-propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries. Chemistry of Materials XX, (2022). 15. Das, P. et al. Dihexyl-Substituted Poly(3,4-Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. 32, 9176–9189 (2020). 16. Rudenko, A. E. & Thompson, B. C. Optimization of direct arylation polymerization (DArP) through the identification and control of defects in polymer structure. J. Polym. Sci. Part A: Polym. Chem. 53, 135– 147 (2015). 17. Enengl, C. et al. Doping-Induced Absorption Bands in P3HT: Polarons and Bipolarons. ChemPhysChem 17, 3836–3844 (2016). 18. Lin, C., Endo, T., Takase, M., Iyoda, M. & Nishinaga, T. Structural, Optical, and Electronic Properties of a Series of 3,4-Propylenedioxythiophene Oligomers in Neutral and Various Oxidation States. J. Am. Chem. Soc. 133, 11339–11350 (2011). 19. Cházaro-Ruiz, L. F., Kellenberger, A. & Dunsch, L. In Situ ESR/UV−vis−NIR and ATR-FTIR Spectroelectrochemical Studies on the p-Doping of Copolymers of 3-Methylthiophene and 3- Hexylthiophene. J. Phys. Chem. B 113, 2310–2316 (2009). 89 Chapter 3 - High-Rate Lithium-Ion Cell with Dihexyl-Substituted Poly (3, 4-Propylenedioxythiophene) Mixed Conductive Polymer As Electrode Binder 1 Abstract The polymer binders used in most lithium ion batteries serve only a structural role, but there are exciting opportunities to increase performance by using polymers with combined electronic and ionic conductivity. Here we examine dihexyl-substituted poly(3,4-propylenedioxythiophene) (PProDOT-Hx 2) as an electrochemically stable π-conjugated polymer that becomes electrically conductive (up to 0.1 S cm -1 ) upon electrochemical doping in the potential range of 3.2 to 4.5 V (vs. Li/Li + ). Because this family of polymer is easy to functionalize, can be effectively fabricated into electrodes, and shows mixed electronic and ionic conductivity, PProDOT-Hx 2 shows promise for replacing the insulating polyvinylidene fluoride (PVDF) commonly used in commercial lithium- ion batteries. A combined experimental and theoretical study is presented here to establish the fundamental mixed ionic and electronic conductivity of PProDOT-Hx 2. Finally, the performance of PProDOT-Hx 2 as a conductive binder for the well-known cathode LiNi 0.8Co 0.15Al 0.05O 2 (NCA) relative to PVDF is presented. PProDOT-Hx 2 based cells displays a five-fold increase in capacity at high rates of discharge compared to PVDF based electrodes at high rates, and also show improved long-term cycling stability. The increased rate capability and cycling stability demonstrate the benefits of using binders such as PProDOT-Hx 2, which show good electronic and ionic conductivity, combined with electrochemical stability over the potential range for standard cathode operation. 90 Introduction Lithium-ion batteries (LIBs) are the dominant energy storage technology used today in electronics and electric vehicles because of the need for high specific energy and power density. 2–4 Traditional polymeric binders, such as polyvinylidene fluoride (PVDF), are used in these energy storage devices to hold the active electrode material together with conductive carbon additives, and to bind both materials to the metallic current collector. PVDF and related binders are chosen for their electrochemical stability, but they suffer from low electronic and Li + ion conductivities, thereby increasing electrode impedance by interfering with charge transport pathways. 5–7 Conjugated polymers that combine the advantages of organic conductors and traditional polymers, serve as excellent candidates for multi-functional binders, offering mixed electronic and ionic conduction along with mechanical integrity, thereby improving electrochemical performance. 8–18 Besides the promise of providing robust mechanical, electrochemical and thermal stability along with good wetting/coating properties, 14,19,20 it is their properties of mixed conduction that lets them function both as a binder and as a conductive additive. Because of this, conjugated polymers have been explored for use as both cathode and anode binders in LIBs. 16,20– 28 Neutral (un-doped) conjugated polymers usually exhibit low electronic conductivity in the range of 10 -10 to 10 -5 S cm -1 , but these values can be increased significantly by electrochemical or chemical doping. 29–32 While high electrical conductivity in conjugated polymers is based on achieving both sufficient levels of doping and sufficient levels of crystallinity, ionic conductivity is dictated by solvated ion motion through the polymer matrix and favored by a lower degree of crystallinity and higher porosity. 33,34 While many doped conjugated polymers show excellent 91 electronic conductivity, in order to augment performance, an important focus is on improving the ionic conductivity, which is often significantly lower. 18 A broader understanding of the critical structure-function relationships in conductive polymer binders is desired in order to leverage their use to increase rate capability, battery cycle life, 21,35–37 , and specific capacity. 20,25,35,38 Regioregular poly(3-alkylthiophenes) (P3ATs) are a broadly-studied family of conjugated polymers possessing relatively narrow band gaps, good processability, and high hole mobility. 38,39 Their solubility in many organic solvents enables formation of uniform thin film coatings on electrodes as opposed to the colloidal nature of doped polymer suspensions like poly(3,4- ethylenedioxythiophene)/poly(4-styrenesulfonic acid) (PEDOT:PSS). Upon doping, poly(3- hexylthiophene) (P3HT) exhibits enhanced electronic conductivity (up to 15 S cm -1 ) within the potential window of a majority of LIB cathodes. 19 Recent work from our team on the use of P3HT as a conductive binder for LiNi 0.8Co 0.15Al 0.05O 2 (NCA) cathodes showed high power density and excellent cycling stability compared to PVDF-NCA electrodes. 40 While P3HT enhances battery performance relative to PVDF, it still lacks Li + solubilizing functionality, limiting ionic conductivity. The semi-crystalline morphology of P3HT poses a further hindrance to Li + transport due to the dense chain packing. While the hexyl side chains of P3HT are not conducive to the diffusion of Li + , other side chains such as oligoethylene-glycol may be more beneficial for promoting Li + hopping as explored in recent work. 41,42 Reichmanis and co-workers 37,43 have also investigated blends of polythiophenes with polyethylene glycol (PEG) as a multi-component binder for anodes in LIBs in order to target mixed conductivity. Unfortunately, polythiophenes also have another drawback as binders for LIB electrodes. Namely, they do not possess the necessary chemical stability required for repeated cycling at 92 high potentials as cathode binders or low potentials as anode binders. 44,45 An alternative to P3ATs and a soluble alternative to PEDOT-PSS is poly(3,4-propylenedioxythiophenes) (PProDOTs), which are known to exhibit remarkable stability to repeated cycling over a broad potential window that is well suited for cathode materials commonly used in LIBs. 46 By functionalizing the 3- and 4- positions of thiophene with an alkylenedioxy bridge, the resulting polymers show a significantly lower onset for oxidation due to the electron donating oxygens. 47 Further, the propylene bridge in ProDOTs allows symmetrical disubstitution and hence good polymer solubility in organic solvents. 26 PProDOTs first garnered interest due to their excellent electrochromic properties, reversible electrochemical doping and fast switching times. 48 In fact, PProDOTs have excellent electrochemical stability when p-doped (oxidized) as well as faster counter-ion influx/efflux upon doping/dedoping than polythiophenes due to their open morphology, making them promising candidates for batteries and capacitors. 49–51 . Reynolds and coworkers have recently reported a ProDOT-based copolymer displaying redox stability over thousands of cycles along with reasonable charge-storage capacity. 52 As a general class, dioxythiophenes have been established for long-term cycling stability when used as supercapacitors 53,54 demonstrating capacity retention over more than 400,000 cycles. 50 As the most well-known dioxythiophene, PEDOT has shown excellent stability when used as a protective coating for LiFePO 4 and other cathodes. 55–57 resulting in enhancement on capacity and charge/discharge behavior. Such conductive coatings are also known to lead to an overall increase in electrical conductivity. In an attempt to provide mixed conductivity in three dimensional electrodes, Hammond and coworkers investigated PEDOT:PSS blended with PEO. 58 93 The favorable interaction of PEO with the PSS matrix suppressed crystallinity and elevated the ionic conductivity, resulting in mixed conductivity at an optimum PEO:PEDOT ratio. However, PEDOT:PSS is intrinsically insoluble, and the colloidal suspension does not provide the same uniform thin film coating abilities offered by soluble polymers. In order to address the current shortcomings of polymer binders for cathode in LIBs and to explore the utility of soluble dioxythiophene polymers, the focus of the current work is the investigation of the mixed electronic and ionic conductivity of dihexyl-substituted poly(3,4- propylenedioxythiophene) (PProDOT-Hx 2). While the propylenedioxythiophene backbone accounts for enhanced electrochemical stability in comparison to polythiophenes, the hexyl side chains enhance the solubility of the polymer in non-polar processing solvents improving the processability of the electrode. An open, partly disordered morphology compared to highly crystalline polymers like P3HT, allows efficient ion insertion/de-insertion, facilitating ionic conductivity. Although the electrochemical and optical characterization of PProDOT-Hx 2 has been reported, 59,60 its properties relevant to cathode binders for LIBs has not been investigated. Therefore, here we present a combined theoretical and experimental study to establish the fundamental mixed ionic and electronic conductivity of PProDOT-Hx 2, and the comparative performance of PProDOT-Hx 2 as a conductive binder for the NCA cathode relative to PVDF. Electrochemical Properties of PProDOT-Hx2 To be an effective conductive binder, it is necessary that PProDOT-Hx 2 be conducting over the potential window of operation for the battery cathode material. Also, the electrochemical doping should have fast kinetics to allow for rapid doping and dedoping during cycling. Thus, the 94 electrochemical doping (oxidation) behavior of the PProDOT-Hx 2 was first investigated in thin film format in a three-electrode cell with 1 M LiTFSI dissolved in 1:1 (by volume) ethylene carbonate (EC)/dimethyl carbonate (DMC) as the electrolyte and lithium foil as the counter and reference electrodes. The cyclic voltammetry (CV) curves between 3 and 4.2 V (vs Li/Li + ) at a scan rate of 10 mV s -1 indicated that the electrochemical doping process was highly reversible (Figure 3.1a). In the first cycle, a shoulder is observed at ~3.3 V which is known to correspond with polaron (radical cation) formation and this is supported by our Electron Paramagnetic Resonance (EPR) studies (vide infra). The first major oxidation peak (oxidation 1) at 3.36 V corresponds to bipolaron formation and is similarly established and supported. 26,48,60–64 Figure 3.1. (a) CV data for PProDOT-Hx 2 thin film in the potential range of 3-4.2 V at a scan rate of 10 mV s -1 for cycles 1-4. (b) CV data for PProDOT-Hx 2 thin film at various potential intervals, at 95 a scan rate of 10 mV s -1 . (c) CV data for PProDOT-Hx 2 thin film as a function of various scan rates from 10 to 100 mV s -1 . (d) Log of the peak current (i) vs. log of the scan rate (v) for the data shown in Figure 3.1c. Upon cycling to higher potentials, a second major oxidation (oxidation 2) is observed at ~3.8 V. This higher oxidation process is not generally reported in dioxythiophene polymers, which are rarely cycled to this high a potential. However, model studies in solution on well-defined ProDOT- Hx 2 oligomers (n = 4-12) indicate that this higher oxidation is due to the generation of additional charge carriers, where the higher doping levels are enabled by the stabilizing effect of the oxygen atoms along the backbone. 65 While the exact nature of the higher oxidation event is not fully established, the oligomer work suggests a more significant role of cation character on the oxygens in the ether bridge in the reversibly generated charge carrier species. Upon subsequent scans, the polaron pre-peak is no longer observed at this scan rate and two pairs of reversible redox peaks were observed at 3.22 to 3.16 V and 3.76 to 3.68 V (vs. Li/Li + ) with the anodic peak listed first in each pair. The reversibility of both oxidation peaks is key to the use of these polymers as binders for cathode materials. The shift of the oxidation peaks to lower potentials between the first and following cycles is attributed to the enhanced conductivity of the polymer after electrochemical doping, which reduced the overpotential. Furthermore, we have calculated the coulombic efficiency for the first cycle. Based on integrating the CV curves (Figure 3.1a), the initial coulombic efficiency (ICE) of PProDOT-Hx2 is 38.4%. The low value of ICE is likely due to side reactions such as the oxidation of the electrolyte and/or the formation of a SEI on the film. The contributions from side reactions will be significant here because the films are only about 50 nm thick. 96 For use as binders in LIBs, stability over a broad range of potentials is key, and so we further expanded the potential window to higher potential to investigate the stable operational range of the polymer. CV curves were obtained with gradually increasing potential ranges (Figure 3.1b). The shapes of the redox peaks were retained up to a voltage cutoff of 4.5 V. At potentials beyond 4.5 V, however, the redox peak features began to change, indicating decomposition of the electrolyte and polymer. The stable and wide operation potential range (3.2 – 4.5 V) of PProDOT- Hx 2 makes it an ideal candidate for use as a binder for a number of cathode materials, including LiCoO 2, LiMn 2O 4 and LiMO 2 (M=Ni, Co, Mn, Al). To ensure that electrochemical doping/de-doping of the PProDOT-Hx 2 binder does not limit cathode cycling, we also examined the kinetics of the electrochemical doping at high rates. Kinetics were analyzed based on a series of CV measurements at various scan rates (Figure 3.1c) from 10 to 100 mV s −1 . The minimal shifts of the redox peaks with increasing scan rates indicated rapid reaction rates. To further quantify the kinetics of polymer doping, we examined the relation between the measured current (i) and scan rate (v). According to Equation 1. 66 𝑖 = 𝑎 𝑣 𝑏 , ( 1) Where b can be determined by the slope of a plot of log (i) vs. log (v) plots for each redox peak. A value of b equal to 0.5 indicates a process controlled by semi-infinite diffusion, while a b value close to 1 indicates a non-diffusion controlled or a surface-controlled charge-storage process. From the plot of log (i) vs. log (v) (Figure 3.1d), the b values of the oxidation 1 anodic and cathodic peaks are 0.96 and 0.89 respectively. For the oxidation 2 anodic and cathodic peaks the b values were 0.88 and 0.94, respectively. These b values indicate rapid redox processes in the polymer 97 thin films. The fast kinetics for the electrochemical p-doping of PProDOT-Hx 2 are expected to facilitate rapid electron transport when used as a conductive binder in cathode composites. To verify the redox stability of the polymer, we tested 100 cycles and even after the long cycling the redox peaks were still evident verifying the polymer stability (Figure 3.2). CV data like that shown in Figure 3.2 also allows us to calculate the specific capacity of the polymer which is 18.5 mAh g - 1 at 10 mV s -1 . While this capacity is not important for battery performance, as binders are used at just a few percent level, it does allow us to calculate the doping level, which is important for electrical conductivity. In this case, the value corresponds to one dopant per 4.5 monomer units Figure 3.2 shows the CV curves of PProDOT-Hx 2 thin film in 3-4.2 V at 50 mV s -1 for 100 cycles. The redox peaks are still very evident, thus verifying the redox stability of the polymer. However, the capacity does decrease on cycling; the capacity retention is 60% after 100 cycles. 98 Electron Paramagnetic Resonance (EPR) EPR measurements were used to confirm the nature of charge carriers in the binder at the molecular level using spin counting. A singly charged polaron is associated with an unpaired electron and has a spin of ½, however, as the charge density increases, polarons can pair to form spinless singlet bipolarons with a charge of 2. 67–70 Here, EPR on the oxidized polymer gives insight into the structure of the charge carriers as a function of the oxidation state by determining the number of unpaired spins associated with carriers. Comparing the spin concentration in the film measured via EPR (Firgure 2b) with the charge concentration measured via CV (Figure 3.3a) gives a direct indication of the relative fraction of polarons in the material (i.e. number of polarons / number of charges). 71 The relative fraction of polarons was measured as a function of the electrode potential of the PProDOT-Hx 2 film (Figure 3.3b). Because the equilibrium open-circuit potential of the undoped PProDOT-Hx 2 was 2.95 V vs. Li/Li + , potentials above this value corresponded to electrochemical oxidation. In Figure 3.3a, the first peak of the doubet in the oxidation 1 peak is clearly observed between 3.1 – 3.2 V due to the slow scan rate of 1 mV s -1 (this is in contrast to the faster cycles in Figure 3.1a). At 3.12 V (vs. Li/Li + ), the first data point in Figure 3.3b, the ratio of spins to charges approached a value of 60%, indicating that polarons are the dominant charge carriers. This is consistent with spectroelectrochemistry reported for ProDOT polymers. 26,48,60 As the electrode potential was increased through the rest of the oxidation 1 peak, the ratio of spins to charges decayed, indicating a shift towards spinless, bipolaronic charge carriers at higher doping levels, an observations that is consistent with previous literature reports on PProDOT-Hx 2. 72 Similar 99 behavior has been found in other polymer systems, 70,71 and is attributed to increased stability of spinless carriers at high charge concentrations due to charge screening. Figure 3.3. (a) An exemplar cyclic voltammogram (CV) of PProDOT-Hx 2 thin film electrode in 1 M LiPF 6 in EC:DMC (50:50 v/v). Total number of charges of each sample were determined by CV of each sample. (b) Ratio of the number of spins to the number of charges in the PProDOT-Hx 2 film as a function of the electrochemical potential determined from the normalized, integrated EPR signal intensity. Beyond 3.4 V (vs. Li/Li + ), the relative fraction of polarons is not observed to change further. Solution studies on well-defined ProDOT-Hx 2 oligomers suggest that the reaction associated with oxidation 2 can be described by two closely spaced redox processes assigned to polaron and bipolaron formation. 65 As the oligomer length increases, however, these two peaks are observed to coalesce as the disproportionation of polarons to bipolarons becomes favorable at high doping levels, effectively rendering oxidation 2 a 2 electron process resulting in the formation of additional bipolarons. This picture is thus fully consistent with the lack of an EPR signal in this range. Overall, EPR measurements of the electrochemical doping of PProDOT-Hx 2 are consistent 100 with polaronic charge carriers dominating at low doping levels (< 3.2 V vs. Li/Li + ), while spinless bipolarons dominate at high doping levels (> 3.6 V vs. Li/Li + ). Electronic and Ionic Conductivity of PProDOT-Hx 2 Thin films of pristine PProDOT-Hx 2 are weak semiconductors with very low electrical conductivity, similar to most undoped conjugated polymers, but show increasing electronic conductivity upon electrochemical doping. 73,74 As the doping process involves the movement of counter anions into the polymer structure, pathways for ionic transport are generated. Such pathways should allow for the transport of both anions and cations through the polymer film. The anions in the polymer structure are expected to change the morphology of the polymer by electrostatic repulsion and thus modulate the electronic and ionic conductivity of the polymer. Finally, favorable interactions between counterions and the solvent should result in swelling of the amorphous regions of the polymer; the solvation of ions by both functional groups in PProDOT-Hx 2 solvent in swollen films is of immense interest for battery application. Thus, we were motivated to develop an in-situ technique to monitor the change in both electronic and ionic conductivity of PProDOT-Hx 2 as a function of electrochemical doping under conditions resembling that of a LIB. Although the ionic and electronic conductivity of various conductive polymer thin-films have been reported in the literature, this is the first report of in-situ measurement of both ionic and electronic conductivity of a polymer thin-film as a function of electrochemical doping in LIB conditions. For example, the electronic conductivity is often measured using a 4-probe method in which the neutral polymer is chemically doped by treatment with a strong oxidizaing agent. 75– 77 Although the electronic conductivity of a conductive block co-polymer has been reported as a function of electrochemical doping, it was performed on a thick polymer film in a solid state 101 system and without the elucidation of ionic conductivity. 78 On the other hand, ionic and electronic conductivity of polymer thin-films in LIB conditions have been reported but, not as a function of chemical or electrochemical doping. 58 In this work, a 50 nm thin-film of PProDOT-Hx 2 formed on an interdigitated gold electrode (Figure 3.4a) was placed in a three-electrode electrochemical cell under conditions similar to those used in standard LIBs (Figure 3.4d). This configuration allows for rapid electrochemical doping and de- doping, as seen from the cyclic voltammograms (Figure 3.4c). To determine the ionic and electronic conductivity as a function of electrochemical doping, the polymer was held at different potentials to achieve various levels of doping. Then, the electrochemical impedance of the PProDOT-Hx 2 film was measured over a sinusoidal frequency range of 0.01 Hz to 100 KHz. The impedance data was then analyzed to determine the conductivity. Since the values of electronic and ionic conductivity sometimes differ by as much as seven orders of magnitude, different electrode geometries were used to determine each conductivity value, as described below. 102 Figure 3.4. (a) SEM image of the spin-coated PProDOT-Hx 2 on interdigitated gold microelectrodes. The light color shows the gold electrodes and the dark color shows the polymer imbedded between the electrodes. (b) The 2-electrode configuration used to measure the electronic conductivity of the polymer as a function of voltage. (c) Cyclic voltammetry plot of PProDOT-Hx 2 at various scan rates. (d) The 3-electrode configuration used to electrochemically dope the polymer and determine its ionic conductivity. Electronic Conductivity as a Function of Degree of Doping Upon electrochemical doping of PProDOT-Hx 2 in the 3-electrode configuration (Figure 3.4d) the electronic resistance of the polymer thin film was determined by analysis of the electrochemical impedance in the 2-electrode configuration (Figure 3.4b). We used the Huggins approach 79 and fit the impedance data to the Debye equivalent circuit model for mixed conduction through the film (Figure 3.5a). 80 The Nyquist plot from the impedance measurements showed two semicircles. The diameter of the semicircle observed at the higher frequencies corresponded to a combination of the ionic and electronic resistances, while the second semicircle at lower frequencies was associated with the electronic resistance of the conducting polymer film. The 103 electronic resistance, R e, was then determined from the intercept on the real axis of the Nyquist plot at near zero frequency. The electronic conductivity was then calculated by accounting for the electrode geometry (equation S3), and the trend in conductivity relative to potential is shown in Figure 3.5c. For comparison, the electronic conductivity of both PProDOT-Hx 2 and P3HT in their neutral, un- doped states were approximately 10 -5 S cm -1 (Figure 3.5c). The electronic conductivity of both polymers increased by several orders of magnitude with doping, reaching a maximum of 1x10 -1 S cm -1 at around 3.6 V corresponding to complete conversion of polarons to bipolarons. The electronic conductivity then slightly decreased for PProDOT-Hx 2 to 5 × 10 -2 S cm -1 above 3.6 V. The slight decrease could be due to potential changes in the nature of the charge carriers associated with oxidation 2, 65 or to observed changes in morphology at higher doping levels, which are discussed below as explored by grazing incidence wide angle X-ray scattering GIWAXS. The electronic conductivity of P3HT reached a maximum of 1 × 10 -1 S cm -1 at 3.55 V and remained stable above 3.55 V with a slight decrease in electronic conductivity above 3.9 V. 104 Figure 3.5. (a) A representative Nyquist impedance plot obtained at 4V vs Li/Li + using the 2- electrode configuration and the fitting circuit (inset) used to obtain the electronic conductivity. (b) A representative Nyquist impedance plot obtained at 4V vs Li/Li + using the 3-electrode configuration to obtain the ionic conductivity. (c) The electronic and (d) ionic conductivities of P3HT and PProDOT-Hx 2. Ionic conductivity For determining ionic conductivity, the electrochemical impedance of the PProDOT-Hx 2 thin-film was measured in the three-electrode configuration following doping to different levels. The impedance data (Figure 3.5b) was analyzed using a transmission line equivalent circuit model used for describing porous electrode structures. 81 The ionic resistance of the polymer phase was separated from the ionic resistance of the solution and the electronic resistance of the polymer 105 film using equation 2. The value of Z real was obtained by extrapolation of the low-frequency tail to meet the real axis, the solution resistance, R s, was obtained from the high frequency intercept, and the electronic resistance R e was obtained from the measurement described in the previous section. 𝑅 𝑖𝑜𝑛𝑖𝑐 = 3 × ( 𝑍 𝑟𝑒𝑎𝑙 − 𝑅 𝑠 )− 𝑅 𝑒 ( 2) The ionic conductivity of the polymer film was then calculated from the measured ionic resistance based on the dimensions of the electrode and the film thickness (equation S2). The ionic conductivity of P3HT is 10 -9 S cm-1 up to 3.2V (vs Li/Li + ) and increased by an order of magnitude to 10 -8 S cm -1 at potentials above 3.2V vs Li/Li + (Figure 3.5d). On the other hand, PProDOT-Hx 2, upon doping showed an ionic conductivity of 10 -7 S/cm which is at least an order of magnitude higher than that of P3HT. Despite a slight decrease beyond 3.2 V, the ionic conductivity of PProDOT-Hx 2 remained relatively constant with increasing levels of electrochemical doping. Changes in morphology upon doping as indicated by the results of GIWAXS studies (vide infra) may be connected to the improved ionic conductivity observed with doping. More directly, as indicated by our theoretical study, the higher ionic conductivity of PProDOT-Hx 2 can be attributed to the Li + solvation abilities of the oxygens present in PProDOT- Hx 2 but absent in P3HT. Morphology Changes of PProDOT-Hx 2 with Electrochemical Doping During electrochemical doping of PProDOT-Hx 2, the packing and microstructure are altered as the polymer chains reorganize to stabilize the formed charge carriers and accommodate the counter-ions necessary for charge balancing. 65 These rearrangements not only influence the 106 polymer’s electrical conductivity, but also help create an ionic environment within the polymer films that may facilitate Li + diffusion. To understand the structural changes induced by doping, we employ grazing incidence wide-angle x-ray scattering (GIWAXS). Full 2-D GIWAXS patterns (Figure 3.6) provide information about the orientation of the polymer with respect to the substrate, while full or selective integrations can be used to quantify lattice constants and scattering intensity. For these semi-crystalline polymers, only two peaks are generally observed: a lamellar or (100) peak, which indicates the distance between polymer chains across the side- chains, and the π-stacking or (010) peak, which indicates the distance between polymer chains roughly in the direction of the π-stacks. PProDOT-Hx 2 showed a lamellar peak at 0.39 Å -1 and a π- stacking peak at 1.40 Å -1 , corresponding to lattice distances of 16.1 Å and 4.49 Å, respectively. While the (100) peak position is typical of many semiconducting polymers, 82–86 the (010) peak position is quite small (large d-spacing), indicating significant tilting of the π-stacking direction relative to the unit cell, which is also observed in MD simulations. 107 Figure 3.6. Representative 2-D diffractograms of PProDOT-Hx 2. The missing wedge of data in each image results from the grazing incidence geometry used in the experiment. With this understanding of the pristine PProDOT-Hx 2 structure, we then began to explore structural changes that occur upon doping at the first oxidation peak following introduction of charge carriers and charge balancing counterions into the film. Here we specifically considered two common battery electrolytes, LiTFSI and LiClO 4, which provided TFSI - and ClO 4 - anions to the doped polymer film (Figure 3.7a). Upon oxidizing the polymer by holding the potential at the first oxidation peak using these electrolytes, a large increase in the intensity of the (100) diffraction 108 peak was observed, indicating an increase in crystallinity (Figure 3.7b). This increase in crystallinity has been observed previously upon doping in disordered polymers such as regiorandom P3HT; for that system, the increase in crystallinity was hypothesized to be driven by the need to delocalize the polaron between polymer chains, which requires well ordered, π- stacked domains. 83,87 Additionally, as shown in the inset of Figures 3.7a, the (100) peak shifts to lower q upon doping, moving to 0.329 Å -1 and 0.327 Å -1 for LiTFSI and LiClO 4, respectively, while the (010) peak shifts to higher q, moving to 1.48 Å -1 and 1.47 Å -1 , for LiTFSI and LiClO 4, respectively. The shifts of both peaks are similar to those observed when P3HT is chemically doped using strong oxidizing agents and can be understood to indicate a reorientation of the unit cell to make room for the dopant anion in the side chain region of the polymer crystallites. 83,84,87– 92 The rearrangement involves a rotation of the π-stack direction to better align with the unit cell axis and an increase in the side-chain spacing to make more room for the counterions. These results indicated that the anion, whether it be TFSI - or ClO 4 - , resides primarily within the lamellar region of the crystallites. We note that while it might seem easier to dope the amorphous regions of the film, previous studies using chemical doping indicate that the crystalline regions, in fact, dope first. 84,92 This likely occurs because increased delocalization within the crystalline regions reduces the doping potential for these sites. To study the structural effect of holding the potential at the second oxidation peak, we considered only TFSI electrolytes, as these are the most stable in moist-air or with water. The (100) peak for this sample showed a further slight increase in intensity and a further shift to lower q (0.31 Å -1 ) compared to the sample that was only doped at oxidation 1 (Figure 3.7b). These changes indicated that the polymer chains further rearranged to accommodate the higher levels 109 of bipolaron expected to be induced at oxidation 2. At the same time, the (010) peak broadened considerably and shifted to lower q, suggesting some disordering of the π-stacking. While it is clear that by oxidizing the polymer further at oxidation 2, a modified structure was obtained, more studies are needed to understand the details of the structure of such highly doped films. To get a better picture of the expected solvent structure during battery operation, we also examined how solvent incorporates into the polymer matrix. In the presence of the organic solvents in a battery, the polymer is expected to swell. 82,93 To verify this with PProDOT-Hx 2, optical ellipsometry measurements were conducted. For the ellipsometry experiments, PProDOT-Hx 2 was spun onto a glass ITO substrate from a solution of 1,2-dichlorbenzene, resulting in films of 50-60 nm in thickness. The films were placed in a customized quartz vial containing propylene carbonate (PC), designed to allow ellispometry data to be collected upon swelling with solvent vapor. For this work, PC was chosen instead of the standard EC/DMC mixture because of the complexity of using a combination of a volatile and non-volatile solvent, and for direct comparison to the molecular dynamics simulations, discussed below. The PProDOT-Hx 2 film was observed to increase in thickness by 30%, indicating significant solvent incorporation into the films, likely mostly into the amorphous regions. 84,93 We note that the ellipsometry cell may not have been able to fully equilibrate with the solvent, so this value should be considered as a lower bound on the amount of solvent that should be expected to incorporate into binder films during battery operation. Nonetheless, these results clearly indicate that PProDOT-Hx 2 films swell in common battery solvents, and so a swollen film should be considered when considering mechanisms for ionic conductivity. 110 Figure 3.7. Full integration of GIWAXS diffractograms for PProDOT-Hx 2 doped under different conditions. Part (a) shows the polymer doped at oxidation 1 using two common battery electrolytes, LiTFSI and LiClO 4, which provide TFSI - and ClO 4 - anions to the polymer film upon doping. Part (b) compares the structure of the polymer doped at oxidation peaks 1 and 2 using the most stable of our battery electrolyte, LiTFSI. Performance as a Cathode Binder To examine the performance of PProDOT-Hx 2 as a polymer binder in Li-ion batteries, LiNi0.8Co0.15Al0.05O2 (NCA) electrodes incorporating PProDOT-Hx 2, Super P carbon black, and multiwalled carbon nanotubes (CNT) with a weight ratio of 90-3-3-4% (NCA-SP-CNT-PProDOT- Hx 2) were prepared. As can be observed from the TEM image of just NCA and PProDOT-Hx2 (96- 4%) (Figure 3.8a), PProDOT-Hx 2 is homogeneously distributed and acts as a binder to the NCA particles, leading to a well-connected network of active materials. Figure 3.8b shows the TEM image of the full composite electrode. 111 Figure 3.8 TEM image of (a) 96-4% NCA-PProDOT-Hx 2 and (b) 90-3-3-4% NCA-SP-CNT-PProDOT- Hx 2 electrodes. To characterize the rate capability of the NCA-PProDOT-Hx2 electrodes, galvanostatic charge−discharge (GCD) tests were performed in coin cells. The coin cells were assembled using lithium as the counter electrode and 1 M LiTFSI in EC/DMC as the electrolyte. The rate capability of the NCA-PProDOT-Hx 2 electrodes were compared to NCA-PVDF control electrodes, and NCA electrodes with-out the addition of any binder for two mass loadings of ~ 6 mg/cm 2 and 11 mg/ cm2 to demonstrate the benefits of PProDOT-Hx 2 even at higher mass loadings (Figure 3.9a, 3.10, and 3.11). Different C-rates were utilized based on 1 C = 160 mA g -1 as previously reported. 44 At a rate of C/5, the NCA-PProDOT-Hx 2 and NCA-PVDF cathodes delivered nearly identical discharge capacities of 168 and 166 mAh g-1, respectively. At slow rates below 1C, they continued to deliver comparable specific capacities. Nevertheless, at higher rates (above 2C), the NCA-PProDOT-Hx 2 cells exhibited enhanced rate capability compared to the NCA-PVDF cells and the cells without binder. At 6C, the NCA-PProDOT-Hx 2 delivered a capacity of 111 mAh g -1 , while NCA-PVDF only provides 20 mAh g -1 (Figure 3.9a). 112 Figure 3.9. Electrochemical properties for NCA-PProDOT-Hx2 and NCA-PVDF electrodes: (a) Rate capability of the NCA-PProDOT-Hx2 and NCA-PVDF. The corresponding galvanostatic charge- discharge (GCD) curves of the NCA-PProDOT-Hx2 (b) and NCA-PVDF (c) at various rates. (d) Long- term cycling for NCA-PProDOT-Hx2 and NCA-PVDF at a rate of 2C. The corresponding GCD curves of the NCA-PProDOT-Hx 2 (e) and NCA-PVDF (f) at different cycles. NCA electrodes without binder We have built NCA-CNT-Super P electrodes without the addition of binder in weight ratio of 94:3:3 respectively and following the same conditions to test their performance in a cell against lithium. Figure 3.10 below shows the rate capability comparison of the NCA-CNT-Super P (No Binder) electrodes against NCA-PVDF electrodes at two different areal mass loadings of (a) 6 mg/cm 2 and (b) 11 mg/cm 2 . The utilization of the binder free electrodes is lower compared to the PVDF electrodes and at higher rates the effect is magnified, suggesting a lack of connectivity within the active material, sluggish kinetics, and a lower active accessible surface area (porosity). 113 Figure 3.12 below shows the corresponding galvanostatic charge/discharge curves for the NCA- CNT-Super P (No Binder) cells at (a) 6 mg/cm 2 and (b) 11 mg/cm 2 . NCA composite electrodes with a higher loading of 11 mg/cm 2 Typical research lab cells generally only have a loading of 1-2 mg/cm 2 , but here we have already used 6 mg/cm 2 which is already close to commercial cells. Nonetheless, we have built NCA- PProDOT-Hx 2 and NCA-PVDF (control) electrodes with a higher areal loading of 11 mg/cm 2 and tested them in a cell against lithium under the same conditions in order to demonstrate the benefits of PProDOT-Hx 2 as a conductive additive/binder. Figure 3.11 below shows the rate capability comparison of the NCA-PProDOT-Hx 2, NCA-PVDF and NCA-CNT-Super P (No Binder) at a higher areal loading of 11 mg/cm 2 . The results show how even at higher loadings of active material the addition of PProDOT-Hx 2 has remarkable benefits in the cell’s capacity at all rates compared to the NCA-PVDF and NCA binder free cells. Figure 3.13 below shows the corresponding galvanostatic charge/discharge curves for the NCA-PProDOT-Hx 2 (a) and NCA- PVDF (b) at various rates with an areal loading of 11 mg cm -2 . The results at higher loadings of NCA (11 mg/cm 2 ) con-firms the benefits of PProDOT-Hx 2 as conductive additive and binder by delivering at 4C a capacity of 30 mAh g -1 while NCA-PVDF only provides 1 mAh g -1 (Figure 3.11). We can attribute these improved results with PProDOT-Hx 2 to the enhanced electronic and ionic conductivity of the doped PProDOT-Hx 2 compared to the insulating PVDF binder. The NCA-PProDOT-Hx 2 cells also showed significantly reduced polarization at all rates above 1C compared to the NCA-PVDF cells and the NCA cells without binder for both active mass loadings (Figure 3.9b,c, 3.12a.b, and 3.13a,b). 114 Figure 3.10 Rate capability comparison of the NCA-PVDF and NCA-CNT-Super P (No Binder) with an active material mass loading of 6 mg cm -2 (a) and 11 mg cm -2 (b). Figure 3.11 Rate capability comparison of the NCA-PProDOT-Hx 2, NCA-PVDF and NCA-CNT- SuperP (No Binder) with an active material mass loading of 11 mg cm -2 . 115 Figure 3.12 Galvanostatic charge-discharge (GCD) curves at various rates of the NCA-CNT-Super P (No Binder) with an active material mass loading of 6 mg cm -2 (a) and 11 mg cm -2 (b). Figure 3.13 Galvanostatic charge-discharge (GCD) curves of the NCA-PProDOT-Hx 2 (a) and NCA- PVDF (b) at various rates with an active material mass loading of 11 mg cm -2 . To study the effect of PProDOT-Hx 2 on the electrode performance, we determined the ICE for the NCA- PProDOT-Hx 2 and NCA-PVDF electrodes (Figure 3.14). The values are 85.4% and 84.8%, respectively, the fact that the ICE values are very similar indicates the negligible effect that the low ICE values have on the doping process for the conductive polymer during charge and discharge. 116 Figure 3.14: The initial cycle charge and discharge profiles for NCA-PProDOT-Hx 2 and NCA-PVDF electrodes at a rate of C/5. To characterize the inner structure of the NCA-PProDOT-Hx 2 electrodes, we have compared against the NCA-PVDF electrodes to have a better understanding of how structure affects the enhanced rate capability. Therefore, we have measured the impedance response as a function of state of charge (SOC) during charge/discharge in order to understand the changes associated with the internal resistances of the cell and the accessible electroactive sur-face area of the electrode. Determination of the electroactive surface area of the composite electrodes with PProDOT-Hx 2 by electrochemical impedance spectroscopy. Porosity and impedance analysis of the NCA electrodes. Figure 3.15 shows the galvanostatic discharge/charge (GCD) curves at C/10 and the intervallic measurement of potentiostatic electrochemical impedance spectroscopy (PEIS) as function of SOC (denoted with green circles) at open circuit potential (OCV) for the NCA-PProDOT-Hx 2 (a) and 117 NCA-PVDF (b) cells. From the GCD curves and PEIS measurements, we can observe a remarkable lower polarization in the NCA-PProDOT-Hx 2 cell reflected in the narrow gap between the OCV and the galvanostatic charging/discharge curve. In contrast, the NCA-PVDF cells showed a higher voltage gap and polarization. The impedance response for both cells has been represented in a 3D Nyquist plot as function of SOC shown below in Figure 3.16. From the 3D charts we can observe that the overall resistance is indirectly proportional to the SOC for both cells. However, the overall resistance for the NCA- PProDOT-Hx 2 cell is approximately 3 times lower than the NCA-PVDF cell. To understand the impedance response as function of SOC, we have decided to utilize a well- known Transmission Line Model for a porous electrode made of cylindrical pores. 94–97 The model denotes that in the presence of a faradaic process at low frequencies (w) the real component (Zre) of the impedance response can be approximated to: 𝑍 𝑟𝑒 = 𝑅 𝑖𝑜𝑛 3 + 𝑅 𝑐𝑡 ( 3) Where R ion is the Li ion mobility within the pore, and R ct is the charge transfer resistance related to the electroactive surface area. Ogihara et al. 96 have shown that R ion is dependent on the electrode thickness. For our experiments, the PProDOT-Hx 2 and PVDF electrodes had approximately the same thickness (28 microns), therefore, we can consider the R ion value to be the same for both cells. With this assumption the analysis of the Zre at low frequencies is now only dependent on the R ct or the electroactive surface area of the electrode. Figure 3.17 shows a comparison of the experimental 118 Zre response as function of SOC at low frequencies for the NCA-PProDOT-Hx 2 cells and the NCA- PVDF cells. We can observe that the Zre of the NCA-PVDF cells is almost 3 times higher than the NCA-PProDOT-Hx 2 cells suggesting that the NCA-PProDOT-Hx 2 electrodes have an electroactive surface area 3 times higher than those made with PVDF. These results suggest an enhanced porosity in the NCA-PProDOT-Hx 2 cells using this typical analysis for electrode porosity. However, we note that effects other than enhanced surface area, such as the enhanced conductivity of PProDOT-Hx 2 relative to PVDF could influence the measurements. Nonetheless, this analysis points to a more porous NCA-PProDOT-Hx 2 cells relative to NCA-PVDF. The results suggest that the overall resistance of the NCA-PVDF cell is about three times higher than the NCA-PProDOT-Hx 2 cells (Figure 3.15 and 3.16). In addition, the analysis of the impedance response in the low frequency region suggests that the NCA- PProDOT- Hx 2 electrodes have a three times higher porosity than those made of PVDF (Figure 3,17). Figure 3.15 (a) Comparison of the charge and discharge curve of the NCA-PProDOT-Hx 2 (a) and NCA-PVDF (b) at at C/10, EIS was measured as function of SOC (denoted with green circles). 119 Figure 3.16 A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of charge for the NCA-PProDOT-Hx 2 (a) and NCA-PVDF (b) cell with a mass loading of 11 mg cm -2 . Figure 3.17 Real part of the impedance response agaisnt frequency (at low freqnuencies) as fucntion of SOC for the NCA-PProDOT-Hx 2 (black) and NCA-PVDF (red) cells. By maintaining the same electrode components ratios, we also made the comparison of PProDOT-Hx 2 and P3HT as electrode additives against commercial PVDF binder to confirm their comparable and superior performance at higher rates (Figure 3.18). 120 Figure 3.18 Rate capability comparison of the NCA-PProDOT-Hx 2, NCA-PVDF, and NCA-P3HT with an active material mass loading of ~ 6 mg cm -2 . While high rate is important for many applications, stability is also key, and PVDF is designed for its chemical inertness. Surprisingly, a long-term cycling comparison at 2C between the PProDOT- Hx 2 and PVDF cells showed that in the first 120 cycles, the capacities are comparable, but after that point, a decay is observed in the NCA-PVDF cells that is not observed in those made with PProDOT-Hx2 (Figure 3.9d). Thus, the NCA-PProDOT-Hx 2 cathode showed better cycling stability with less decay of the potential drop in the GCD curves compared to NCA-PVDF electrodes. Overall, this indicates that the PProDOT-Hx2 polymer was more stable during long-term cycling (Figure 3.9e and 3.9f). Conclusions PProDOT-Hx 2 is an electrochemically stable conjugated polymer that is amenable to large scale synthesis and facile processing as a conductive electrode binder. Thin films of PProDOT-Hx 2 undergo reversible electrochemical doping in the potential range suitable for most state-of-art lithium-ion battery cathode materials and can be cycled up to +4.5 V (vs. Li/Li + ) without any 121 noticeable change in electrochemical reversibility. The rapid kinetics of electrochemical doping make PProDOT-Hx 2 an ideal binder for cathode composites. Importantly, PProDOT-Hx 2 showed dual ion/electron conducting properties, with excellent electronic conductivity of 0.1 S cm -1 over the entire potential window of the doped state, rivalling P3HT. The ionic conductivity of PProDOT-Hx 2 in the doped state was at least one order of magnitude higher than that of P3HT, consistent with the opportunity for increased solvation of ions induced by the oxygen atoms present in PProDOT-Hx 2. Molecular dynamics simulations indicate that the nature of the Li + ion solvation environments change considerably more with solvent exposure, than morphology does, indicating that solvent swelling plays an important role in ionic conductivity. The undoped polymer exhibited short range ordering at two length scales corresponding to the lamellar spacing between the chains and the π-π stacking distance with a significant tilting of the π-stacking direction relative to the unit cell. These structure assignments were consistent with the equilibrium ordering predicted by molecular dynamics. It was found that doping caused significant shifts in the lamellar and -stacking spacing, indicating that the dopant anions lie in the lamellar region and that the unit cell reoriented to accommodate the dopant. Upon extensive doping, the polymer chains were observed to further rearrange to accommodate additional bipolaron formation with a resultant disordering of the π-stacking in the polymer. Importantly, the performance of NCA electrodes was significantly enhanced when PProDOT-Hx 2 was used as a binder. Specifically, NCA-PProDOT-Hx 2 electrodes showed at least five times higher capacity compared to NCA-PVDF electrodes at 6C. Furthermore, when cycled at 2C for 200 cycles, 122 NCA-PVDF electrodes showed significant decay in their capacity while NCA-PProDOT-Hx 2 electrodes maintained their specific capacity, indicating that PProDOT-Hx 2 improved the cycle stability of NCA when used as a conductive binder. Considering the ease of chemical functionalization in PProDOTs, it is clear that they are a family of conjugated polymers that is well-suited as conductive binders in cathodes for lithium-ion batteries. Their relatively high electronic and ionic conductivities, coupled with their electrochemical stability and ease of functionalization, paves the way for high-performance cathodes with conductive binders that are specifically tailored to enhance electrode stability and rate capability. References 1. Das, P. et al. Dihexyl-Substituted Poly(3,4-Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. acs.chemmater.0c02601 (2020) doi:10.1021/acs.chemmater.0c02601. 2. Yang, Z. et al. Electrochemical Energy Storage for Green Grid. Chem. Rev. 111, 3577–3613 (2011). 3. Thackeray, M. M., Wolverton, C. & Isaacs, E. D. Electrical energy storage for transportation— approaching the limits of, and going beyond, lithium-ion batteries. Energy Environ. Sci. 5, 7854 (2012). 4. Dunn, B., Kamath, H. & Tarascon, J.-M. Electrical Energy Storage for the Grid: A Battery of Choices. Science 334, 928–935 (2011). 5. Chen, H. et al. Exploring Chemical, Mechanical, and Electrical Functionalities of Binders for Advanced Energy-Storage Devices. Chem. Rev. 118, 8936–8982 (2018). 6. Dudney, N. J. & Li, J. Using all energy in a battery. Science 347, 131–132 (2015). 7. Lopez, J., Mackanic, D. G., Cui, Y. & Bao, Z. Designing polymers for advanced battery chemistries. Nat Rev Mater 4, 312–330 (2019). 8. Bounioux, C. et al. Thermoelectric composites of poly(3-hexylthiophene) and carbon nanotubes with a large power factor. Energy Environ. Sci. 6, 918 (2013). 9. Cho, S. I. & Lee, S. B. Fast Electrochemistry of Conductive Polymer Nanotubes: Synthesis, Mechanism, and Application. Acc. Chem. Res. 41, 699–707 (2008). 10. Das, P. R., Komsiyska, L., Osters, O. & Wittstock, G. PEDOT: PSS as a Functional Binder for Cathodes in Lithium Ion Batteries. J. Electrochem. Soc. 162, A674–A678 (2015). 123 11. Higgins, T. M. et al. A Commercial Conducting Polymer as Both Binder and Conductive Additive for Silicon Nanoparticle-Based Lithium-Ion Battery Negative Electrodes. ACS Nano 10, 3702–3713 (2016). 12. Hughes, M., Chen, G. Z., Shaffer, M. S. P., Fray, D. J. & Windle, A. H. Electrochemical Capacitance of a Nanoporous Composite of Carbon Nanotubes and Polypyrrole. Chem. Mater. 14, 1610–1613 (2002). 13. Kaufman, J. H., Chung, T. ‐C., Heeger, A. J. & Wudl, F. Poly(Thiophene): A Stable Polymer Cathode Material. J. Electrochem. Soc. 131, 2092–2093 (1984). 14. Shi, Y., Zhou, X. & Yu, G. Material and Structural Design of Novel Binder Systems for High-Energy, High-Power Lithium-Ion Batteries. Acc. Chem. Res. 50, 2642–2652 (2017). 15. Wang, H. et al. Cationic polymer binder inhibit shuttle effects through electrostatic confinement in lithium sulfur batteries. J. Mater. Chem. A 6, 6959–6966 (2018). 16. Wu, M. et al. Toward an Ideal Polymer Binder Design for High-Capacity Battery Anodes. J. Am. Chem. Soc. 135, 12048–12056 (2013). 17. Ye, L. et al. Quantitative relations between interaction parameter, miscibility and function in organic solar cells. Nature Mater 17, 253–260 (2018). 18. Zeng, W. et al. Enhanced Ion Conductivity in Conducting Polymer Binder for High-Performance Silicon Anodes in Advanced Lithium-Ion Batteries. Adv. Energy Mater. 8, 1702314 (2018). 19. Kwon, Y. H. et al. Toward Uniformly Dispersed Battery Electrode Composite Materials: Characteristics and Performance. ACS Appl. Mater. Interfaces 8, 3452–3463 (2016). 20. Patnaik, S. G., Vedarajan, R. & Matsumi, N. BIAN based functional diimine polymer binder for high performance Li ion batteries. J. Mater. Chem. A 5, 17909–17919 (2017). 21. Park, S.-J. et al. Side-Chain Conducting and Phase-Separated Polymeric Binders for High- Performance Silicon Anodes in Lithium-Ion Batteries. J. Am. Chem. Soc. 137, 2565–2571 (2015). 22. Minnici, K. et al. Carboxylated Poly(thiophene) Binders for High-Performance Magnetite Anodes: Impact of Cation Structure. ACS Appl. Mater. Interfaces 11, 44046–44057 (2019). 23. Cao, Y. et al. Conductive Polymers Encapsulation To Enhance Electrochemical Performance of Ni- Rich Cathode Materials for Li-Ion Batteries. ACS Appl. Mater. Interfaces 10, 18270–18280 (2018). 24. Li, X., An, H., Strzalka, J., Lutkenhaus, J. & Verduzco, R. Self-Doped Conjugated Polymeric Binders Improve the Capacity and Mechanical Properties of V2O5 Cathodes. Polymers 11, 589 (2019). 25. Kwon, Y. H. et al. SWNT Networks with Polythiophene Carboxylate Links for High-Performance Silicon Monoxide Electrodes. ACS Appl. Energy Mater. 1, 2417–2423 (2018). 26. Welsh, D. M. et al. Regiosymmetric Dibutyl-Substituted Poly(3,4-propylenedioxythiophene)s as Highly Electron-Rich Electroactive and Luminescent Polymers. Macromolecules 35, 6517–6525 (2002). 27. Wu, H. et al. Stable Li-ion battery anodes by in-situ polymerization of conducting hydrogel to conformally coat silicon nanoparticles. Nat Commun 4, 1943 (2013). 28. Zhao, H. et al. Toward Practical Application of Functional Conductive Polymer Binder for a High- Energy Lithium-Ion Battery Design. Nano Lett. 14, 6704–6710 (2014). 124 29. Kumar, D. & Sharma, R. C. Advances in conductive polymers. European Polymer Journal 34, 1053– 1060 (1998). 30. Harima, Y., Kunugi, Y., Yamashita, K. & Shiotani, M. Determination of mobilities of charge carriers in electrochemically anion-doped polythiophene film. Chemical Physics Letters 317, 310–314 (2000). 31. Kunugi, Y., Harima, Y., Yamashita, K., Ohta, N. & Ito, S. Charge transport in a regioregular poly(3- octylthiophene) film. J. Mater. Chem. 10, 2673–2677 (2000). 32. Jiang, X., Patil, R., Harima, Y., Ohshita, J. & Kunai, A. Influences of Self-Assembled Structure on Mobilities of Charge Carriers in π-Conjugated Polymers. J. Phys. Chem. B 109, 221–229 (2005). 33. Sun, C., Gobetto, R. & Nervi, C. Recent advances in catalytic CO 2 reduction by organometal complexes anchored on modified electrodes. New J. Chem. 40, 5656–5661 (2016). 34. Liew, C., Durairaj, R. & Ramesh, S. Rheological Studies of PMMA–PVC Based Polymer Blend Electrolytes with LiTFSI as Doping Salt. PLoS ONE 9, e102815 (2014). 35. Kim, S.-M. et al. Poly(phenanthrenequinone) as a conductive binder for nano-sized silicon negative electrodes. Energy Environ. Sci. 8, 1538–1543 (2015). 36. Minnici, K. et al. Tuning Conjugated Polymers for Binder Applications in High-Capacity Magnetite Anodes. ACS Appl. Energy Mater. 2, 7584–7593 (2019). 37. Kwon, Y. H. et al. Electron/Ion Transport Enhancer in High Capacity Li-Ion Battery Anodes. Chem. Mater. 28, 6689–6697 (2016). 38. Liu, D. et al. Novel conductive binder for high-performance silicon anodes in lithium ion batteries. Nano Energy 36, 206–212 (2017). 39. Ma, W., Yang, C., Gong, X., Lee, K. & Heeger, A. J. Thermally Stable, Efficient Polymer Solar Cells with Nanoscale Control of the Interpenetrating Network Morphology. Adv. Funct. Mater. 15, 1617–1622 (2005). 40. Lai, C.-H. et al. Application of Poly(3-hexylthiophene-2,5-diyl) as a Protective Coating for High Rate Cathode Materials. Chem. Mater. 30, 2589–2599 (2018). 41. Dong, B. X. et al. Influence of Side-Chain Chemistry on Structure and Ionic Conduction Characteristics of Polythiophene Derivatives: A Computational and Experimental Study. Chem. Mater. 31, 1418–1429 (2019). 42. Schmode, P. et al. The Key Role of Side Chain Linkage in Structure Formation and Mixed Conduction of Ethylene Glycol Substituted Polythiophenes. ACS Appl. Mater. Interfaces 12, 13029–13039 (2020). 43. Minnici, K. et al. High capacity Li-ion battery anodes: Impact of crystallite size, surface chemistry and PEG-coating. Electrochimica Acta 260, 235–245 (2018). 44. Pud, A. A. Stability and degradation of conducting polymers in electrochemical systems. Synthetic Metals 66, 1–18 (1994). 45. Tsai, E. W., Basak, S., Ruiz, J. P., Reynolds, J. R. & Rajeshwar, K. Electrochemistry of some β‐ Substituted Polythiophenes: Anodic Oxidation, Electrochromism, and Electrochemical Deactivation. J. Electrochem. Soc. 136, 3683–3689 (1989). 125 46. Welsh, D. M., Kumar, A., Meijer, E. W. & Reynolds, J. R. Enhanced Contrast Ratios and Rapid Switching in Electrochromics Based on Poly(3,4-propylenedioxythiophene) Derivatives. Adv. Mater. 11, 1379–1382 (1999). 47. Groenendaal, B., Friedrich, J., Dieter, F., Pielartzik, H. & Reynolds, J. R. Poly(3,4- ethylenedioxythiophene) and Its Derivatives: Past, Present, and Future. Advanced Materials 12, (2000). 48. Kumar, A. et al. Conducting Poly(3,4-alkylenedioxythiophene) Derivatives as Fast Electrochromics with High-Contrast Ratios. Chem. Mater. 10, 896–902 (1998). 49. Österholm, A. M., Ponder, J. F., Kerszulis, J. A. & Reynolds, J. R. Solution Processed PEDOT Analogues in Electrochemical Supercapacitors. ACS Appl. Mater. Interfaces 8, 13492–13498 (2016). 50. Lang, A. W. et al. Flexible, aqueous-electrolyte supercapacitors based on water-processable dioxythiophene polymer/carbon nanotube textile electrodes. J. Mater. Chem. A 5, 23887–23897 (2017). 51. Ponder, J. F., Menon, A. K. & Dasari, R. R. Conductive, Solution-Processed Dioxythiophene Copolymers for Thermoelectric and Transparent Electrode Applications. Advanced Energy Materials 9, (2019). 52. Savagian, L. R. et al. Balancing Charge Storage and Mobility in an Oligo(Ether) Functionalized Dioxythiophene Copolymer for Organic- and Aqueous- Based Electrochemical Devices and Transistors. Advanced Materials 30, (2018). 53. Acharya, S. et al. Ultrahigh stability of high-power nanofibrillar PEDOT supercapacitors. Sustainable Energy Fuels 1, 482–491 (2017). 54. Österholm, A. M., Shen, D. E., Dyer, A. L. & Reynolds, J. R. Optimization of PEDOT Films in Ionic Liquid Supercapacitors: Demonstration As a Power Source for Polymer Electrochromic Devices. ACS Appl. Mater. Interfaces 5, 13432–13440 (2013). 55. Lepage, D., Michot, C., Liang, G., Gauthier, M. & Schougaard, S. B. A Soft Chemistry Approach to Coating of LiFePO4 with a Conducting Polymer. Angew. Chem. 123, 7016–7019 (2011). 56. Liu, X., Li, H., Li, D., Ishida, M. & Zhou, H. PEDOT modified LiNi 1/3 Co 1/3 Mn 1/3 O 2 with enhanced electrochemical performance for lithium ion batteries. Journal of Power Sources 243, 374–380 (2013). 57. Her, L.-J., Hong, J.-L. & Chang, C.-C. Preparation and electrochemical characterizations of poly(3,4- dioxyethylenethiophene)/LiCoO2 composite cathode in lithium-ion battery. Journal of Power Sources 157, 457–463 (2006). 58. McDonald, M. B. & Hammond, P. T. Efficient Transport Networks in a Dual Electron/Lithium- Conducting Polymeric Composite for Electrochemical Applications. ACS Appl. Mater. Interfaces 10, 15681– 15690 (2018). 59. Siju, C. R., Saravanan, T. R., Rao, K. N. & Sindhu, S. Optical, Electrochemical, and Structural Properties of Spray Coated Dihexyl Substituted Poly (3,4 Propylene Dioxythiophene) Film for Optoelectronics Devices. International Journal of Polymeric Materials and Polymeric Biomaterials 63, 374– 379 (2014). 60. Reeves, B. D. et al. Spray Coatable Electrochromic Dioxythiophene Polymers with High Coloration Efficiencies. Macromolecules 37, 7559–7569 (2004). 126 61. Hwang, J. et al. In situ measurements of the optical absorption of dioxythiophene-based conjugated polymers. Phys. Rev. B 83, 195121 (2011). 62. Kalagi, S. S. & Patil, P. S. Secondary electrochemical doping level effects on polaron and bipolaron bands evolution and interband transition energy from absorbance spectra of PEDOT: PSS thin films. Synthetic Metals 220, 661–666 (2016). 63. Wei, Q., Mukaida, M., Kirihara, K., Naitoh, Y. & Ishida, T. Photoinduced Dedoping of Conducting Polymers: An Approach to Precise Control of the Carrier Concentration and Understanding Transport Properties. ACS Appl. Mater. Interfaces 8, 2054–2060 (2016). 64. Yamamoto, J. & Furukawa, Y. Electronic and Vibrational Spectra of Positive Polarons and Bipolarons in Regioregular Poly(3-hexylthiophene) Doped with Ferric Chloride. J. Phys. Chem. B 119, 4788– 4794 (2015). 65. Lin, C., Endo, T., Takase, M., Iyoda, M. & Nishinaga, T. Structural, Optical, and Electronic Properties of a Series of 3,4-Propylenedioxythiophene Oligomers in Neutral and Various Oxidation States. J. Am. Chem. Soc. 133, 11339–11350 (2011). 66. Choi, C. et al. Achieving high energy density and high power density with pseudocapacitive materials. Nat Rev Mater 5, 5–19 (2020). 67. Kivelson, S. & Heeger, A. J. First-order transition to a metallic state in polyacetylene: A strong- coupling polaronic metal. Phys. Rev. Lett. 55, 308–311 (1985). 68. Miller, L. L. & Mann, K. R. π-Dimers and π-Stacks in Solution and in Conducting Polymers. Acc. Chem. Res. 29, 417–423 (1996). 69. van Haare, J. A. E. H. et al. Redox States of Long Oligothiophenes: Two Polarons on a Single Chain. Chemistry A European J 4, 1509–1522 (1998). 70. Enengl, C. et al. Doping-Induced Absorption Bands in P3HT: Polarons and Bipolarons. ChemPhysChem 17, 3836–3844 (2016). 71. Nowak, M. J., Spiegel, D., Hotta, S., Heeger, A. J. & Pincus, P. A. Charge storage on a conducting polymer in solution. 22, 10 (1989). 72. Reeves, B. D., Unur, E., Ananthakrishnan, N. & Reynolds, J. R. Defunctionalization of Ester- Substituted Electrochromic Dioxythiophene Polymers. Macromolecules 40, 5344–5352 (2007). 73. Bertho, D. & Jouanin, C. Polaron and bipolaron excitations in doped polythiophene. Phys. Rev. B 35, 626–633 (1987). 74. Cao, J. & Curtis, M. D. Polarons, Bipolarons, and π-Dimers of Bis(3,4-ethylene-dioxythiophene)- (4,4‘-dialkyl-2,2‘-bithiazole)- co -Oligomers. Direct Measure of the Intermolecular Exciton Transfer Interaction. Chem. Mater. 15, 4424–4430 (2003). 75. Aubry, T. J. et al. Dodecaborane-Based Dopants Designed to Shield Anion Electrostatics Lead to Increased Carrier Mobility in a Doped Conjugated Polymer. Advanced Materials 31, (2019). 76. Liu, L., Qui, L. & Alessandri, R. Enhancing Molecular n-Type Doping of Donor–Acceptor Copolymers by Tailoring Side Chains. Advanced Materials 30, (2018). 127 77. Obrzut, J. & Page, K. A. Electrical conductivity and relaxation in poly(3-hexylthiophene). Phys. Rev. B 80, 195211 (2009). 78. Patel, S. N., Javier, A. E. & Balsara, N. P. Electrochemically Oxidized Electronic and Ionic Conducting Nanostructured Block Copolymers for Lithium Battery Electrodes. ACS Nano 7, 6056–6068 (2013). 79. Huggins, R. A. Simple method to determine electronic and ionic components of the conductivity in mixed conductors a review. Ionics 8, 300–313 (2002). 80. Jamnik, J. & Maier, J. Treatment of the Impedance of Mixed Conductors Equivalent Circuit Model and Explicit Approximate Solutions. J. Electrochem. Soc. 146, 4183–4188 (1999). 81. Pickup, P. G. Alternating current impedance study of a polypyrrole-based anion-exchange polymer. Faraday Trans. 86, 3631 (1990). 82. Manley, E. F. et al. In Situ GIWAXS Analysis of Solvent and Additive Effects on PTB7 Thin Film Microstructure Evolution during Spin Coating. Adv. Mater. 29, 1703933 (2017). 83. Yee, P. Y., Scholes, D. T., Schwartz, B. J. & Tolbert, S. H. Dopant-Induced Ordering of Amorphous Regions in Regiorandom P3HT. J. Phys. Chem. Lett. 10, 4929–4934 (2019). 84. Scholes, D. T. et al. The Effects of Crystallinity on Charge Transport and the Structure of Sequentially Processed F 4 TCNQ‐Doped Conjugated Polymer Films. Adv. Funct. Mater. 27, 1702654 (2017). 85. Cochran, J. E. et al. Molecular Interactions and Ordering in Electrically Doped Polymers: Blends of PBTTT and F 4 TCNQ. Macromolecules 47, 6836–6846 (2014). 86. Mazaheripour, A., Thomas, E. M., Segalman, R. A. & Chabinyc, M. L. Nonaggregating Doped Polymers Based on Poly(3,4-Propylenedioxythiophene). Macromolecules 52, 2203–2213 (2019). 87. Lim, E., Glaudell, A. M., Miller, R. & Chabinyc, M. L. The Role of Ordering on the Thermoelectric Properties of Blends of Regioregular and Regiorandom Poly(3‐hexylthiophene). Adv. Electron. Mater. 5, 1800915 (2019). 88. Fontana, M. T. et al. Evaporation vs Solution Sequential Doping of Conjugated Polymers: F 4 TCNQ Doping of Micrometer-Thick P3HT Films for Thermoelectrics. J. Phys. Chem. C 123, 22711–22724 (2019). 89. Hynynen, J. et al. Enhanced Electrical Conductivity of Molecularly p-Doped Poly(3- hexylthiophene) through Understanding the Correlation with Solid-State Order. Macromolecules 50, 8140–8148 (2017). 90. Lim, E., Peterson, K. A., Su, G. M. & Chabinyc, M. L. Thermoelectric Properties of Poly(3- hexylthiophene) (P3HT) Doped with 2,3,5,6-Tetrafluoro-7,7,8,8-tetracyanoquinodimethane (F 4 TCNQ) by Vapor-Phase Infiltration. Chem. Mater. 30, 998–1010 (2018). 91. Karpov, Y., Kiriy, N. & Formanek, P. Sequentially Processed P3HT/CN6-CP•−NBu4+ Films: Interfacial or Bulk Doping? Advanced Electronic Materials 6, (2020). 92. Hamidi-Sakr, A., Biniek, L., Bantignies, J.-L. & Maurin, D. A Versatile Method to Fabricate Highly In-Plane Aligned Conducting Polymer Films with Anisotropic Charge Transport and Thermoelectric Properties: The Key Role of Alkyl Side Chain Layers on the Doping Mechanism. Advanced Functional Materials 27, (2017). 128 93. Fontana, M. T. et al. Low-Vapor-Pressure Solvent Additives Function as Polymer Swelling Agents in Bulk Heterojunction Organic Photovoltaics. J. Phys. Chem. C 122, 16574–16588 (2018). 94. de Levie, R. On porous electrodes in electrolyte solutions. Electrochimica Acta 8, 751–780 (1963). 95. Ogihara, N. et al. Theoretical and Experimental Analysis of Porous Electrodes for Lithium-Ion Batteries by Electrochemical Impedance Spectroscopy Using a Symmetric Cell. J. Electrochem. Soc. 159, A1034–A1039 (2012). 96. Ogihara, N., Itou, Y., Sasaki, T. & Takeuchi, Y. Impedance Spectroscopy Characterization of Porous Electrodes under Different Electrode Thickness Using a Symmetric Cell for High-Performance Lithium-Ion Batteries. J. Phys. Chem. C 119, 4612–4619 (2015). 97. Itagaki, M., Suzuki, S., Shitanda, I. & Watanabe, K. Electrochemical Impedance and Complex Capacitance to Interpret Electrochemical Capacitor. Electrochemistry 75, 649–655 (2007). 129 Chapter 4 - High-Energy/High-Rate π-Conjugated Polymers as Binders for Lithium Metal Batteries 1 Abstract Developing lithium-ion batteries with both high specific energy and high-power capability is a challenging task because of the necessity for meeting conflicting design requirements. We show that high-energy and high-rate capability can be achieved by using various π-conjugated p- dopable polymers as binders at the cathode and by lowering the mass fraction of all the inactive components of the cell. We report a lithium-metal battery that can deliver 320 Wh kg -1 at C/2 using a mass-efficient cell design. To this end, three conducting polymers with different ionic and electronic conductivities have been studied; dihexyl-substituted poly(3,4- propylenedioxythiophene) (PProDOT-Hx 2), poly(3-hexylthiophene) (P3HT), and a new Random Copolymer (Hex:OE)(80:20) PProDOT. These conducting polymers are compared against a conventional polymer binder, PVDF. We show that under the mass-efficient conditions required for achieving high specific energy and rate capability, the conducting polymers play a crucial role by providing electronic and ionic conductivity, protection against rapid growth of solid electrolyte interphase (SEI), and access to a large electrochemically active surface area. Thus, the use of conducting polymers with appropriate molecular structure as binders opens a viable pathway to maximizing the specific energy and rate capability of lithium-ion battery cathodes. Introduction The continuous development and improvement of high-energy and high-rate lithium-ion batteries (LIBs) is crucial to satisfy the ever-growing demand for lighter portable devices with 130 long run-times. 2–5 This persistent need has been driving the scientific research towards new materials, compositions, and cell configurations. 6–11 To this end, conducting polymers (CPs) are attractive as additives/binders for electrodes in LIBs. These CP binders reduce the electrode resistance, protect the surface of the intercalation materials, and enhance lithium-ion transport, leading to increased rate capability, discharge capacity and cycle life. 12–16 Furthermore, pairing of a high-capacity cathode material with a lithium metal anode, to take advantage of the high theoretical capacity (3860 mAh g -1 ) and low density (0.535 g mL -1 ) of lithium, has been getting more attention. 17–20 Although there have been several reports demonstrating the benefits of CPs as binder/additives with cathode and anode materials 16,21–28 , many of these studies employ low areal mass loadings of active material (thin electrodes, 1 to 5 mg cm -2 ), with high mass fractions of carbon additive (3 to 10%) and/or excessive amounts of electrolyte and anode material. For example, previous work from our team has explored the stability and electrochemical properties of dihexyl-substituted poly(3,4-propylenedioxythiophene) (PProDOT-Hx 2) 29 and poly(3- hexylthiophene) (P3HT) 14 as an electrode additive. These studies were on cathodes with ample amounts of carbon additive (6 wt%), an excess of electrolyte (5.5 µL mg -1 ), and a low areal mass loading (1.5 and 6 mg cm -2 ). Such “non-limiting” conditions, common in the scientific literature, differ significantly from those used in cells built for practical applications. Specifically, in the practical design of commercial LIBs, the amount of “inactive” materials is curtailed to the extreme, and a mass loading of 10-15 mg cm -2 is commonly used to maximize the specific energy. As a result, translating the improvements showed by “non-limiting” designs to realize a high energy cell with high power density is not often feasible despite the valid claims of a new material or an improved formulation. 30–33 Improvements directed at high-energy and high rate-capability 131 must focus on “mass-efficient” designs where the mass of inactive materials namely everything other than the mass of the cathode and anode materials are minimized and the electrodes have a high areal mass loading. In this study, we demonstrate how various π-conjugated p-dopable polymers binders/additives can achieve a high specific energy/high-rate lithium metal-based battery at the cell level. Replacing common cathode binders such as polyvinylidene fluoride (PVDF) with CPs allows us to achieve an impressive specific energy (at the cell component level, excluding the cell packaging) of 320 Wh kg -1 at a practical high rate of C/2. Most importantly, we have found that the CPs have a multifunctional role in preserving the specific energy of the cell under “mass-efficient” conditions, compared to the “non-limiting” designs used in scientific studies. We have revisited the use of PProDOT-Hx 2 and P3HT as binders for a mass-efficient cell design, and we also introduce a third new polymer with enhanced ionic conductivity denoted as (Hex:OE)(80:20) PProDOT. We have utilized extensively electrochemical impedance spectroscopy (EIS), differential capacity (dQ dV -1 vs V) analysis and characterization of electrode morphology to gain insight into the enabling role of the CPs in modifying the underlying chemical and transport phenomena for achieving high specific energy and high rate capability. 132 Cathodes with new π-conjugated polymers for a high-energy lithium battery A comparison of three types of CPs allowed us to understand the benefits of the role of polymer structure in achieving high specific energy and rate capability. Accordingly, in addition to PProDOT-Hx 2 and P3HT, we synthesized a new variety of dihexylPProDOT polymer with 20% oligoether (OE) chains referred to here as (Hex:OE)(80:20) (Figure 2.5- 2.8) 14,29 . The new (Hex:OE)(80:20) underwent reversible electrochemical doping in the potential range of 3.2 to 4.5 V vs Li + /Li (Figure 4.1a-c). The slight asymmetry of the peaks in the voltammograms during oxidation and reduction is attributed to the reversible morphological changes that occur in the polymer film due to solvent uptake during cycling 29 When doped, (Hex:OE)(80:20) had a maximum electronic conductivity of ~ 10 -2 S cm -1 and ionic conductivity of ~ 10 -6 S cm -1 over the potential range of 3.5 V to 4.2 V vs Li + /Li as measured by the methods developed previously in our group. 34 The addition of the OE side group in the dihexylPProDOT structure enhanced the ionic conductivity compared to PProDOT-Hx 2 and P3HT (Figure 4.1d). The polymers stayed doped with small changes to the conductivity values over the potential range of 3.3 V to 4.2 V vs Li + /Li that encompasses the potential window of cycling of typical cathode materials such as NCA and NMC622. 133 Figure 4.1 CVs of (Hex:OE)(80:20) a) at various potential intervals at 50 mV s -1 , b) at various scan rates, and c) over 100 cycles at 10 mV s -1 . d) Electronic and Ionic Conductivities of P3HT, PProDOT- Hx 2 and (Hex:OE)(80:20). We have tested these polymers with NCA as the electrode active material due to NCA’s excellent reported stability 35,36 , high theoretical charge capacity (265 mAh g -1 ) 35 , high electronic conductivity (~ 10 -4 S cm -1 lithiated,~ 10 -3 S cm -1 delithiated) 37 , high lithium-ion diffusivity ( ∼2 × 10 -10 cm 2 s -1 ) 37 and commercial availability. 36 The NCA cathode was formulated with 95% of NCA, 1% of carbon, and 4% of the CP with an areal mass loading of approximately 14 mg cm -2 to reflect a commercially-relevant electrode composition that maximized the fraction of active materials. With the goal of increasing the specific energy, the mass of all the additional electrode and cell 3.0 3.5 4.0 4.5 -20 0 20 3.0 3.5 4.0 -100 0 100 3.0 3.5 4.0 -20 0 20 10 -5 10 -3 10 -1 3.0 3.5 4.0 10 -9 10 -7 Current ( A) Potential (V vs Li/Li + ) 4.0 V 4.3 V 4.1 V 4.4 V 4.2 V 4.5 V Current ( A) Potential (V vs Li/Li + ) 100 mV s -1 100 mV s -1 100 mV s -1 100 mV s -1 100 mV s -1 Current ( A) Potential (V vs Li/Li + ) 1st Cycle 60th Cycle 20th Cycle 80th Cycle 40th Cycle 100th Cycle (d) (c) (b) ionic P3HT PProDOT-Hx2:OE(80:20) PProDOT-Hx2 Conductivity (S cm -1 ) electronic (a) Potential (V vs Li/Li + ) P3HT PProDOT-Hx2:OE(80:20) PProDOT-Hx2 134 components was reduced significantly, viz., a negative/positive electrode capacity ratio (N/P) of ≈ 3 with metallic lithium as the negative electrode, and an electrolyte to cathode capacity ratio (E/C) of ≈ 3 g (Ah) -1 (Table 4.1). The properties of these Li-NCA-CP cells with this mass-efficient formulation were compared with cells of the same electrode composition and loading but with excess electrolyte and excess anode material (lithium) denoted as “non-limited” cells. The anode and the electrolyte contributed to just 28% in the mass-efficient design whereas it was 75% in the non-limited type cells (Figure 4.2a). The mass-efficient and non-limiting formulations were also assembled with the conventional PVDF binder for comparison. Table 4.1 Cell parameters/performance for the high energy NCA-lithium metal cells with π-conjugated polymers. Cell Parameters Li-NCA- PProDOT-Hx 2 Li-NCA- (Hex:OE)(80:20) Li-NCA- P3HT Li-NCA- PVDF NCA electrode Active Material mass fraction 95% 95% 95% 95% Conductive Polymer mass fraction 4% 4% 4% 4% Areal Loading [mg cm -2 ] 14.2 14.2 14.2 14.5 Electrode Thickness [μm] 41 41 41 46 Al foil thickness [μm] 15 15 15 15 Reversible Capacity [mAh g -1 ] 160 160 160 160 Li electrode Negative/Positive ratio N/P 3.3 3.3 3.1 3.2 Areal Capacity [mAh cm -2 ] 7.5 7.5 7.0 7.5 Specific Capacity [mAh g -1 ] 3860 3860 3860 3860 Electrolyte Electrolyte to Capacity Ratio E/C 3.1 3.1 3.1 3.0 Weight [mg] 10.8 10.8 9.6 10.8 Separator Thickness [μm] 25 25 25 25 Capacity Discharge Capacity at C/2 [mAh g -1 ] 176 158 163 149 Discharge Capacity at C/10 [mAh g -1 ] 194 188 186 174 Specific Energy Specific Energy C/2 [Wh kg -1 ] 323 271 294 250 Specific Energy C/10 [Wh kg -1 ] 359 336 336 312 135 Figure 4.2 Specific energy at the component level of Li-NCA-π-conjugated polymer cells under mass-efficient and non-limited conditions. a) Mass distribution of cell components, b) Specific energy at the cell-level as a function of cycle number for Li-NCA-PProDOT-Hx 2, Li-NCA- (Hex:OE)(80:20), Li-NCA-P3HT, and Li-NCA-PVDF cells with an areal loading of 14 mg cm -2 under mass-efficient and non-limited conditions discharged at C/2 and charged at C/5. NCA electrodes were fabricated with a minimal amount of carbon additive (1 wt%). Two formation cycles at C/10 were conducted before cycling at 23 °C. Non-limited cells had an excess of lithium and electrolyte with an approximate N/P ratio of 45 and an E/C ratio of 17.2. In contrast, mass-efficient cells were assembled with a minimum amount of metallic lithium and electrolyte with an approximate N/P ratio of 3.3 (≈ 3 mg of Li/cell) and E/C ratio of 3.1 (≈ 9-13 µL/cell). Conventional carbonate/LiPF 6 electrolyte was utilized. c)-e) Corresponding galvanostatic charge-discharge curves for the mass-efficient cells in b). The specific energy for the Li-NCA-CP cells of the mass-efficient type with Li-NCA-PProDOT-Hx 2, Li-NCA-(Hex:OE)(80:20), Li-NCA-P3HT, and Li-NCA-PVDF cells reached 359, 336, 336, and 312 Wh kg -1 , respectively, delivering over 3 times the specific energy of their non-limited counterparts (Figure 4.2b). 136 The CPs not only provide a higher mixed (electronic and ionic) conductivity and reversibility 14,16,24 but can also form a protective thin film over the cathode particles. 14–16 We have found that the replacement of the conventional PVDF binder with CPs reduced the electrode resistance and the growth of the solid electrolyte interphase (SEI) on the NCA particles. The benefit of the mixed conductivity exhibited by the CP electrodes became evident as the discharge rate is increased; at the C/2 rate, cells of the mass-efficient type with NCA-PProDOT-Hx 2, NCA-(Hex:OE)(80:20), NCA- P3HT and NCA-PVDF cathodes delivered an impressive initial specific energy of 323, 271, 294, 250 Wh kg -1 respectively and the average capacity decay per cycle after the formation cycles was 1.7, 1.5, 2.2, and 3.1 %, respectively. However, for the cells of the non-limited type the specific energy ranged from 90 to 70 Wh kg -1 although the cathodes exhibited higher overall values of specific capacity than the mass-efficient type (Figure 4.3-4.6). While the high mass fractions of the inactive components in the non-limited cells ensured good utilization of the material, low electrode resistance and longer cycle life, the specific energy of the cell was reduced in proportion to the mass fraction of the inactive components in the electrodes. Thus, the specific energy of the cell was low because of the large mass of the inactive cell components. Unlike the mass-efficient cells, the capacity reduction with increasing rate of discharge in the non-limited cells was not steep. Thus, there is a trade-off between the mass fraction of the inactive materials, their role in enhancing the performance, and the achievable specific energy. In this design optimization of the N/P ratio, the active material loading, porosity/tortuosity of the electrode, electrode thickness, and the electrolyte to capacity (E/C) ratio are the crucial parameters. 30,32,38– 40 137 Figure 4.3. Galvanostatic charge-discharge curves as a function of cycling for the Li-NCA-PVDF cell under mass-efficient conditions charged at C/5 and discharged at C/2. Two formation cycles were carried at C/10 before cycling at 25 °C. 138 Figure 4.4 Corresponding capacity as function of cycle number for the mass-efficient cells in Figure 1a. Specific capacity, and coulombic efficiency as function of cycle number for a) Li-NCA- PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li-NCA-P3HT, and d) Li-NCA-PVDF cells under mass- efficient conditions with an areal loading of approximately 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. Figure 4.5 Corresponding galvanostatic charge-discharge curves for the non-limited cells in Figure 1a. Specific energy at the cell level for a) Li-NCA-PProDOT-Hx 2 cell, b) Li-NCA- (Hex:OE)(80:20) cell, c) Li-NCA-P3HT cell, and d) Li-NCA-PVDF cell under non-limited conditions with an areal loading of 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. 139 Figure 4.6 Corresponding capacity as function of cycle number for the non-limited cells in Figure 1a. Specific capacity and coulombic efficiency as function of cycle number for a) Li-NCA-PProDOT- Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li-NCA-P3HT, and d) Li-NCA-PVDF cells under non-limited conditions with an areal loading of approximately 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. Furthermore, GCD cycling of solely the CPs (without NCA) under lean electrolyte conditions over 100 cycles in the potential range of 2.7 V to 4.2 V vs Li + /Li exhibited a capacity retention of 98.3, 99.8, and 99.5% for PProDOT-Hx 2, (Hex:OE)(80:20), and P3HT respectively, confirming the CPs excellent cyclability and stability (Figure 4.7). 140 Figure 4.7 Galvanostatic characterization of the conducting polymers over the potential window of 2.7 V to 4.2 V vs Li + /Li. a) Percentage capacity retention and coulombic efficiency as a function of cycle number for PProDOT-Hx 2, (Hex:OE)(80:20), and P3HT in a half-cell configuration with a conducting polymer loading of approximately 0.65 mg cm -2 under limited electrolyte conditions with an approximate electrolyte to polymer ratio of 9 µL per mg of polymer, discharged and charged at 0.26 mA cm 2 b)-c) corresponding galvanostatic charge-discharge curves for the cells in a. To prepare the conducting polymer electrodes a slurry composed of conducting polymer and carbon (Super P) in a weight ratio of 40:60 was prepared in OCDB. The slurry was coated onto Al foil using the doctor blading, followed by vacuum drying at 110 °C for two hours. The electrodes were punched into 1.539 cm 2 discs. Understanding the impact of CPs on the cell’s internal processes. We have measured the impedance response for all the non-limited and mass-efficient cells (Figure 4.2b) as a function of state of charge (SOC) during charge and discharge for the 1 st cycle (after formation) and the 26 th cycle. The percentage SOC during charge is defined as the percentage of the charge capacity until the cut-off of 4.2 V vs. Li + /Li. The percentage SOC during discharge is defined as the percentage of the discharge capacity delivered up to a cut-off voltage 141 of 2.7 V vs Li + /Li. On the first galvanostatic charge-discharge (GCD) curves (Figure 4.8a,b) the yellow circles indicate the various SOC values at which the electrochemical impedance (EIS) was measured as a function of frequency for the mass-efficient Li-NCA-PProDOT-Hx 2 and Li-NCA-PVDF cells. We observed a higher overall polarization for the PVDF cell compared to the PProDOT-Hx 2 cell due to the mixed electronic and ionic conductivity that the conducting polymer provides over the insulating PVDF binder. The EIS response was fitted to a Randles-type equivalent electric circuit (EEC) model with an added Voigt element (Figure 4.8c). The EEC model is composed of the ohmic resistance of electrode and electrolyte (R 1), a Voigt element representing the resistance and capacitance of the SEI layer at the electrodes (R 2 and C 1), the charge-transfer resistance (R 3) in parallel with the double layer capacitance (C 2) in series with a semi-infinite Warburg element (W) representing the diffusion processes associated with the electrochemical intercalation reaction. The corresponding impedance measurements at various SOCs during charge are presented as colored spheres in the two 3D-Nyquist plots (Figure 4.8d,e) along with the black lines that correspond to the simulation fit of the EEC model. The 3D-Nyquist plots for all mass- efficient and non-limited CPs cells are shown in Figures 4.9-4.16. Mass-efficient conditions Mass-efficient conditions a) b) c) d) e) Figure 4.8 Impedance characterization for the cells. Galvanostatic charge-discharge profiles for 142 the a) Li-NCA-PProDOT-Hx 2 and b) Li-NCA-PVDF “mass-efficient” cells at C/20 with a charging cut- off of 4.2 V vs Li + /Li. EIS was measured as function of SOC at OCV (denoted with yellow circles). c) equivalent electrical circuit used for fitting the impedance response. Corresponding three- dimensional Nyquist plots of the experimental (denoted with color spheres) and fitted (denoted with black lines) impedance response as function of SOC during charge for the d) Li-NCA- PProDOT-Hx 2 cell and e) Li-NCA-PVDF cells in a and b. Figure 4.9. Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- PProDOT-Hx 2 cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 143 Figure 4.10Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- PProDOT-Hx 2 cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 144 Figure 4.11 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- (Hex:OE)(80:20) cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 145 Figure 4.12 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- (Hex:OE)(80:20) cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 146 Figure 4.13 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- P3HT cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 147 Figure 4.14 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- P3HT cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 148 Figure 4.15 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- PVDF cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . 149 Figure 4.16 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA- PVDF cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . The values of impedance for Li-NCA-PProDOT-Hx 2 cells were 3 to 5 times lower than those for the Li-NCA-PVDF cells. The impedance decreased at higher cell voltage values, for both the mass- efficient and non-limited cells. The impedance elements fitted to the EEC (Figure 4.8c) were found to vary with the SOC (Figure 4.17). Tables 4.2 and 4.3 show the best fit of the EEC elements at mid-SOC for all the Li-NCA-π-conjugated polymer cells under mass-efficient and non-limited conditions. 150 Figure 4.17 Impedance characterization for the Li-NCA-π-conjugated polymer cells under mass- efficient and non-limited conditions. Evolution of the a) electrolyte resistance R 1, b) SEI resistance R 2, c) charge transfer resistance R 3, d) SEI capacitance C 1 e), double layer capacitance C 2, and f) Warburg coefficient S as a function of SOC for the 1 st cycle and the 26 th cycle during charge and discharge for Li-NCA-PProDOT-Hx 2 (denoted by blue and turquoise connected dots), Li-NCA- (Hex:OE)(80:20) (olive and green), Li-NCA-P3HT (purple and magenta), and Li-NCA-PVDF (red and pink). Relative fitting errors are indicated with error bars. The equivalent electrical circuit utilized for the impedance fitting is shown in Figure 4.8c. 0 20 40 60 80 100 0 20 40 60 80 100 120 0 20 40 60 80 100 0 20 40 60 80 100 120 100 80 60 40 20 0 0 20 40 60 80 100 120 100 80 60 40 20 0 0 20 40 60 80 100 120 0 20 40 60 80 100 0 20 40 60 80 100 120 140 160 0 20 40 60 80 100 0 20 40 60 80 100 120 140 160 100 80 60 40 20 0 0 20 40 60 80 100 100 80 60 40 20 0 0 20 40 60 80 100 0 20 40 60 80 100 0 20 40 60 80 100 120 140 0 20 40 60 80 100 0 20 40 60 80 100 120 140 100 80 60 40 20 0 0 20 40 60 80 100 120 140 100 80 60 40 20 0 0 20 40 60 80 100 120 140 0 20 40 60 80 100 0 10 20 30 40 50 0 20 40 60 80 100 0 10 20 30 40 50 100 80 60 40 20 0 0 10 20 30 40 50 100 80 60 40 20 0 0 10 20 30 40 50 0 20 40 60 80 100 1E-6 1E-5 1E-4 0.001 0.01 0.1 0 20 40 60 80 100 1E-6 1E-5 1E-4 0.001 0.01 0.1 100 80 60 40 20 1E-6 1E-5 1E-4 0.001 0.01 0.1 100 80 60 40 20 1E-6 1E-5 1E-4 0.001 0.01 0.1 0 20 40 60 80 100 0.0E+0 6.0E-5 1.2E-4 1.8E-4 2.4E-4 0 20 40 60 80 100 0.0E+0 6.0E-5 1.2E-4 1.8E-4 2.4E-4 100 80 60 40 20 0 0.0E+0 1.0E-4 2.0E-4 3.0E-4 4.0E-4 100 80 60 40 20 0 0.0E+0 1.0E-4 2.0E-4 3.0E-4 4.0E-4 1 st Cycle Charge 26 th Cycle Charge Discharge Discharge R 3 (ohm) % state of charge 1 st Cycle Charge Discharge Discharge b) 26 th Cycle Charge % state of charge % state of charge 1 st Cycle Charge R 2 (ohm) c) R 1 (ohm) 26 th Cycle Charge % state of charge % state of charge % state of charge Discharge Discharge NCA-PVDF (Mass-efficient) NCA-PProDOT-Hx 2 (Mass-efficient) NCA-(Hex:OE)(80:20) (Mass-efficient) NCA-P3HT (Mass-efficient) NCA-PVDF NCA-PProDOT-Hx 2 NCA-(Hex:OE)(80:20) NCA-P3HT 1 st Cycle Charge d) 26 th Cycle Charge Discharge Discharge 1 st Cycle Charge 26 th Cycle Charge Discharge C 1 (F) C 2 (F) S 1 (ohm s -0.5 ) Discharge 1 st Cycle Charge Discharge f) e) a) 26 th Cycle Charge Discharge 151 Table 4.2. Fitting parameters and their corresponding relative error for the 1 st cycle of the Li- NCA- π-conjugated polymers cells at approximately 4.0 V vs Li + /Li. R 1 [Ohm] Error % C 1 [F] Error % R 2 [Ohm] Error % C 2 [F] Error % R 3 [Ohm] Error % S 1 [Ohm S 0.5 ] Error % Li-NCA-PProDOT-Hx 2 (Mass-efficient) 6.1 1.7 1.9E-03 4.0E-03 1.5 0.6 3.8E-06 4.1E-06 6.4 0.6 1.4 0.4 Li-NCA-PProDOT-Hx 2 6.3 1.4 3.7E-03 1.0E-02 1.2 0.7 1.2E-05 1.9E-05 3.0 0.7 1.2 0.4 Li-NCA-(Hex:OE)(80:20) (Mass-efficient) 5.2 1.2 2.9E-04 3.1E-04 5.2 1.7 2.3E-06 3.8E-07 24.2 2.0 3.9 0.3 Li-NCA-(Hex:OE)(80:20) 6.6 1.4 1.6E-03 2.2E-03 2.4 0.7 4.1E-06 1.8E-06 9.8 0.4 1.9 0.4 Li-NCA-P3HT (Mass- efficient) 6.4 1.4 7.2E-03 1.3E-02 1.5 0.5 9.1E-06 8.9E-06 4.1 1.0 1.8 0.5 Li-NCA-P3HT 6.0 0.9 6.3E-03 1.5E-02 1.3 0.6 2.5E-05 1.8E-05 2.6 0.2 1.4 0.5 Li-NCA-PVDF (Mass- efficient) 5.3 0.7 2.3E-06 7.4E-08 47.9 4.1 6.4E-05 3.2E-05 19.7 3.9 4.3 0.3 Li-NCA-PVDF 7.3 0.6 3.0E-06 7.2E-08 38.7 2.7 1.0E-04 3.9E-05 19.3 2.5 4.8 0.3 Table 4.3 Fitting parameters and their corresponding relative error for the 26 th cycle of the Li- NCA- π-conjugated polymers cells at approximately 4.0 V vs Li + /Li. R 1 [Ohm] Error % C 1 [F] Error % R 2 [Ohm] Error % C 2 [F] Error % R 3 [Ohm] Error % S 1 [Ohm S 0.5 ] Error % Li-NCA-PProDOT-Hx 2 (Mass-efficient) 19.6 1.1 1.2E-04 9.1E-05 9.5 2.5 6.5E-06 2.6E-06 9.7 2.9 6.1 0.3 Li-NCA-PProDOT-Hx 2 12.9 1.3 1.6E-04 1.5E-04 6.8 2.0 6.4E-06 4.0E-06 7.8 2.6 4.6 0.3 Li-NCA-(Hex:OE)(80:20) (Mass-efficient) 15.5 1.2 2.0E-06 3.0E-07 31.7 4.9 2.9E-05 1.3E-05 23.6 4.1 7.1 0.3 Li-NCA-(Hex:OE)(80:20) 9.9 1.1 4.0E-06 8.8E-07 17.6 3.8 4.3E-05 1.8E-05 20.9 3.3 9.5 0.3 Li-NCA-P3HT (Mass- efficient) 19.8 1.5 8.8E-05 5.7E-05 9.2 1.7 3.0E-06 1.7E-06 12.1 2.9 6.5 0.3 Li-NCA-P3HT 9.3 1.4 7.3E-05 8.1E-05 6.6 2.2 5.8E-06 8.1E-06 5.8 4.4 6.7 0.2 Li-NCA-PVDF (Mass- efficient) 31.4 1.3 5.8E-06 9.9E-07 80.1 5.9 1.0E-06 1.0E-07 58.1 7.1 8.3 0.2 Li-NCA-PVDF 10.4 0.3 2.9E-06 1.1E-07 55.7 2.0 6.8E-05 7.4E-06 55.9 1.9 9.5 0.4 152 The ohmic resistance R 1 was in the range of 5 to10 ohms for all the freshly-prepared cells and did not vary significantly with SOC. However, after 26 cycles the mass-efficient cells exhibited an increase in their ohmic resistance that we attribute to electrolyte insufficiency. The limited amount of electrolyte in these cells is mainly consumed by reaction at the lithium electrode due to the continuous growth of the SEI during repeated cycling. 41–43 Thus, by avoiding electrolyte consumption at the anode we can ensure a long-cycle life cell with a mass-efficient design. The value of R 1 following cycling for the mass-efficient Li-NCA-PVDF cell (red dots) was about two times larger than that of the Li-NCA-CP cells, indicating a higher rate of electrolyte consumption in the PVDF cells compared to the mass-efficient cells with CPs. This finding is consistent with the protective role of the CPs coated on the NCA particles forming a more electrochemically-stable SEI, preventing the electrolyte consumption at the cathode. 14–16 In stark contrast, in the non- limited cells because of the abundance of electrolyte, we observe only a slight increase in the ohmic resistance after cycling. The value of R 2 for the NCA-PVDF electrodes (red and pink dots) was 4 to 10 times higher than the cells with CPs, confirming the benefit of formation of a stable SEI over the NCA particles with the CP-based binders (Figure 4.17b). However, we observe that after 26 cycles there is a slight increase in R 2 for just the mass-efficient Li-NCA-(Hex:OE)(80:20) cell (olive dots). These results suggest that the differences in electronic conductivity (Figure 4.1d) and the effect of degree of swelling of the polymers became evident when the availability of the electrolyte is restricted. Figure 4.17c shows the changes of R 3, the polarization resistance associated with the electrochemical reaction kinetics at the electrode. For all the polymer binders R 3 decreased with increasing SOC. Further, the CP containing cells exhibited 70% lower values compared to the 153 PVDF cells. The polarization resistance, R 3, is related to the exchange current density and electrochemically active surface area by R 3=RT (nFA oi o) -1 , where i o is the exchange current density, A o is the active surface area, F the Faraday constant, n is the number of electrons transferred, T is the absolute temperature, and R is the ideal gas constant. Thus, the decrease in R 3 with increasing SOC is consistent with the increase in i o observed with the charging of NCA. 44,45,36 However, after 26 cycles there is a significant increase in R 3 for the PVDF cells indicating that the active area (A o) of the electrode available for electrochemical intercalation of lithium decreased as a function of cycling, most likely due to the continuous growth of a thicker/insulating SEI layer on the cathode. The lower values of R 3 for the CP cathodes in comparison to PVDF-containing cathodes is attributed to the electronic conductivity of the conjugated polymers that results in an increase of the electrochemically accessible surface area. We attribute the differences in R 3 observed among the various Li-NCA-CP cells, to the porosity variations of the electrodes. From the SEM studies on the electrodes, we found that all electrodes consisted of macro-spherical agglomerates of NCA with a diameter ranging from 4-10 µm, where each sphere is comprised by smaller NCA particles (100 to 500 nm in size) creating a hierarchical morphology (Figure 4.18a-d). 154 Figure 4.18. SEM characterization of the NCA-π-conjugated polymer electrodes. SEM of the electrode composite films as prepared with a composition of 95 wt% NCA, 4 wt% conducting polymer and 1 wt% of carbon at various scales ranges from 100 microns to 500 nm for a) NCA- PProDOT-Hx 2, b) NCA-(Hex:OE)(80:20), c) NCA-P3HT and d) NCA-PVDF. Despite the same mass fraction of the various constituents and the uniform distribution of electrode components, the different CPs gave rise to different values of final porosity and pore sizes. The primary (macro) and secondary (meso) pore distribution of the NCA electrodes, analyzed using the ImageJ software (Figure 4.19-4.26), showed that NCA-P3HT electrodes had a larger fraction of macro and meso pores compared to the other CPs and PVDF-based electrodes. The effective porosity (as measured by the imbibition/Archimedes’ method, Table 4.4) for the 155 NCA-P3HT, NCA-PProDOT-Hx 2, NCA-(Hex:OE)(80:20), and NCA-PVDF electrodes was 80, 76, 68, and 61% respectively. Using Energy-dispersive X-ray Fluorescence Spectroscopy (EDS) we confirmed the uniform distribution of the CPs (Figure 4.27-4.34). Therefore, we conclude that the lower value of R 3 for PProDOT-Hx 2 (blue and turquoise dots) and P3HT (purple and magenta dots) was attributable to the higher porosity of these electrodes (Figure 4.17c). Figure 4.19 Macro pores characterization of the NCA-PProDOT-Hx 2 electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. Figure 4.20 Meso pores characterization of the NCA-PProDOT-Hx 2 electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. 156 Figure 4.21 Macro pores characterization of the NCA-(Hex:OE)(80:20) electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. Figure 4.22 Meso pores characterization of the NCA-(Hex:OE)(80:20) electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. Figure 4.23 Macro pores characterization of the NCA-P3HT electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. 157 Figure 4.24 Meso pores characterization of the NCA-P3HT electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. Figure 4.25 Macro pores characterization of the NCA-PVDF electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores and c) bare outlines form the particle analysis. Figure 4.26 Meso pores characterization of the NCA-PVDF electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores and c) bare outlines form the particle analysis. 158 Energy-dispersive X-ray spectroscopy (EDS) mapping. Electrode particles are made of NCA (Ni,Co,O,Al) (Figure 4.27e,g,f,h, 4.29e,g,f,h, 4.31e,g,f,h, 4.33e,g,f,h). Elemental mapping of sulfur showed the uniform distribution of the π-CPs containing thiophene groups on the NCA particles (Figure 4.27d, 4.29d, 4.31d). The fluorine mapping confirmed that the PVDF binder was also uniformly distributed (Figure 4.33d). Figure 4.27 EDS Elemental mapping of NCA-PProDOT-Hx 2 electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% PProDOT-Hx 2, and 1% Carbon, b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. 159 Figure 4.28 EDS spectrum of NCA-PProDOT-Hx 2 electrode surface. Figure 4.29 EDS Elemental mapping of NCA-(Hex:OE)(80:20) electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% (Hex:OE)(80:20), and 1% Carbon b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. 160 Figure 4.30 EDS spectrum of NCA-(Hex:OE)(80:20) electrode surface. Figure 4.31 EDS Elemental mapping of NCA-P3HT electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% P3HT, and 1% Carbon, b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. 161 Figure 4.32 EDS spectrum of NCA-P3HT electrode surface Figure 4.33 EDS Elemental mapping of NCA-PVDF electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% PVDF, and 1% Carbon b) all combined elements mapping, c) carbon mapping, d) fluorine mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. 162 Figure 4.34 EDS spectrum of the NCA-PVDF electrode surface. Figure 4.35 Pore parameters of the NCA-π-conjugated polymers electrodes from SEM. a) Estimation of the surface % area covered by macro pores and b) meso pores for the NCA-π- conjugated polymers electrodes. c) Mean Feret Diameter of the macro pores and d) meso pores NCA-π-conjugated polymers electrodes. 163 Determination of the effective Porosity (Imbibition/Archimedes’ Method) Table 4.4 Effective Porosity by Imbibition/Archimedes’ Method for the NCA- π-conjugated polymers electrodes. Porosity was determined using the average value of triplicated measurements using isopropyl alcohol as the fluid. Effective Porosity Standard deviation (SDV) Standard Error (SEM) NCA-PProDOT-Hx 2 76% 0.4% 0.2% NCA-(Hex:OE)(80:20) 68% 0.8% 0.4% NCA-P3HT 80% 1.7% 0.9% NCA-PVDF 61% 0.5% 0.3% During the 1 st cycle, CP cells showed C 1 values of 100-1000 times that of PVDF cells. The higher value of C 1 was attributed to the ability of the CP layers to store charge during the polymer doping process, not possible with the PVDF cells (Figure 4.17d). The value of C 1 represents the faradaic pseudo-capacitance of the conducting polymer with a value of about 0.01 F/cm 2 consistent with the polymer being present only to the extent of 5% of the total electrode mass. Upon cycling the increasing thickness of the SEI on the cathode leads to a decrease of the values of C 1 for all the polymers. Changes to the active surface area for the charge-transfer reactions is reflected in C 2. During the first cycle PVDF electrodes tended to have higher values of C 2 attributed to higher active area due to a more active SEI formation (Figure 4.17e). For all the cells, we found the Warburg coefficient, S 1, increased with decreasing SOC values (Figure 4.17f). S 1 is inversely related to the diffusion coefficient of the lithium ion in the intercalation material. At lower SOCs, when NCA is highly lithiated, only a limited number of lattice sites are available for lithium-ion diffusion. With progressive de-lithiation of the NCA, more pathways become available for 164 diffusion leading to an increase in the apparent diffusion coefficient of the lithium ion. Additionally, towards the end of discharge, the PVDF cells exhibited a higher value of S 1 compared to the CP electrodes. This observation suggests a higher interfacial area of contact for the mass transport of lithium ions in the CP cells compared to the PVDF cells. We have also characterized the impedance response of LiNi 0.6Mn 0.2Co 0.2O2 (NMC-622) with PProDOT-Hx 2 and PVDF binders (Figures 4.36-4.38). The data shows that electrodes made with PProDOT-Hx 2 exhibited 6 to 7 times lower charge transfer resistance compared to the electrodes made with PVDF confirming the benefit of utilizing conducting polymers with other lithium-ion intercalating materials. Figure 4.36 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. Nyquist plots as function of SOC during charge for a) Li-NMC-622-PVDF cell b) for the Li-MNC-622-PProDOT-Hx 2 cell under mass-efficient conditions with an approximate areal loading of 13 mg cm 2 . Furthermore, we have fitted the impedance response a function of SOC based on the ECC from Figure 4.8c to obtain the charge transfer resistance of the electrodes. Figure 4.37 shows the 165 charge transfer resistance comparison as function of SOC for cells made with NCA-PVDF, NCA- PProDOT-Hx 2, NMC-622-PVDF and NMC-622-PProDOT-Hx 2. We can observe the electrodes made with CP exhibited 6 to 7 times lower charge transfer resistance compared to the electrodes made with PVDF confirming the benefit of utilizing conducting polymers. Figure 4.37. Charge transfer resistance as function of SOC obtained by ECC fitting from Figure 4.10c for NCA and NMC-622 based electrodes with PVDF and PProDOT-Hx 2. We assembled a NMC-622-PProDOT-Hx 2 cell under non-limited conditions, moderate loading, and with 1% carbon to demonstrate the benefits of the CPs with other lithium-ion intercalating materials. Figure 4.38 shows the rate capability testing of NCM-622-PProDOT-Hx 2. At extremely high rates (4C) the cell was able to deliver a remarkable specific capacity of 130 mAh g -1 . 166 Figure 4.38 a) Galvanostatic charge-discharge curves as function of C-rate for a Li-NMC-622- PProDOT-Hx 2 under non-limited conditions with an approximate mass loading of 5.5 mg cm -2 . Two formation cycles where carried at C/20. C-rate was calculated based on the reversible capacity of NMC-622 being 1C = 180 mA g -1 . b), corresponding specific capacity at the cell-level as function of cycle number at the different rates from the discharge curves plotted in a). Phase Equilibria and Polarization via Differential capacity analysis To analyze cell degradation (capacity loss) as a function of cycle number, we determined the differential capacity (dQ dV -1 vs V) from the GCD data for all the Li-NCA-CP cells and Li-NCA-PVDF cells. Differential capacity is indicative of the phase equilibria and changes to the internal resistance of NCA during de-lithiation and lithiation 46–49 . The shift in peak position and changes to peak heights can provide insight into the structure evolution of NCA during lithium ion intercalation/deintercalation. 48,50,51 During charging in the first cycle of formation for Li-NCA- PProDOT-Hx 2 cell showed (Figure 4.39a, red line) a peak between 3.55 and 3.65 V vs Li + /Li attributed to the phase coexistence of a hexagonal (H1) and monoclinic (M) lattice (H1+M), followed by a second peak between 3.95 and 4.05 V associated with the coexistence of M with a second hexagonal (H2) lattice, (M+H2) 47,52 . During discharge we observed a peak between 4.2 167 and 4.12 V attributed to the phase equilibria of a third hexagonal lattice (H3) with H2, (H2 + H3), followed by peaks for the coexistence of M and H2 between 4.05 and 3.85 V, and the (H1 + M) equilibria between 3.85 and 3.55 V. The results of differential capacity analysis indicate that the NCA particles coated with CPs retain their phase composition over cycling, suggesting a protective role of the CPs to mitigate the material transformation of the NCA particles 45,46,53,54 . In situ XRD studies combined with XPS would be necessary to confirm this hypothesis. We observed that as a function of cycling, all cells in Figure 4.39 exhibited a significant shifting of the discharge peak for H1 + M phase equilibria to lower potentials that can be correlated with a higher overall impedance (Figure 4.9-4.16). This shift is most significant with the mass-efficient PVDF cells (Figure 4.39d). We found that the non-limited cells maintained their phase transitions over cycling except for the non-limited-PVDF cell that exhibited a notable decrease in the height of the peaks and shift of peak positions consistent with the observed capacity fade and higher impedance of the PVDF containing cells. We also found that the mass-efficient-type cells with PProDOT-Hx 2 (Figure 4.39a) and (Hex:OE)(80:20) (Figure 4.39b) cells maintained their discharge peaks over cycling. However, with the P3HT cells (Figure 4.39c) their charge and discharge peaks became smaller with cycling, consistent with a higher rate of capacity fade and impedance increase (Figure 4.17a) for these cells. While all the CPs appear superior to PVDF in ensuring phase stability, PProDOT-Hx 2 and (Hex:OE)(80:20) seem to provide beneficial properties over P3HT. Thus, under lean electrolyte conditions, regions of the P3HT electrode could become devoid of electrolyte due to its higher porosity (Figure 4.19-4.35 and Table 4.4). Such a situation would lead to insufficient electrolyte availability to all parts of the electrode during cycling, and consequently a lower utilization of the active material. 168 Figure 4.39 Differential capacity vs voltage analysis (dQ dV -1 vs V) for Li-NCA-π-conjugated polymer cells. dQ dV -1 vs V curves measured as a function of number of cycles with a C/2 rate for discharge and C/5 for charging including the first formation cycle at C/10 (denoted with a red curve) under mass-efficient and non-limited conditions between 2.7 and 4.2 V vs Li + /Li for a) Li- NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20) cells, c) Li-NCA-P3HT cells, and d) Li-NCA-PVDF cells. Rate Capability under mass-efficient and non-limited conditions We studied the discharge specific energy and specific capacity as function of cycle number and discharge rate for the Li-NCA-CP and Li-NCA-PVDF cells under mass-efficient and non-limited conditions (Figure 4.40a,b). Figures 4.40c-4.40f are the corresponding GCD curves for the mass- efficient cells from Figure 7a. The GCD curves for the non-limited cells in Figure 7b are provided in Figure 4.41 For all the mass-efficient cells (Figure 4.40a) we found a comparable utilization of the cathode materials during formation. However, as the C-rate is increased the beneficial role of the CPs and 169 the electrode morphology became apparent. Even at C/2 we observed a dramatic decrease in the capacity for the Li-NCA-PVDF cell. Two main factors contribute to this drop in capacity; a lower electrode conductivity, and a more aggressive consumption of the limited amount of electrolyte during the SEI build-up that is consistent with the impedance results (Figure 4.17a,b). At 1C we observe that the PProDOT-Hx 2 and (Hex:OE)(80:20) cells retain an impressive specific energy of ≈ 200 Wh kg -1 followed by P3HT (150 Wh kg -1 ) while the PVDF cells stay at a low value of 50 Wh kg -1 . Assuming that all cells have a similar consumption of electrolyte at the lithium electrode, the differences among the CP cells can be attributed to the mixed ionic/electronic conductivity and the morphological differences of the NCA electrode. Specifically, higher ionic conductivity is an advantage at high rates under electrolyte-lean conditions. Thus, the CP binders are particularly suited for the sustaining electronic and ion transport as the electrolyte is consumed, a feature and benefit that is unavailable with PVDF as a binder. While porosity is usually an advantage, in the NCA-P3HT electrode the higher porosity and wider macro pores could be detrimental because of the limited amount of available electrolyte. Consequently, there is an early onset of impedance increase with the mass-efficient P3HT cells especially when discharged at the higher rates (C/2-1C). The impact of re-distribution and consumption of the electrolyte due to the higher porosity of the P3HT cell can also be observed in the higher ohmic resistance after 26 cycles (Figure 4.17a), and the peak shifts following cycling in the differential capacity analysis (Figure 4.39c). At the 2C rate these effects are amplified further. The cells with (Hex:OE)(80:20) showed the highest utilization and retention of all the cells due to its enhanced ionic conductivity and lower porosity. These limitations were not observed when the cells were assembled with excess of electrolyte (Figure 4.40b and 4.40b). Cells made of the CPs with higher 170 electronic conductivity (PProDOT-Hx 2 and P3HT) retained a higher capacity as rate was increased. Therefore, we have identified that the benefits of low electrode porosity and high polymer ionic conductivity are the critical parameters affecting rate capability under mass-efficient conditions, whereas the electronic conductivity is the main factor under non-limited conditions. Figure 4.40 Rate Capability for Li-NCA-π-conjugated polymer cells under mass-efficient and non- limited conditions. a) Specific energy at the cell-level as a function of cycling with increments in the discharge rate with a constant charging rate (C/5) and two formation cycles charged/discharged at C/10 for the Li-NCA-π-conjugated polymer cells under mass-efficient conditions. Mass-efficient cells had an approximate N/P ratio of 3 and an E/C ratio of 3. b) Specific capacity as a function of cycling at various discharge rates with a constant charging rate and two formation cycles for the Li-NCA-π-conjugated polymer cells under non-limited conditions. Corresponding galvanostatic discharge-charge curves in a) for c) Li-NCA-PProDOT-Hx 2, d) Li-NCA- (Hex:OE)(80:20), e) Li-NCA-P3HT, and f) Li-NCA-PVDF cells 171 Figure 4.41 Rate Capability for Li-NCA-π-conjugated polymer cells under non-limited conditions. Corresponding galvanostatic discharge-charge curves from Figure 6b for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li-NCA-P3HT, and d) Li-NCA-PVDF cells in non-limited conditions. Conclusions We show that lithium metal batteries with CPs as cathode binders can deliver a high specific energy (at the cell component level excluding packaging) of 320 Wh kg -1 at C/2 and 208 Wh kg -1 at 1C. The high-specific energy/high-rate performance is achieved by reducing the weight of the inactive components enabled by the benefits of higher electronic/ionic conductivity and protection of the intercalating particles by the π-conjugated polymers in the cathode. The benefits of polymer structure and conductivity was exemplified in the behavior of PProDOT-Hx 2, (Hex:OE)(80:20), and P3HT at higher discharge rates. The electrode morphology characterization, 172 rate capability testing, EIS analysis, and differential capacity analysis indicated that under the conditions of limited electrolyte availability, polymers with higher-ionic conductivity such as (Hex:OE)(80:20) are preferred for achieving higher specific energy and higher rate. Under these mass-efficient conditions, the electrodes of higher porosity with wide macro pores without the protection of the NCA particles leads to poor cell performance. In contrast, for the cells where the electrolyte content is not limited, the pore size and porosity differences of the electrodes did not have a significant contribution to the rate capability whereas the higher electronic/ionic conductivity were important. The rigorous attention to mass-efficient cell designs and control of parameters such as the N/P and E/C ratios during cell fabrication facilitate the translation of our findings to practical designs. With the progress in the design of less reactive electrolytes and protected lithium anodes, the conducting-polymer-based cathodes provide a new pathway to lithium-ion batteries with high specific energy, high rate capability, and long cycle life. References 1. Elizalde-Segovia, R. et al. Understanding the Role of π-Conjugated Polymers as Binders in Enabling Designs for High-Energy/High-Rate Lithium Metal Batteries. J. Electrochem. Soc. 168, 110541 (2021). 2. Choi, S. & Wang, G. Advanced Lithium-Ion Batteries for Practical Applications: Technology, Development, and Future Perspectives. Adv. Mater. Technol. 3, 1700376 (2018). 3. Placke, T., Kloepsch, R., Dühnen, S. & Winter, M. Lithium ion, lithium metal, and alternative rechargeable battery technologies: the odyssey for high energy density. J Solid State Electrochem 21, 1939–1964 (2017). 4. Zeng, X. et al. Commercialization of Lithium Battery Technologies for Electric Vehicles. Adv. Energy Mater. 9, 1900161 (2019). 5. Zubi, G., Dufo-López, R., Carvalho, M. & Pasaoglu, G. The lithium-ion battery: State of the art and future perspectives. Renewable and Sustainable Energy Reviews 89, 292–308 (2018). 6. Berg, E. J., Villevieille, C., Streich, D., Trabesinger, S. & Novák, P. Rechargeable Batteries: Grasping for the Limits of Chemistry. J. Electrochem. Soc. 162, A2468–A2475 (2015). 173 7. Kim, U.-H. et al. Pushing the limit of layered transition metal oxide cathodes for high-energy density rechargeable Li ion batteries. Energy Environ. Sci. 11, 1271–1279 (2018). 8. Radin, M. D. et al. Narrowing the Gap between Theoretical and Practical Capacities in Li-Ion Layered Oxide Cathode Materials. Adv. Energy Mater. 7, 1602888 (2017). 9. Shreenivasa, L. et al. Scalable chemical approach to prepare crystalline Mn2V2O7 nanoparticles: introducing a new long-term cycling cathode material for lithium-ion battery. J Mater Sci: Mater Electron 31, 19638–19646 (2020). 10. Shreenivasa, L. et al. An introduction of new nanostructured Zn0.29V2O5 cathode material for lithium ion battery: a detailed studies on synthesis, characterization and lithium uptake. Materials Research Express 6, 13 (2019). 11. Sadhasivam, T. et al. High charge acceptance through interface reaction on carbon coated negative electrode for advanced lead-carbon battery system. Electrochimica Acta 295, 367–375 (2019). 12. Nguyen, V. A. & Kuss, C. Review—Conducting Polymer-Based Binders for Lithium-Ion Batteries and Beyond. J. Electrochem. Soc. 167, 065501 (2020). 13. Snook, G. A., Kao, P. & Best, A. S. Conducting-polymer-based supercapacitor devices and electrodes. Journal of Power Sources 196, 1–12 (2011). 14. Lai, C.-H. et al. Application of Poly(3-hexylthiophene-2,5-diyl) as a Protective Coating for High Rate Cathode Materials. Chem. Mater. 30, 2589–2599 (2018). 15. Zhao, Y. et al. A Conductive Binder for High-Performance Sn Electrodes in Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 10, 1672–1677 (2018). 16. Wu, F. et al. Surface Modification of Li-Rich Cathode Materials for Lithium-Ion Batteries with a PEDOT:PSS Conducting Polymer. ACS Appl. Mater. Interfaces 8, 23095–23104 (2016). 17. Chen, S. et al. Critical Parameters for Evaluating Coin Cells and Pouch Cells of Rechargeable Li- Metal Batteries. Joule 3, 1094–1105 (2019). 18. Ren, X. et al. Enabling High-Voltage Lithium-Metal Batteries under Practical Conditions. Joule 3, 1662–1676 (2019). 19. Niu, C. et al. High-energy lithium metal pouch cells with limited anode swelling and long stable cycles. Nat Energy 4, 551–559 (2019). 20. Niu, C. et al. Self-Smoothing Anode for Achieving High-Energy Lithium Metal Batteries under Realistic Conditions. 10. 21. Patel, S. N., Javier, A. E. & Balsara, N. P. Electrochemically Oxidized Electronic and Ionic Conducting Nanostructured Block Copolymers for Lithium Battery Electrodes. ACS Nano 7, 6056–6068 (2013). 22. Higgins, T. M. et al. A Commercial Conducting Polymer as Both Binder and Conductive Additive for Silicon Nanoparticle-Based Lithium-Ion Battery Negative Electrodes. ACS Nano 10, 3702–3713 (2016). 23. Eliseeva, S. N., Levin, O. V., Tolstopyatova, E. G., Alekseeva, E. V. & Kondratiev, V. V. Effect of addition of a conducting polymer on the properties of the LiFePO4-based cathode material for lithium-ion batteries. Russ J Appl Chem 88, 1146–1149 (2015). 174 24. Lee, J. & Choi, W. Surface Modification of Over-Lithiated Layered Oxides with PEDOT:PSS Conducting Polymer in Lithium-Ion Batteries. J. Electrochem. Soc. 162, A743–A748 (2015). 25. Zhong, H. et al. Carboxymethyl chitosan/conducting polymer as water-soluble composite binder for LiFePO4 cathode in lithium ion batteries. Journal of Power Sources 336, 107–114 (2016). 26. Kim, J.-M. et al. Conducting Polymer-Skinned Electroactive Materials of Lithium-Ion Batteries: Ready for Monocomponent Electrodes without Additional Binders and Conductive Agents. ACS Appl. Mater. Interfaces 6, 12789–12797 (2014). 27. Zeng, W. et al. Enhanced Ion Conductivity in Conducting Polymer Binder for High-Performance Silicon Anodes in Advanced Lithium-Ion Batteries. Adv. Energy Mater. 8, 1702314 (2018). 28. Park, K.-S., Schougaard, S. B. & Goodenough, J. B. Conducting-Polymer/Iron-Redox- Couple Composite Cathodes for Lithium Secondary Batteries. Adv. Mater. 19, 848–851 (2007). 29. Das, P. et al. Dihexyl-Substituted Poly(3,4-Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. 32, 9176–9189 (2020). 30. Cao, Y., Li, M., Lu, J., Liu, J. & Amine, K. Bridging the academic and industrial metrics for next- generation practical batteries. Nat. Nanotechnol. 14, 200–207 (2019). 31. Albertus, P., Babinec, S., Litzelman, S. & Newman, A. Status and challenges in enabling the lithium metal electrode for high-energy and low-cost rechargeable batteries. Nat Energy 3, 16–21 (2018). 32. Liu, J. et al. Pathways for practical high-energy long-cycling lithium metal batteries. Nat Energy 4, 180–186 (2019). 33. Ding, N. et al. Rational design of a high-energy LiNi0.8Co0.15Al0.05O2 cathode for Li-ion batteries. Solid State Ionics 323, 72–77 (2018). 34. Zayat, B., Das, P., Thompson, B. C. & Narayan, S. R. In Situ Measurement of Ionic and Electronic Conductivities of Conductive Polymers as a Function of Electrochemical Doping in Battery Electrolytes. J. Phys. Chem. C acs.jpcc.0c08934 (2021) doi:10.1021/acs.jpcc.0c08934. 35. Xia, Y., Zheng, J., Wang, C. & Gu, M. Designing principle for Ni-rich cathode materials with high energy density for practical applications. Nano Energy 49, 434–452 (2018). 36. Julien, C. M. & Mauger, A. NCA, NCM811, and the Route to Ni-Richer Lithium-Ion Batteries. Energies 13, 6363 (2020). 37. Amin, R., Ravnsbæk, D. B. & Chiang, Y.-M. Characterization of Electronic and Ionic Transport in Li 1- x Ni 0 . 8 Co 0.15 Al 0.05 O 2 (NCA). J. Electrochem. Soc. 162, A1163–A1169 (2015). 38. Zheng, H., Li, J., Song, X., Liu, G. & Battaglia, V. S. A comprehensive understanding of electrode thickness effects on the electrochemical performances of Li-ion battery cathodes. Electrochimica Acta 71, 258–265 (2012). 39. Singh, M., Kaiser, J. & Hahn, H. A systematic study of thick electrodes for high energy lithium ion batteries. Journal of Electroanalytical Chemistry 782, 245–249 (2016). 40. Delattre, B. et al. Impact of Pore Tortuosity on Electrode Kinetics in Lithium Battery Electrodes: Study in Directionally Freeze-Cast LiNi 0.8 Co 0.15 Al 0.05 O 2 (NCA). J. Electrochem. Soc. 165, A388–A395 (2018). 175 41. Lu, D. et al. Failure Mechanism for Fast-Charged Lithium Metal Batteries with Liquid Electrolytes. Adv. Energy Mater. 5, 1400993 (2015). 42. Chen, K.-H. et al. Dead lithium: mass transport effects on voltage, capacity, and failure of lithium metal anodes. J. Mater. Chem. A 5, 11671–11681 (2017). 43. Zhang, S. S. Problem, Status, and Possible Solutions for Lithium Metal Anode of Rechargeable Batteries. ACS Appl. Energy Mater. 1, 910–920 (2018). 44. Zhu, X. et al. Electrochemical impedance study of commercial LiNi0.80Co0.15Al0.05O2 electrodes as a function of state of charge and aging. Electrochimica Acta 287, 10–20 (2018). 45. Tsai, P.-C. et al. Single-particle measurements of electrochemical kinetics in NMC and NCA cathodes for Li-ion batteries. Energy Environ. Sci. 11, 860–871 (2018). 46. Li, J. et al. Dependence of Cell Failure on Cut-Off Voltage Ranges and Observation of Kinetic Hindrance in LiNi 0.8 Co 0.15 Al 0.05 O 2. J. Electrochem. Soc. 165, A2682–A2695 (2018). 47. Zhu, X.-J. et al. Preparation and characterization of LiNi0.80Co0.20– x Al x O2 as cathode materials for lithium ion batteries. J Electroceram 17, 645–649 (2006). 48. Nam, G. W. et al. Capacity Fading of Ni-Rich NCA Cathodes: Effect of Microcracking Extent. ACS Energy Lett. 4, 2995–3001 (2019). 49. Park, K.-J. et al. Degradation Mechanism of Ni-Enriched NCA Cathode for Lithium Batteries: Are Microcracks Really Critical? ACS Energy Lett. 4, 1394–1400 (2019). 50. Ryu, H.-H., Park, K.-J., Yoon, C. S. & Sun, Y.-K. Capacity Fading of Ni-Rich Li[Ni x Co y Mn 1– x – y ]O 2 (0.6 ≤ x ≤ 0.95) Cathodes for High-Energy-Density Lithium-Ion Batteries: Bulk or Surface Degradation? Chem. Mater. 30, 1155–1163 (2018). 51. Natarajan, S., Moodakare, S. B., Haridoss, P. & Gopalan, R. Concentration Gradient-Driven Aluminum Diffusion in a Single-Step Coprecipitation of a Compositionally Graded Precursor for LiNi 0.8 Co 0.135 Al 0.065 O 2 with Mitigated Irreversibility of H2 ↔ H3 Phase Transition. ACS Appl. Mater. Interfaces 12, 34959–34970 (2020). 52. Xie, H. et al. Synthesis of LiNi 0.8 Co 0.15 Al 0.05 O 2 with 5-sulfosalicylic acid as a chelating agent and its electrochemical properties. J. Mater. Chem. A 3, 20236–20243 (2015). 53. Kim, J.-H. et al. Variation of Electronic Conductivity within Secondary Particles Revealing a Capacity-Fading Mechanism of Layered Ni-Rich Cathode. ACS Energy Lett. 3, 3002–3007 (2018). 54. Makimura, Y., Zheng, S., Ikuhara, Y. & Ukyo, Y. Microstructural Observation of LiNi 0.8 Co 0.15 Al 0.05 O 2 after Charge and Discharge by Scanning Transmission Electron Microscopy. J. Electrochem. Soc. 159, A1070–A1073 (2012). 176 177 Chapter 5 - Enhancing the Ionic Conductivity of Poly(3,4- propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries 1 Abstract Mixed electron and ion conducting polymers serve as excellent candidates for polymer binders in lithium-ion batteries (LIBs) due to an extension of functionality beyond simple mechanical adhesion. Such dual conduction was observed in our recent report on dihexyl-substituted poly(3,4- propylenedioxythiophene) (PProDOT-Hx 2) that showed excellent performance as a cathode binder for LiNi 0.8Co 0.15Al 0.05O 2 (NCA). However, ionic conductivity was significantly lower than its electronic counterpart. To enhance mixed conduction, here we report a family of synthetically tunable, electrochemically stable, random copolymers based on PProDOT-Hx2, in which the hexyl (Hex) side chains are replaced to varying extents with oligoether (OE) side chains, generating a series of (Hex:OE) PProDOTs. When the OE content was varied from 5% to 35%, the resulting copolymers were insoluble in the battery electrolyte and were stable after 100 electrochemical doping/de- doping cycles. Electron paramagnetic resonance and electrochemical kinetics studies were performed to illustrate the reversible and fast electrochemical doping process of (Hex:OE) PProDOTs. Electronic and ionic conductivity measurements as a function of electrode potential, showed a decrease in electronic conductivity and concurrent increase in ionic conductivity with increasing incorporation of oligoether side chains. X-ray scattering studies on electrochemically- doped polymers indicate a decline in crystalline ordering with increase in OE content of the (Hex:OE) PProDOTs, suggesting that decreasing crystallinity is responsible for both the increased ionic and 178 reduced electronic conductivity. Compounding these structural changes, swelling studies show a linear mass increase with OE content upon electrolyte exposure, indicating that solvent-induced swelling and electrolyte uptake plays a significant role in the ability of these polymers to conduct ions. Finally, the performance of the cells was studied via electrochemical impedance spectroscopy (EIS), galvanostatic charge-discharge, rate capability testing, differential capacity vs voltage analysis, using NCA cathodes to understand the role of these polymers as mixed electron and Li + ion conducting polymer binders in LIBs in comparison to the commonly used polyvinylidene fluoride, PVDF. It was observed that (75:25) PProDOT containing 25% of oligoether side chains, achieved the highest rate capability and fastest charging and discharging characteristics under symmetric testing conditions. The synthetic flexibility to fine-tune electronic and ionic conductivity make (Hex:OE) PProDOTs, a promising new class of mixed conducting polymers for electrochemical energy storage application. Introduction Mixed ionic/electronic conductors (MIECs) are materials, typically obtained from ceramics and polymers, allowing simultaneous transport of electrons and ions. 2–6 Such dual conduction is critical for the electrochemical processes in devices such as sensors, 7 actuators, 8 batteries, 9–11 organic electrochemical transistors (OECTs), 12 fuel cells, 13 light emitting electrochemical cells (LEECs) 14 and other bioelectronics and optoelectronics. 15,16 In this regard, organic mixed ionic-electronic conductors (OMIECs) have gained tremendous attention due to their soft nature, flexibility, synthetic tunability, processability and potential for high throughput. 2 Organic semiconductors, specifically conducting polymers, widely popular in the field of optoelectronics, 17 have found significant application in batteries and supercapacitors due to their ability to solvate and transport 179 ions along with electronic conduction. 18–20 While electronic conduction of conjugated polymers is enabled by their rigid, semi-crystalline morphology, ion transport on the other hand, demands an amorphous and less dense structure for accommodation of the larger ion carriers. 21 Such contradictory requirements for electronic and ionic conduction pose challenges to design materials with mixed conduction. Recent years have seen rapid advancement and development in the organic mixed conduction arena, addressing new design rules, structure-property relationships and morphological effects on dual conduction. 21,22 Lithium-ion batteries (LIBs) are one of the leading energy storage technologies23,24 dominating the field of renewable energy storage grids, consumer electronics and electric vehicles, where mixed electron and Li+ ion conduction plays a major role in determining overall energy and power density. The continuous increase in demand for LIB, calls for improvement in their energy and power density. Typically, LIB cathodes comprise of the active material (transition metal oxides), carbon additives for providing electron transport pathways and a polymer binder (usually polyvinylidene fluoride, PVDF) for ensuring mechanical integrity. While high- performance active materials can significantly boost the cell efficiency, minimizing inactive components like the electrically conductive carbon and polymer binder,25 can enhance the energy density of the battery. This focus has paved the way for investigating conducting polymer binders26,27 as multifunctional binders in LIBs, with electronic conduction as an added advantage alongside mechanical binding and reducing the overall electrode impedance. Over the last few years, numerous conducting polymers have been utilized as cathode and anode binders and/or additives in LIBs. 28–32 However, while these conjugated polymers leveraged electronic conductivity and displayed much better performance in comparison to PVDF, they lacked 180 the synthetic handle to fine tune ionic conductivity for Li + ion diffusion. Enhancing mixed conduction in conducting polymer binders is important for alleviating electrode impedance and enhancing cycle life. Employing OMIECs as polymer binders is potentially an effective strategy to optimize such mixed conduction. 29,33–35 While OMIECs can have varied architectures such as block copolymers 36 or homopolymers, blends of electron-conducting and ion-conducting polymers represent the most well-known approach. 37,38 A number of polymer blends have been reported for electrochemical energy storage applications exhibiting dual conduction. 39,40 The most common approach to facilitate ionic conduction is blending Li + ion conducting polyethylene glycol 41 (PEG) with electronically conducting conjugated polymers. 38,42–44 For example, Kwon et al. reported a PEG-coated composite magnetite anode using poly[3-(potassium-4- butanoate) thiophene] (PPBT) as the conducting polymer binder, resulting in enhanced cycle life and rate capabilities as a result of concurrent electronic and ionic charge transport. 42 Utilization of polymer blends however, or a multicomponent system, results in lower the electronic conductivity due to the increase in the amount of the non-conjugated component. In this regard, we recently reported dihexyl-substituted poly(3,4-propylenedioxythiophene) (PProDOT-Hx2) as a dual electronic and ionic conductive cathode binder for LIBs.45 PProDOT-Hx2 based cells displayed a five-fold enhancement in capacity in comparison to PVDF at high rates of discharge. The improved rate capability and cycling stability were attributed to dual electronic and ionic conductivity. The polymer, PProDOT-Hx2 surpassed P3HT as a conductive polymer binder by exhibiting an order of magnitude enhancement in ionic conductivity on electrochemical doping.45 Furthermore, molecular dynamics simulations explained such improved Li+ ion conduction due to solvent 181 swelling, which in turn is enabled by the open morphology and increased solvation due to the oxygen containing propylenedioxythiophene backbone. While PProDOT-Hx 2 was a step closer to achieving a single component multifunctional polymeric binder system, there still is a lot of room for performance improvement, specifically in the improvement of Li + ion conduction. Ionic conductivity of conjugated polymers is generally several orders of magnitude lower than the electronic conductivity. 46,47 At the same time, ionic diffusion distances required are generally much shorter than the distances required for effective electrode based electronic conductivity, so the ideal compromise is not well defined. Elevating ionic conductivity in order to enhance mixed conduction could significantly improve the rate capability and capacity retention of LIBs, but a chemically tunable system is required to optimize both. One of the ways of achieving this is by appending ion-transporting, polar oligoether side chains to the conjugated backbone, 28,48–50 where, similar to crown ethers, the oxygen atoms along the side chains coordinate with Li + ion. For example, Reynolds et al. have reported an oligoether functionalized propylenedioxythiophene (ProDOT)-based copolymer with redox stability for thousands of redox cycles along with excellent charge-storage capacity. 48 It has also been well reported that incorporation of oligoether side chains endows the conducting polymer with a larger surface area, decreased π-stacking distance 51,52 and a larger volumetric capacitance. 48,53,54 Thus, we directed our efforts towards incorporating oligoether side chains into our previously reported PProDOT-Hx 2 system. We investigated a series of random copolymers based on dihexyl-substituted propylenedioxythiophenes, whereby the hexyl side chains of the homopolymer (PProDOT-Hx2) have been replaced to varying extents with oligoether side chains, named as (Hex:OE) PProDOTs. We 182 have presented a detailed synthesis and experimental study for a comprehensive understanding of the balance between electronic and ionic conduction, their associated structural change upon electrochemical doping and their ultimate usage as a mixed conductive binder for LiNi0.8Co 0.15Al0.05O2 (NCA) cathodes. We have shown extensive and rigorous cell testing to elucidate the role of mixed conduction in this random copolymer series when used as a polymer binder for NCA cathodes and the effect on battery performance relative to our previously reported PProDOT-Hx2 and commercial PVDF. 183 Synthesis of (Hex:OE) PProDOT Random Copolymers using Direct Arylation Polymerization (DArP) The (Hex:OE) PProDOT random copolymer series was synthesized using our previously reported 45 green method of direct arylation polymerization (DArP) (Scheme 2.3). 45,55,56 We have incorporated oligoether side chains into PProDOT-Hx 2 by employing a random copolymerization strategy, which is a facile way to integrate different functionalities into a single polymeric system. 57,58 The hexyl (Hex) versus oligoether (OE) side chain ratio (Hex:OE) was tuned from (65:35) to (95:5) to generate a family of four random copolymers, which were subsequently used for further characterization and ultimate investigation for cathode binder applications in LIBs with NCA cathodes. Number-average molecular weight (M n) and dispersity (Ð) of the resulting polymers were 13.7 kDa (1.49), 16.3 kDa (1.78), 13.9 kDa (1.81) and 17.6 kDa (1.74) for (Hex:OE) (65:35), (75:25), (85:15) and (95:5), respectively. NMR and GPC traces for monomers and polymers can be found in Supporting Information (Figure 2.9-2.21) Figure 2.24-2.28). Incorporation of 50 percent or more of OE relative to hexyl side chains (Figure 2.18-2.21), results in dissolution of the polymer in battery electrolyte and therefore is unsuitable for polymer binder applications. The selected four (Hex:OE) random copolymers (Scheme 5.1), along with PProDOT- Hx 2, were found insoluble in battery electrolyte and were synthesized on the gram scale and analyzed relative to PVDF. 184 Scheme 5.1. Synthesis of (Hex:OE) PProDOT Random Copolymers using DArP Electrochemical Properties of (Hex:OE) PProDOT Random Copolymers To determine the suitability of these polymers for use in LIB cathodes, we examined the electrochemical doping (oxidation), potential window of stability, kinetics, and long-term cycling of the polymer films in battery electrolyte. The cyclic voltammogram (CV) of (85:15) PProDOT (Figure 5.1a) and all other ProDOT variants (Figure 5.2-4a), revealed two main oxidation peaks with the first peak at 3.21 V vs Li/Li + and a second oxidation peak at 3.74V vs Li/Li + . Two main reduction peaks were also observed at 3.12V and 3.61 V vs Li/Li + . Similarly, two redox peaks were observed for all the other polymers at similar potentials (Figure 5.2). Furthermore, when the potential window limit was expanded gradually from 4.0 V to 4.5V vs Li/Li + (Figures 5.1a and 5.2), no change in redox peak shape was observed, indicating that there is no degradation upon polarization to 4.5 V vs Li/Li + . The wide operating window of 2.8 to 4.5 V makes the (Hex:OE) PProDOT copolymer family ideally suited to work with various commercially relevant LIB cathodes such as NCA. 185 Figure 5.1. Electrochemical performance of (85:15) PProDOT. (a) Cycling Voltammograms at 10 mV s-1 in various potential windows up to 4.5 V, (b) CV scans at different scan rates, and (c) long- term cycling at 10 mV s-1. Figure 5.2. CV data at 10 mV s -1 for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT-Hx 2 at various potential intervals. 186 The CV scans collected at scan rates from 20 to 100 mV s -1 between 2.8 V and 4.1 V vs Li/Li + (Figures 5.1b and 5.3) showed that the electrochemical doping process was highly reversible up to 100 mV s -1 . In addition, the minimal shift in the redox peaks with increasing scan rates indicated that the electrochemical doping process was rapid. Furthermore, b-value analysis 59 yielded a b-value close to 1 that indicated a non-diffusion controlled process. A b-value of 0.95 and 0.93 was obtained for the first and second oxidation peaks while a b-value of 0.96 and 0.98 was obtained for first and second reduction peaks, respectively. Similarly, the b-values of redox peaks for all the polymers tested were found to be close to 1. This b-value analysis indicated that the kinetics of doping are rapid, and the polymers can facilitate rapid electron and ion transport. Figure 5.3 CV scan at varying scan rates between 100 and 20 mV s -1 for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT-Hx 2. 187 The long-term electrochemical stability of these polymers was studied by repeated cycling over the potential range 2.8 V to 4.1 V at 10 mV s -1 for 100 cycles. As can be seen from Figure 5.1c, a significant fraction of the capacity (71%) was retained after 100 cycles. Furthermore, no major peak shifts were observed indicating stable electrochemistry for these polymers. This electrochemical stability was observed for all (Hex:OE) PProDOTs (Figure 5.4). However, it has been observed that even 35% of oligoether content ((65:35) PProDOT) led to the dissolution of the polymer film and loss in capacity during extended cycling (Figure 5.5). Figure 5.4 Long-term CV curves for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT-Hx 2 thin films at 10 mV s -1 . 188 Figure 5.5 (a) Capacity retention of (Hex:OE) PProDOT copolymer family after 100 cycles at 10 mV s -1 as a function of oligo ether content. Electron Paramagnetic Resonance (EPR). Electrochemical doping of conjugated polymers generates polarons and bipolarons, the movement of which along or between the polymer chains invariably results in enhanced electronic conductivity. 60–62 Hence, the charge transport properties in these systems are a function of degree of doping. The electronic conductivity further depends on the charge carrier density and mobility, which is again dependent on the degree of doping. EPR measurements were used to determine the nature of charge carriers in the (Hex:OE) PProDOTs at the molecular level. A singly charged polaron is associated with an unpaired electron and has a spin of ½. As the charge density increases, additional charge carrier structures can be formed, including bipolarons, which are spinless and have a charge of 2. 63–66 EPR can be used to count spins in the oxidized polymer, which in turn provides insight into the structure of the charge carriers as a function of the oxidation state. Comparing the spin concentration in the films measured via EPR (Figure 5.6b) with the charge concentration measured via CV (Figure 5.6a) gave a direct 189 indication of the relative fraction of polarons in the material (i.e. number of polarons / dopants) 67 as a function of electrode potential. These studies were conducted by Gordon Pace of UCSB in Rachel Segalman’s laboratory. It is worth noting that the spin character of these films is quite stable over time, where Figure 2.22 shows EPR results for (75:25) PProDOT measured on the day of oxidation and after 1 week. The most apparent result is that all PProDOT variants show a similar trend in polaron character vs potential. At 3.2 V (vs. Li/Li + ), the first data point in Figure 5.6b, the ratio of spins to charges is between 60% and 90%, indicating that polarons are the dominant charge carriers in all variants. This finding is consistent with the charge carriers in all variants and also with the spectroelectrochemistry reported for ProDOTs. 68–70 The ratio of spins to charges subsequently decayed as the potential increased, indicating a shift towards spinless (bipolaronic) charge carriers at higher doping levels, which is consistent with the PProDOT literature. 68 Similar behavior has been found in other polymers, 66,67 and is attributed to increased stability of spinless carriers at high charge concentrations due to charge screening. By 3.4 V (vs. Li/Li + ), the relative fraction of polarons is below 10% for all polymers and remains below 3% after 3.8V. 190 Figure 5.6 Representative cyclic voltammograms for each of the PProDOT variants, acquired at 1 mV/s in 1 M LiPF 6 in EC:DMC (1:1 v/v). (b) The number of radicals (or equivalently polarons) detected via EPR, normalized to the number of dopants, as measured by CV, as a function of electrochemical potential. The PProDOT-Hx 2 data (black curves) is adopted from our previous work. 45 Model studies on ProDOT-Hx 2 oligomers 71 and our previous report on the homopolymer, PProDOT-Hx 2, 45 indicated that the polaron-bipolaron transition occurs with two closely aligned oxidation processes around 3.2-3.4 V vs Li/Li + , often manifesting themselves in a single apparent peak in voltammograms due to overlapping of the individual oxidation peaks. This is consistent with the data presented in Figure 5.6. Additionally, a smaller, second oxidation peak is often observed in PProDOTs and is apparent in Figure 5.6a around 3.8-4.0 V vs Li/Li + . Given the lack of EPR signal in this regime, this is likely a two-electron transfer process which forms additional 191 bipolarons. Thus, regardless of the side chain chemistry, all PProDOT variants exhibit primarily polaronic charge carriers at low doping levels (3.2 V and below), and then transition to have 90% or more bipolaron character by 3.3V. EPR data aligns well with the electronic conductivity data (Figure 5.7b) shown in the following section. We observe orders of magnitude enhancement in electronic conductivity on low levels of electrochemical doping. This is indicated by a peak where there is a maximum concentration of the highly mobile polarons. On further doping, the electronic conductivity drops due to the increasing conversion of polarons to spinless bipolarons. Such decline is attributed to the lower mobility of bipolarons 72 and disruption of the polymer chain connectivity due to solvent-induced swelling of the (Hex:OE) PProDOTs due to polar oligoether side chains. Finally, at high levels of doping, the conductivity plateaus where all the polarons have been totally converted to bipolarons, owing to the increased stability of spinless charge carriers at high charge concentration. 61 This implies that any electronic conductivity variation along the (Hex:OE) PProDOT series does not stem from polaron/bipolaron character, but rather from other phenomena correlated to incorporation of polar oligoether side chains, such as an increase in free volume or solvent swelling. Electronic and Ionic Conductivity of (Hex:OE) PProDOT Random Copolymers Electronic Conductivity. To determine the effect of oligoether side chains of the (Hex:OE) PProDOTs on the electron transport within the polymer structure, electronic conductivity of the polymer thin films was measured using electrochemical impedance spectroscopy (EIS) in 1 M LiTFSI in EC/DMC as a function of electrochemical doping (electrode potential) using a previously- 192 described method, 73 where we have made a pioneering contribution to report a method that simultaneously determines electronic and ionic conductivity of conducting polymers as a function of electrochemical doping. At low electrochemical doping levels (below 3.2 V vs Li/Li + ) all (Hex:OE) PProDOTs had a similar electronic conductivity starting at ~5 × 10 -6 S cm -1 at 2.8 V vs Li/Li + (Figure 5.7b). The electronic conductivity of the polymer films increased steadily to reach 10 -2 S cm -1 at 3.2 V vs Li/Li + which aligned with the first oxidation peak. The conducting polymer films reached their peak electronic conductivity beyond 3.2 V, where the values started to diverge, with (65:35) PProDOT having the lowest maximum electronic conductivity of 1.4 × 10 -1 S cm -1 at 3.4 V vs Li/Li + followed by (75:25) PProDOT which has an electronic conductivity of 2.4 × 10 -1 S cm -1 at 3.4 V vs Li/Li + . The trend continued with (85:15) PProDOT and (95:05) PProDOT with a maximum electronic conductivity of 6.1 × 10 -1 S cm -1 and 1.1 S cm -1 , respectively. Interestingly, the electronic conductivity of all (Hex:OE) PProDOTs dropped at potentials higher than 3.4 V vs Li/Li + to reach a plateau of ~2 × 10 -2 S cm -1 at high doping levels. Although a similar overall trend was observed with the PProDOT-Hx 2 homopolymer, a higher electronic conductivity was observed at low doping levels (below 3.2V vs Li/Li + ) and high doping levels (above 3.4 V vs Li/Li + ). However, PProDOT-Hx 2 reaches a maximum conductivity of 1 S cm -1 , similar to that of (95:5) PProDOT between 3.3 and 3.5 V vs Li/Li + . Electronic conductivity values for PProDOT:Hx 2 are in agreement with that reported in our previous work. 45 Thus, we can conclude that introducing oligoether side chains into PProDOT-Hx 2 lowered its electronic conductivity, and there is an inverse relationship between the oligoether content and the maximum electronic conductivity that is attained between 3.2 and 3.5 V vs Li/Li + . 193 Ionic Conductivity. EIS was also used to measure the ionic conductivity of polymer thin films in 1 M LiTFSI in EC/DMC as a function of electrochemical doping (Figure 5.7c). Upon doping, the ionic conductivity of PProDOT-Hx 2 reached 1 × 10 -7 S cm -1 . However, upon introduction of OE side chains, the ionic conductivity of (Hex:OE) PProDOTs increased by as much as a factor of four with (65:35) PProDOT reaching 4 × 10 -7 S cm -1 . The incorporation of only 5% oligoether side chains doubled the ionic conductivity compared to PProDOT-Hx 2. This trend in ionic conductivity is well supported by swelling studies (Figure 5.12, see below), as we observe that ion conduction is proportional to the amount of Li + coordinating OE, which in turn is proportional to the amount of electrolyte uptake, indicated by the linear increase in % mass on swelling in propylene carbonate (PC) electrolyte. We note that while the ionic conductivity is improved with OE incorporation, it is still orders of magnitude lower than the electronic conductivity. 194 Figure 5.7 (a) CV scans of (Hex:OE) PProDOTs at 5 mV s -1 in 1M LiTFSI dissolved in EC/DMC. (b) Electronic and (c) Ionic conductivities of the copolymer series. 195 Morphology Changes of (Hex:OE) PProDOTs with Electrochemical Doping To understand the trends in conductivity and how they change as a function of potential for the various (Hex:OE) PProDOTs, we investigated the polymer structure in both neutral (Figure 5.8) and electrochemically doped forms using grazing incidence wide-angle X-ray scattering (GIWAXS). This experiment provides information about total scattering intensity (which correlates with crystallinity), lattice constants (which change upon doping), and about orientation of the polymer chains with respect to the substrate. The GIWAX experiments were conducted by Charlene Salamat of Prof. Sarah Tolbert’s laboratory. The results are cited here to support the learnings from electrochemical studies. All of these polymers are fairly disordered, with only two semi-crystalline peaks observed in most cases, and almost no texture. At ~0.3-0.4 Å, the lamellar (100) peak, which corresponds to the side chain spacing between the polymer chains, is observed. At ~1.4 Å, the π-stacking (010) peak appears, which corresponds to the distance between polymer chains along the lattice vector closest to the π-stacking direction. 196 Figure 5.8 The integrated GIWAXS diffractograms of all the polymers in their neutral state. The neutral polymers are all relatively similar: there are two dominant peaks: the lamellar peak and the pi-stacking peak, indicating that they are not very crystalline. Figure 5.9a shows the fully integrated diffractograms of the neutral polymers normalized to the lamellar (100) peak. In the neutral form of the polymers, the two peaks are at roughly the same peak position, with a slight shift of the lamellar peak to larger q (smaller d-spacing) as the oligoether content increases, attributed to the curling of ether side-chains. 74,75 All polymers appear similarly disordered, and show similar diffraction intensity when un-normalized, and all polymers are much more disordered than similar polymers such as poly(3-hexylthiophene) (P3HT), confirming the disordered nature of PProDOTs. Moreover, as seen in the raw date (Figure 5.10), the diffraction rings are fairly isotropic, in contrast to more crystalline semiconducting polymer such as P3HT, which shows strong in- and out-of-plane texturing. 76–78 Because of the lack of texture in these polymers, the 2D GIWAXS data can be integrated to 1D patterns without loss of information. 197 Figure 5.9 a) Full integration of GIWAXS diffractograms for (Hex:OE) PProDOTs, normalized to the lamellar peak under neutral conditions b) at the first oxidation peak, c) at the second oxidation peak, normalized to the π-stacking peak, d-f) select polymers with varying amounts of oligoether under neutral conditions, and doped at the first and second oxidation peak 198 Figure 5.10 2D diffractograms of the various polymers with differing amounts of oligoether content, at neutral (1 st column), at first oxidation peak (2 nd column), and at 2 nd oxidation peak (3 rd column). Electrochemical doping of (Hex:OE) PProDOTs was carried out in a mixture of ethylene carbonate/dimethyl carbonate (EC/DMC) with LiTFSI, where the TFSI - anion participates as the counter-ion to the doped polymer. Figure 5.9b shows un-normalized diffraction data for the full polymer series with OE content ranging from 5% to 35% upon doping at the first oxidation peak at ~3.2 V vs. Li/Li + . Figure 5.9c shows the corresponding diffraction data, normalized to the π- stacking peak, upon being doped at the second oxidation peak at ~3.8 V vs. Li/Li + . Figures 5.9(d- 199 f) present un-normalized GIWAXS profiles at different doping states for select (Hex:OE) PProDOTs with high (35%), low (5%) and intermediate (15%) amounts of OE under neutral conditions and when doped at the first and second oxidation peaks (Figure 5.11 for (75:25) PProDOT). This data allows for comparison of relative crystallinity, using diffraction peak heights, in a single sample as a function of doping state. We observe that doping at the first oxidation peak dramatically increases the crystalline order of many of these polymers, as evidenced by the significant increase in the (100) diffraction peak intensity and significant peak narrowing. This suggests that electrochemical doping induces order in these naturally disordered (Hex:OE) PProDOTs. Similar results have been observed previously for regio-random P3HT, which shows a dramatic increase in crystallinity upon doping. 78,79 Interestingly, Figure 5.9 indicates that the enhancement is not uniform across the series. In moving from (Hex:OE) PProDOT (95:5) to (65:35), the magnitude of the increase in peak intensity goes steadily down. For the (65:35) PProDOT containing the most OE content (Figure 5.9f), the lamellar peak still narrows upon doping, but no increase in diffraction intensity is observed. Figure 5.11 Full integration of GIWAXS diffractograms for (75:25) PProDOT under neutral conditions, doped at the first and second oxidation peaks. 200 In addition to the intensity increase, which is indicative of an overall increase in crystallinity when doped at the first oxidation potential, there is also a shift of the lamellar peak to lower q or higher d-spacing. This is accompanied by a shift towards higher q (and smaller d-spacing) for the π- stacking peak. Both changes are commonly observed in thiophene based polymers and are the result of reorientation of the π-stacking direction with respect to the unit cell axis to allow room for the dopant counterions in the lamellar side-chain region. 79–81 On doping further to the second oxidation peak, the intensity of both the lamellar and π-stacking peaks decease, indicating disruption of the ordered structure (Figure 5.9c). This correlates well with the electronic and ionic conductivity trend of the (Hex:OE) PProDOT family (Figure 5.7b and 5.7c). At the first oxidation peak at ~3.2 V vs Li/Li + , the increase in crystallinity at lower OE contents (Figure 5.9b) leads to high electronic conductivity (Figure 5.7b). In contrast, the higher OE content materials show both smaller changes in crystallinity upon electrochemical doping and lower electronic conductivity. The most crystalline order can be induced in the (95:5) PProDOT system upon electrochemical doping, which aids in electronic conduction, leading to the highest electronic conductivity among the copolymer family. On the other hand, OE side-chains are well-known to be Li + ion conductors in disordered form. 82,83 Because doping is not able to induce significant crystallinity in the high OE fraction samples, we find that the ionic conductivity shows the opposite trend compared to the electronic conductivity. Higher OE fraction samples are the most disordered and thus show higher ionic conductivity. For example, Figure 5.7c shows that (65:35) PProDOT has the highest ionic conductivity in the series and PProDOT-Hx 2 has the lowest. 201 Solvent Swelling of (Hex:OE) PProDOTs Using Battery Electrolytes While there is a strong anti-correlation between the diffraction intensity and electronic conductivity and a strong correlation between the diffraction and the ionic conductivity, the root cause of that correlation remains unclear. It is possible that the OE side chains intrinsically disorder the material, and that disorder favors ionic conductivity and disfavors electronic conductivity. Alternatively, the polar nature of the OE side chains, combined with the increased disorder, could facilitate solvent swelling, and it is the presence of the extra solvent that dominantly increases ionic conductivity and decreases electronic conductivity. Although clean separation of these trends is likely impossible, insight can be gained by exposing each polymer to electrolyte vapor to measure the equilibrium solvent swelling as a function of OE content in these (Hex:OE) PProDOTs (Figure 2.23). From Figure 5.12, we observe that there is an almost a linear increase in swelling with increasing OE content, culmination in a near 2x mass increase in the polymers with the highest OE content. Electrolyte uptake of this magnitude can easily generate ion-transport pathways, improving ionic conductivity. Unfortunately, swelling simultaneously disrupts π-conjugation and thereby, electron-transport pathways, which explains the inverse trend in both electronic and ionic conductivity with increase in OE content within the random copolymer family. Importantly, this result suggests that tuning electrolyte swelling may be the primary method for optimizing the interplay between electronic and ionic conductivity in Li + battery systems, where polar organic solvents will always have some propensity to swell polymer binders. 202 Figure 5.12 Swelling study of the (Hex:OE) PProDOTs using propylene carbonate (PC) electrolyte, showing an almost linear trend in % mass increase with increase in oligoether incorporation. Electrochemical cycling of NCA cathodes incorporating (Hex:OE) PProDOTs as binders. To test the function and efficacy of these polymers in a practical battery electrode, NCA cathodes employing (Hex:OE) PProDOTs, PProDOT-Hx 2, and PVDF polymers as binders were repeatedly cycled at a 1C rate for 200 cycles. All the electrodes with conducting polymer (CP) binders, barring (65:35) PProDOT, retained 17 to 40% higher capacity over the cells with PVDF binder (Figures 5.13 and 5.14). While all the (Hex:OE) PProDOTs are found to be beneficial in retaining a higher capacity compared to PVDF, (75:25) PProDOT retained the highest capacity. We attribute the improved cycling stability to the ability of these polymers to maintain inter-particle electronic and ionic connectivity. 203 Figure 5.13 Specific capacity as a function of cycle number at constant charge/discharge rate of 1C with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOTs cells with a cathode mass composition of 90% NCA, 4% binder, 6% carbon and an areal loading of 3.1 ± 0.4 mg cm -2 . Figure 5.14 Corresponding galvanostatic charge-discharge curves from Figure 5.13; a) Li-NCA- (85:15) PProDOT and Li-NCA-(95:05) PProDOT, b) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, c) Li-NCA-PProDOT-Hx 2 and Li-NCA-PVDF. 204 It appears that solvent-induced and doping-induced swelling of oligoether containing conducting polymers is beneficial in maintaining good electrical interconnectivity and binding of the electrode particles during the repeated expansion and contraction during charge/discharge. However, the results in Figure 5.13 indicate that the benefits require an optimal level of oligoether group incorporation. Excessive swelling is clearly undesirable and leads to loss of binding action as seen in the steep decay of cycle life with (65:35) PProDOT. The rapid decay in capacity of electrodes with (65:35) PProDOT is consistent with the results of cyclic voltammetry studies on the polymer film that suggested partial dissolution of the polymer during extended cycling. The above understanding is supported by the results of impedance analysis, differential capacity analysis, cyclic voltammetry and wide-angle X-ray scattering. Impact of the CP binders on Rate Capability Electrochemical impedance of the various cells with (Hex:OE) PProDOTs after 200 cycles indicated differences at low frequencies (0.01 Hz to 1 Hz ) where the impedance is governed by the mass transport of lithium ions by diffusion in and out of the cathode particles (Figure 5.15, 3D-Nyquist plots as a function of State-of-Charge). The Warburg coefficients measured at any particular SOC (Figure 5.16 and 5.17), indicated three to five times higher interfacial contact area for lithium diffusion with the CP polymer cells relative to the PVDF cells. Warburg coefficient analysis At low frequencies in the range of 0.01 Hz to 1 Hz, the impedance is governed by the diffusional processes involved in the mass transport of lithium ions in and out of the cathode particles. Such diffusion can be modeled by the finite length diffusion element (for a reflective boundary) 84 , whose impedance as a function of frequency is given by 205 𝑍 ̂ 𝜔 = 𝜎 √ 𝜔 𝑐𝑜𝑡 ℎ ( √ 𝑗𝜔 𝐷 𝑙 ) ( 1 − 𝑗 ) ( 1) where ω is the angular frequency, l is a characteristic length of the active particle, D is the diffusion coefficient of the lithium ions, j = (‐1) 0.5 . σ is the Warburg coefficient denoted as 𝜎 = 𝑅𝑇 𝑛 2 𝐹 2 𝐴 √ 2 ( 1 𝐷 0.5 ∙ 𝐶 ∙ 𝑆𝑂𝐶 ∙ ( 1 − 𝑆𝑂𝐶 ) ) ( 2) where R is the ideal gas constant, T is temperature, n is the number electrons, F is the faraday constant, A is the active area of contact, C is the concentration of the lithium ions and SOC is the state‐of‐charge. The Nyquist plot for this type of circuit element is characterized by a sloping 45‐ degree line at high frequency transitioning to a constant value for the real component of impedance and a vertical line at low frequencies (Figure 5.16). 85 At low frequencies, the frequency dependence of impedance is given by 𝑍 ̂ 𝜔 = √ 2𝜎𝑙 3√ 𝐷 − 𝑗 √ 2𝐷𝜎 𝜔𝑙 ( 3) Consequently, Z real*(‐Z imag)= 2  2 /3 . Using the impedance values at low frequency we can determine the value of  for the various states of charge for all the polymers. Once  has been determined, the value of D/L 2 can be determined either from real or imaginary components. Of the various CPs studied, the (75:25) PProDOT cell exhibited the highest interfacial contact area. Thus, we concluded that the cells with (75:25) PProDOT exhibit an optimal level of swelling (Figure 5.12) for the necessary resilience and stable interfacial contact area for ion diffusion even after 200 cycles. 206 Figure 5.15 Three-dimensional chart of the impedance response after 200 cycles as function of SOC. Nyquist plots as function of SOC for a) NCA-PProDOT-Hx 2, b) NCA-(95:05) PProDOT, c) NCA-(85:15) PProDOT, d) NCA-(75:25) PProDOT, e) NCA-(65:35) PProDOT, and f) NCA-PVDF cells. Figure 5.16 Nyquist plots of total impedances for (a) semi-infinite linear, (b) transmissive, and (c) reflective boundary. 85 207 Figure 5.17 Calculated Warburg coefficient σ as function of state of charge (SOC) for the NCA- (Hex:OE) PProDOTs, NCA-PProDOT-Hx 2 and NCA-PVDF cells after 200 cycles. The benefit of the CP polymers in reducing the polarization losses was most prominent in the part of the charge/discharge curve where the diffusion coefficient for lithium ions is the lowest, namely, in the potential range of 3.55 to 3.65 V or close to 0% SOC. 86–91 Analysis of the peaks in the differential capacity plots of dQ/dV vs V (Figure 5.18) reveal that with multiple cycling, the cells with NCA-(75:25) PProDOT and NCA-(85:15) PProDOT exhibited a lower shrinkage of the peak in the range of 3.55 to 3.65 V suggesting that the improved ionic conductivity and higher interfacial contact area for diffusion seen with these polymers are specifically beneficial in reducing the polarization losses. 92,93 208 Figure 5.18 Differential capacity vs voltage analysis (dQ dV ‐1 vs V) for the Li‐NCA‐(Hex:OE) PProDOTs cells from Figure 5.13-14 as a function of cycle number charged‐discharged at 1C with calculated between 2.7 and 4.2 V vs Li + /Li for a) Li‐NCA‐PProDOT‐Hx 2, b) Li‐NCA‐(65:35) PProDOT, c) Li‐NCA‐(75:25) PProDOT, d) Li‐NCA‐(85:15) PProDOT, e) Li‐NCA‐(95:05) PProDOT, and f) Li‐NCA‐ PVDF cells. After two formation cycles at C/10, the impedance response was measured as a function of SOC at various points on the GCD curves (Figure 5.19a-c). PProDOT-Hx 2 and (Hex:OE) PProDOT cells exhibited a lower resistance than the NCA-PVDF cell at all states of charge (Figure 5.19d-f). The impedance response was normalized by the electrode geometrical area (Nyquist plots) and the electrode active mass loading (ohm mg -1 ) due to the possible differences in loading and area that are present from one electrode to other. Among the conducting polymers the impedance at low frequencies followed the order NCA-(65:35) > NCA-(75:25) > NCA-(85:15) > NCA-PProDOT-Hx 2 > NCA-(95:05). These trends in impedance suggest that the higher electronic and ionic conductivity 209 (Figure 5.7b and 5.7c) combined with optimal degree of swelling and large interfacial contact area (Figure 5.12-5.17) resulted in the least impedance for charge-transfer and mass-transfer processes in the cells. These differences in impedance are also reflected in the discharge capacity realized at different rates. Figure 5.19 a) Galvanostatic charge-discharge profiles for Li-NCA-PProDOT-Hx 2 and Li-NCA- (75:25) PProDOT cells at C/20 with an areal mas loading of ∼ 5 mg cm -2 and a charging cut-off of 4.2 V vs Li + /Li. The impedance response was measured at OCV as function of SOC (denoted with yellow circles). b) and c) Corresponding three-dimensional Nyquist plots of the experimental impedance response (denoted with color spheres) as function of SOC during charge for the same Li-NCA-(75:25) PProDOT (b) and Li-NCA-PProDOT-Hx 2 (c)cells used to generate the data in part (a). Real part of the impedance response against frequency for the Li-NCA-(65:35) PProDOT, Li- NCA-(75:25) PProDOT, Li-NCA-(85:15) PProDOT, Li-NCA-(95:05) PProDOT, Li-NCA-PProDOT-Hx 2, and Li-NCA-PVDF cells at d) 3.7 V e) 3.9 V, and f) 4.08 V vs Li + /Li. To determine the discharge rate capability, the cells were first charged up at a low rate (C/5) and then discharged at various rates ranging from C/5 to 6C. At C/10 and C/5 rates the NCA- 210 conducting polymer cells had a higher discharge capacity than the NCA-PVDF cells. However, beyond 1C rate of discharge, the NCA-(Hex:OE) PProDOT cells exhibited a lower discharge capacity compared to NCA-PProDOT-Hx 2 and even the NCA-PVDF cells (Figures 5.20a and 5.21a- c). These results suggested that the discharge processes at high rates was specifically hindered in the cells with NCA-(Hex:OE) PProDOT polymers owing to their lower electronic conductivity, higher solvent swelling and polymer dissolution specifically for (65:35) PProDOT. This trend is consistent with the higher impedance at low frequencies exhibited by (Hex:OE) PProDOT relative to PProDOT-Hx 2. We attribute the higher capacity at the high rates of discharge for PProDOT-Hx 2 to its high electronic conductivity compared to the (Hex:OE) PProDOTs. 211 Figure 5.20 a) Specific capacity as a function of cycle number at various discharge rates with a constant charging rate at C/5 for Li-NCA-(Hex:OE) PProDOTs cells with an areal loading of 5.2 ± 0.8 mg cm -2 . b) Specific capacity as a function of cycle number with symmetric discharge/charge rate for Li-NCA-(Hex:OE) PProDOTs cells with an areal loading of 5.0 ± 0.8 mg cm -2 . In both cases, two formation cycles at C/10 were performed prior to rate testing. 212 Figure 5.21 Corresponding galvanostatic charge-discharge curves from Figure 5.20a; (a) Li-NCA- (85:15) PProDOT and Li-NCA-(95:05) PProDOT, (b) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, (c) Li-NCA-PProDOT-Hx 2 and Li-NCA-PVDF. Corresponding galvanostatic charge- discharge curves from Figure 5.20b; (d) Li-NCA-(85:15) PProDOT and Li-NCA-(95:05) PProDOT, (e) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, (f) Li-NCA-PProDOT-Hx 2 and Li-NCA-PVDF. In addition to testing for high discharge rates, ability to charge at high rates is of practical importance. To assess the ability of the cells to accept charge at high rates the cells were charged and discharged at the same rate (up to 6C) and referred to here as “symmetric charge/discharge rate capability”. We found that the NCA-(Hex:OE) PProDOT and NCA-PProDOT-Hx 2 cells yielded a higher discharge capacity compared to the NCA-PVDF cells (Figures 5.20b and 5.21d-f). Beyond 213 1C the NCA-(75:25) PProDOT retained the highest specific capacity among the NCA-(Hex:OE) PProDOT cells. We find that the ionic conductivity and degree of swelling of the (Hex:OE) PProDOT polymers has a significant impact on the fast charging process. We have found that the (75:25) PProDOT polymer exhibited the lowest polarization (Figures 8, 5.22 and 5.23) and highest discharge capacity consistent with the highest ionic conductivity and an optimal degree of swelling. Whereas, when the degree of swelling of the polymers became excessive, polymer dissolution occurred and the rate capability was compromised, as exemplified by NCA-(65:35) PProDOT (Figures 5.21e, 5.4a and 97 % mass increase on swelling, Figure 5.12). Thus, we find that the discharge rate capability can be enhanced by increasing electronic conductivity, while the fast charge capability is enhanced by the high ionic conductivity and an optimal level swelling of the (Hex:OE) PProDOT polymers. 214 Figure 5.22 Differential capacity vs voltage analysis (dQ dV -1 vs V) for the Li-NCA-(Hex:OE) PProDOT family cells from Figure 5.20a with increments in the discharge rate and a constant charging rate of C/5 with an areal mass loading of 5.2 ± 0.8 mg cm -2 calculated between 2.7 and 4.2 V vs Li + /Li for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(65:35) PProDOT, c) Li-NCA-(75:25) PProDOT, d) Li-NCA-(85:15) PProDOT, e) Li-NCA-(95:05) PProDOT, and f) Li-NCA-PVDF cells. 215 Figure 5.23 Differential capacity vs voltage analysis (dQ dV -1 vs V) for the Li-NCA-(Hex:OE) PProDOT family cells from Figure 5.20b with increments in the discharge/charge rate with an areal mass loading of 5.0 ± 0.8 mg cm -2 calculated between 2.7 and 4.2 V vs Li + /Li for a) Li-NCA- PProDOT-Hx 2, b) Li-NCA-(65:35) PProDOT, c) Li-NCA-(75:25) PProDOT, d) Li-NCA-(85:15) PProDOT, e) Li-NCA-(95:05) PProDOT, and f) Li-NCA-PVDF cells. Effect of low binder content and carbon content. Since swelling to an optimal level appeared to be beneficial, we decided to modulate the swelling effects, by lowering the binder content from the baseline of 4% to 2% and 1% while keeping the carbon content at 3%. As with the baseline binder content of 4%, the NCA-(75:25) PProDOT and NCA-PProDOT-Hx 2 cells with 2% and 1% binder exhibited a significantly lower impedance at all frequencies compared to the NCA-PVDF cells (Figure 5.24a-b and 5.25). This effect is attributed to the higher electronic and ionic conductivity of the PProDOTs compared to PVDF. 216 Figure 5.24 Real part of the impedance response Zre against frequency for the NCA-(75:25) PProDOT, NCA-PProDOT-Hx 2, and NCA-PVDF cells with cathode composition with a) 2% binder, 3% carbon, 95% NCA, b) 1% binder, 3% carbon, 96% NCA and c) 1.5% carbon, 4% binder, 94.5% NCA at 50% SOC. 217 For the cells with low binder content, at low rates of discharge (C/10 and C/5) all the cells with the conducting polymers had comparable values of discharge capacity but higher than the NCA- PVDF cells (Figure 5.26a-b and 5.27a-f). However, at the higher rates from 1C to 6C, the NCA- (75:25) PProDOT cells had the highest capacity. At 4C, NCA-(75:25) PProDOT cells with 2% and 1% of binder both delivered a remarkable specific capacity of 88 mAh g -1 . Furthermore, the benefits of (75:25) PProDOT were seen even at a reduced binder content of 1%. We find that the lower polymer fraction in the electrode reduced the volume changes resulting from the swelling, and led to an improved performance under the conditions of repeated fast charging and fast discharge. Figure 5.25 Three-dimensional chart of the impedance response as function of SOC. Nyquist plots as function of SOC for a) 2% binder NCA-PVDF, b) 2% binder NCA-PProDOT-Hx 2, c) 2% binder NCA- 218 (75:25) PProDOT, d) 1% binder NCA-PVDF, e) 1% binder NCA-PProDOT-Hx 2, f) 1% binder NCA- (75:25) PProDOT, g) 1.5% Carbon NCA-PVDF, h) 1.5% Carbon NCA-PProDOT-Hx 2, and i) 1.5% Carbon NCA-(75:25) PProDOT cell. Figure 5.26 Specific capacity as a function of cycle number with increments in the discharge/charge rate (symmetric charge/discharge rate capability) with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOT cells with a cathode mass composition of a) 95% NCA, 2% binder, 3% carbon with an areal loading of 5.5 ± 0.3 mg cm -2 and b) 96% NCA, 1% binder, 3% carbon with a loading of 5.3 ± 0.6 mg cm -2 . 219 Figure 5.27 Corresponding galvanostatic charge-discharge curves from Figure 5.26a; (a) Li-NCA- (75:25) PProDOT, (b) Li-NCA-PProDOt-Hx 2, and (c) Li-NCA-PVDF. Corresponding galvanostatic charge-discharge curves from Figure 5.26b; d) Li-NCA-(75:25) PProDOT, e) Li-NCA-PProDOT-Hx 2, and f) Li-NCA-PVDF. The reduction of carbon content from 3% to 1.5% while keeping the binder content at 4%, resulted in higher impedance values for the NCA-(75:25) PProDOT at mid- and high frequencies (Figure 5.24c). The discharge capacities realized with (75:25) PProDOT are reduced by over 50% (Figure 5.28). At 1.5% carbon, the values of impedance for the NCA-PVDF and NCA-PProDOT-Hx 2 cells were comparable, indicating that the carbon content of at least 3% is required for achieving the long-range electrical connectivity of the particles and adding porosity to the electrode. The conducting polymers on the other hand provided the increased interfacial contact area and local 220 electronic and ionic transport. These findings point to the role of carbon being distinct from that of the conducting polymer, in addressing the long-range and short-range charge transport processes, respectively. Figure 5.28 a) Specific capacity as a function of cycle number with increments in the discharge/charge rate with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOT cells with a cathode mass composition of 94.5% NCA, 4% binder, 1.5% carbon and an areal loading of 4.9 ± 0.5 mg cm -2 . Corresponding galvanostatic charge-discharge curves in a) for; b) Li-NCA- (75:25) PProDOT, c) Li-NCA-PVDF, and d) Li-NCA-PProDOT-Hx 2. Conclusions A series of (Hex:OE) PProDOT random copolymers were synthesized via Direct Arylation Polymerization on the gram scale. The hexyl side chains of PProDOT-Hx 2 was replaced with varying content of oligoether side chains, in order to facilitate Li + ion conduction and thereby 221 fine-tune the electronic and ionic conductivity by regulating the side chain composition. When the composition of oligo ether was varied from 5% to 35%, the resulting copolymers were insoluble in the battery electrolyte and were electrochemically stable up to 4.5 V vs Li/Li + with good stability after 100 cycles. The electrochemical doping process was highly reversible up to 100 mV s -1 . Rapid kinetics of the electrochemical doping process enables fast ionic and electronic transport, justifying their relevance as cathode binders. Electronic and ionic conductivity measurements of the (Hex:OE) PProDOTs clearly demonstrated their dual conduction and the effect of oligoether side chains on the overall electron and ion transport. At low electrochemical doping levels, with increase in the percentage incorporation of oligoether side chains, the maximum electronic conductivity decreased along the copolymer series, with (95:5) PProDOT having the maximum value of 1.1 S cm -1 and (65:35) PProDOT having the minimum value of 0.14 S cm -1 . This is supported by the GIWAXS studies, which indicate a reduction in crystalline order in the doped polymers on moving from (95:5) to (65:35) PProDOT. On the other hand, ionic conductivity increased by a factor of four, with (65:35) PProDOT reaching a maximum value of 4 × 10 -7 S cm -1 , which is well supported by both the increased disorder with increasing EO content and the linear trend in electrolyte uptake with EO content observed in the swelling studies of the (Hex:OE) PProDOTs. In device studies, NCA electrodes made with (Hex:OE) PProDOTs as binders exhibited lower impedance compared with the electrodes with PVDF binder. At low rates of discharge, the benefit of the conducting polymers was significant. Oligoether containing PProDOTs led to an increased interfacial contact area and lower capacity decay over extended cycling, in addition to a higher discharge capacity after multiple cycles. Thus, electrodes fabricated with conducting 222 polymers with an optimal level of swelling and high active contact area ((75:25) PProDOT in this case) were found to be useful for achieving higher rate capability over a wider range of rates and can facilitate the charging process, as seen from symmetric rate capability studies and undiminished performance at reduced binder content. However, PProDOTs with a high degree of swelling, despite the high ionic conductivity, negatively impacted the electrode morphology and access to the surface area of the cathode particles. Therefore, reducing the mass fraction of the conducting polymers with high ionic conductivity is beneficial in avoiding the undesirable long-term consequences of swelling especially under repeated charge and discharge. Furthermore, the studies on varying carbon content and binder content confirm that the conducting polymer and carbon additive in the electrodes play distinct roles in maintaining the electrode performance. As such, PProDOTs containing optimum contents of oligoether side chains are a new class of synthetically tunable, electrochemically stable mixed electron and ion conducting cathode binders for LIBs, offering the advantage to fine-tune electronic and ionic conductivity through compositional variation. References 1. Das, P. et al. Enhancing the Ionic Conductivity of Poly(3,4-propylenedioxythiophenes) with Oligoether Side Chains for Use as Conductive Cathode Binders in Lithium-Ion Batteries. Chem. Mater. XX, (2022). 2. Paulsen, B. D., Tybrandt, K., Stavrinidou, E. & Rivnay, J. Organic mixed ionic–electronic conductors. Nat. Mater. 19, 13–26 (2020). 3. Paulsen, B. D., Fabiano, S. & Rivnay, J. Mixed Ionic-Electronic Transport in Polymers. Annu. Rev. Mater. Res. 51, annurev-matsci-080619-101319 (2021). 223 4. Thomas, E. M., Nguyen, P. H., Jones, S. D., Chabinyc, M. L. & Segalman, R. A. Electronic, Ionic, and Mixed Conduction in Polymeric Systems. Annu. Rev. Mater. Res. 51, annurev-matsci-080619-110405 (2021). 5. Papac, M., Stevanović, V., Zakutayev, A. & O’Hayre, R. Triple ionic–electronic conducting oxides for next-generation electrochemical devices. Nat. Mater. 20, 301–313 (2021). 6. Chen, C.-C., Fu, L. & Maier, J. Synergistic, ultrafast mass storage and removal in artificial mixed conductors. Nature 536, 159–164 (2016). 7. Shan, K., Yi, Z.-Z., Yin, X.-T., Dastan, D. & Garmestani, H. Conductivity and mixed conductivity of a novel dense diffusion barrier and the sensing properties of limiting current oxygen sensors. Dalton Trans. 49, 6682–6692 (2020). 8. Muñoz-Castro, M., Walter, N., Prüßing, J. K., Pernice, W. & Bracht, H. Self-Holding Optical Actuator Based on a Mixed Ionic–Electronic Conductor Material. ACS Photonics 6, 1182–1190 (2019). 9. Moy, D. & Narayanan, S. R. Mixed Conduction Membranes Suppress the Polysulfide Shuttle in Lithium-Sulfur Batteries. J. Electrochem. Soc. 164, A560–A566 (2017). 10. Wang, M. J., Wolfenstine, J. B. & Sakamoto, J. Mixed Electronic and Ionic Conduction Properties of Lithium Lanthanum Titanate. Adv. Funct. Mater. 30, 1909140 (2020). 11. Elizalde-Segovia, R., Irshad, A., Zayat, B. & Narayanan, S. R. Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature. J. Electrochem. Soc. 167, 140529 (2020). 12. Maria, I. P. et al. The Effect of Alkyl Spacers on the Mixed Ionic-Electronic Conduction Properties of N-Type Polymers. Adv. Funct. Mater. 31, 2008718 (2021). 13. Wang, X. et al. Charge-Transfer Modeling and Polarization DRT Analysis of Proton Ceramics Fuel Cells Based on Mixed Conductive Electrolyte with the Modified Anode–Electrolyte Interface. ACS Appl. Mater. Interfaces 10, 35047–35059 (2018). 14. Youssef, K., Li, Y., O’Keeffe, S., Li, L. & Pei, Q. Fundamentals of Materials Selection for Light- Emitting Electrochemical Cells. Adv. Funct. Mater. 30, 1909102 (2020). 15. Isaksson, J. et al. Electronic control of Ca2+ signalling in neuronal cells using an organic electronic ion pump. Nat. Mater. 6, 673–679 (2007). 16. van de Burgt, Y. et al. A non-volatile organic electrochemical device as a low-voltage artificial synapse for neuromorphic computing. Nat. Mater. 16, 414–418 (2017). 17. Dou, L., Liu, Y., Hong, Z., Li, G. & Yang, Y. Low-Bandgap Near-IR Conjugated Polymers/Molecules for Organic Electronics. Chem. Rev. 115, 12633–12665 (2015). 18. Katz, H. E., Searson, P. C. & Poehler, T. O. Batteries and charge storage devices based on electronically conducting polymers. J. Mater. Res. 25, 1561–1574 (2010). 19. Qin, K., Huang, J., Holguin, K. & Luo, C. Recent advances in developing organic electrode materials for multivalent rechargeable batteries. Energy Environ. Sci. 13, 3950–3992 (2020). 20. Song, H.-K. & Palmore, G. T. R. Redox-Active Polypyrrole: Toward Polymer-Based Batteries. Adv. Mater. 18, 1764–1768 (2006). 224 21. Onorato, J. W. & Luscombe, C. K. Morphological effects on polymeric mixed ionic/electronic conductors. Mol. Syst. Des. Eng. 4, 310–324 (2019). 22. Rivnay, J. et al. Structural control of mixed ionic and electronic transport in conducting polymers. Nat. Commun. 7, 11287 (2016). 23. Wang, H. et al. Recent Advances in Conversion-Type Electrode Materials for Post Lithium-Ion Batteries. ACS Mater. Lett. 956–977 (2021) doi:10.1021/acsmaterialslett.1c00043. 24. Tang, W. et al. Recent Advances of Mesoscale-Structured Cathode Materials for High Energy Density Lithium-Ion Batteries. ACS Appl. Energy Mater. 4, 2962–2975 (2021). 25. Shi, Y., Zhou, X. & Yu, G. Material and Structural Design of Novel Binder Systems for High-Energy, High-Power Lithium-Ion Batteries. Acc. Chem. Res. 50, 2642–2652 (2017). 26. Nguyen, V. A. & Kuss, C. Review—Conducting Polymer-Based Binders for Lithium-Ion Batteries and Beyond. J. Electrochem. Soc. 167, 065501 (2020). 27. Higgins, T. M. et al. A Commercial Conducting Polymer as Both Binder and Conductive Additive for Silicon Nanoparticle-Based Lithium-Ion Battery Negative Electrodes. ACS Nano 10, 3702–3713 (2016). 28. Hwang, C. et al. The rational design of a redox-active mixed ion/electron conductor as a multi- functional binder for lithium-ion batteries. J. Mater. Chem. A 9, 4751–4757 (2021). 29. Liu, G. et al. Polymers with Tailored Electronic Structure for High Capacity Lithium Battery Electrodes. Adv. Mater. 23, 4679–4683 (2011). 30. Zhao, Y. et al. A Conductive Binder for High-Performance Sn Electrodes in Lithium-Ion Batteries. ACS Appl. Mater. Interfaces 10, 1672–1677 (2018). 31. Lai, C.-H. et al. Application of Poly(3-hexylthiophene-2,5-diyl) as a Protective Coating for High Rate Cathode Materials. Chem. Mater. 30, 2589–2599 (2018). 32. Elizalde-Segovia, R. et al. Understanding the Role of π-Conjugated Polymers as Binders in Enabling Designs for High-Energy/High-Rate Lithium Metal Batteries. J. Electrochem. Soc. (2021) doi:10.1149/1945- 7111/ac3850. 33. Fu, Y. & Manthiram, A. Enhanced Cyclability of Lithium–Sulfur Batteries by a Polymer Acid-Doped Polypyrrole Mixed Ionic–Electronic Conductor. Chem. Mater. 24, 3081–3087 (2012). 34. Park, S.-J. et al. Side-Chain Conducting and Phase-Separated Polymeric Binders for High- Performance Silicon Anodes in Lithium-Ion Batteries. J. Am. Chem. Soc. 137, 2565–2571 (2015). 35. Wu, M. et al. Toward an Ideal Polymer Binder Design for High-Capacity Battery Anodes. J. Am. Chem. Soc. 135, 12048–12056 (2013). 36. Patel, S. N., Javier, A. E., Stone, G. M., Mullin, S. A. & Balsara, N. P. Simultaneous Conduction of Electronic Charge and Lithium Ions in Block Copolymers. ACS Nano 6, 1589–1600 (2012). 37. McDonald, M. B. & Hammond, P. T. Efficient Transport Networks in a Dual Electron/Lithium- Conducting Polymeric Composite for Electrochemical Applications. ACS Appl. Mater. Interfaces 10, 15681– 15690 (2018). 38. Zeng, W. et al. Enhanced Ion Conductivity in Conducting Polymer Binder for High-Performance Silicon Anodes in Advanced Lithium-Ion Batteries. Adv. Energy Mater. 8, 1702314 (2018). 225 39. Minnici, K. et al. Tuning Conjugated Polymers for Binder Applications in High-Capacity Magnetite Anodes. ACS Appl. Energy Mater. 2, 7584–7593 (2019). 40. Ghosh, S. Networks of Electron-Conducting Polymer in Matrices of Ion-Conducting Polymers Applications to Fast Electrodes. Electrochem. Solid-State Lett. 3, 213 (1999). 41. Homann, G. et al. Poly(Ethylene Oxide)-based Electrolyte for Solid-State-Lithium-Batteries with High Voltage Positive Electrodes: Evaluating the Role of Electrolyte Oxidation in Rapid Cell Failure. Sci. Rep. 10, 4390 (2020). 42. Kwon, Y. H. et al. Electron/Ion Transport Enhancer in High Capacity Li-Ion Battery Anodes. Chem. Mater. 28, 6689–6697 (2016). 43. Minnici, K. et al. Carboxylated Poly(thiophene) Binders for High-Performance Magnetite Anodes: Impact of Cation Structure. ACS Appl. Mater. Interfaces 11, 44046–44057 (2019). 44. Oriňáková, R., Fedorková, A. & Oriňák, A. Effect of PPy/PEG conducting polymer film on electrochemical performance of LiFePO4 cathode material for Li-ion batteries. Chem. Pap. 67, (2013). 45. Das, P. et al. Dihexyl-Substituted Poly(3,4-Propylenedioxythiophene) as a Dual Ionic and Electronic Conductive Cathode Binder for Lithium-Ion Batteries. Chem. Mater. 32, 9176–9189 (2020). 46. Ren, X. & Pickup, P. G. Ionic and Electronic Conductivity of Poly‐(3‐methylpyrrole‐4‐carboxylic Acid). J. Electrochem. Soc. 139, 2097 (1992). 47. Ren, X. & Pickup, P. G. Coupling of ion and electron transport during impedance measurements on a conducting polymer with similar ionic and electronic conductivities. J. Chem. Soc. Faraday Trans. 89, 321 (1993). 48. Savagian, L. R. et al. Balancing Charge Storage and Mobility in an Oligo(Ether) Functionalized Dioxythiophene Copolymer for Organic- and Aqueous- Based Electrochemical Devices and Transistors. Adv. Mater. 30, 1804647 (2018). 49. Chen, X. et al. n-Type Rigid Semiconducting Polymers Bearing Oligo(Ethylene Glycol) Side Chains for High-Performance Organic Electrochemical Transistors. Angew. Chem. Int. Ed. 60, 9368–9373 (2021). 50. Dong, B. X. et al. Influence of Side-Chain Chemistry on Structure and Ionic Conduction Characteristics of Polythiophene Derivatives: A Computational and Experimental Study. Chem. Mater. 31, 1418–1429 (2019). 51. Meng, B. et al. Replacing Alkyl with Oligo(ethylene glycol) as Side Chains of Conjugated Polymers for Close π–π Stacking. Macromolecules 48, 4357–4363 (2015). 52. Chen, X., Zhang, Z., Ding, Z., Liu, J. & Wang, L. Diketopyrrolopyrrole-based Conjugated Polymers Bearing Branched Oligo(Ethylene Glycol) Side Chains for Photovoltaic Devices. Angew. Chem. Int. Ed. 55, 10376–10380 (2016). 53. Moser, M. et al. Side Chain Redistribution as a Strategy to Boost Organic Electrochemical Transistor Performance and Stability. Adv. Mater. 32, 2002748 (2020). 54. Moser, M. et al. Ethylene Glycol-Based Side Chain Length Engineering in Polythiophenes and its Impact on Organic Electrochemical Transistor Performance. Chem. Mater. 32, 6618–6628 (2020). 226 55. Ye, L. & Thompson, B. C. Improving the efficiency and sustainability of catalysts for direct arylation polymerization (DArP). J. Polym. Sci. n/a,. 56. Pankow, R. M. & Thompson, B. C. Approaches for improving the sustainability of conjugated polymer synthesis using direct arylation polymerization (DArP). Polym. Chem. 11 (2020). 57. Rudenko, A. E., Khlyabich, P. P. & Thompson, B. C. Random Poly(3-hexylthiophene-co-3- cyanothiophene) Copolymers via Direct Arylation Polymerization (DArP) for Organic Solar Cells with High Open-Circuit Voltage. ACS Macro Lett. 6 (2014). 58. Rudenko, A. E., Wiley, C. A., Stone, S. M., Tannaci, J. F. & Thompson, B. C. Semi-random P3HT analogs via direct arylation polymerization. J. Polym. Sci. Part Polym. Chem. 50, 3691–3697 (2012). 59. Choi, C. et al. Achieving high energy density and high power density with pseudocapacitive materials. Nat. Rev. Mater. 5, 5–19 (2020). 60. Jiang, X.; Harima, Y.; Yamashita, K.; Tada, Y.; Ohshita, J.; Kunai, A. Doping-induced change of carrier mobilities in poly(3-hexylthiophene) films with different stacking structures. Chem. Phys. Lett. 2002, 364, 616– 620, DOI: 10.1016/S0009-2614(02)01383-0 61. Chance, R. R., Brédas, J. L. & Silbey, R. Bipolaron transport in doped conjugated polymers. Phys. Rev. B 29, 4491–4495 (1984). 62. Bargigia, I., Savagian, L. R., Österholm, A. M., Reynolds, J. R. & Silva, C. Charge-Transfer Intermediates in the Electrochemical Doping Mechanism of Conjugated Polymers. J. Am. Chem. Soc. 143, 294–308 (2021). 63. Kivelson, S. & Heeger, A. J. First-order transition to a metallic state in polyacetylene: A strong- coupling polaronic metal. Phys. Rev. Lett. 55, 4 (1985). 64. Miller, L. L. & Mann, K. R. π-Dimers and π-Stacks in Solution and in Conducting Polymers. 7. 65. Haare, J. A. E. H. van et al. Redox States of Long Oligothiophenes: Two Polarons on a Single Chain. Chem. – Eur. J. 4, 1509–1522 (1998). 66. Enengl, C. et al. Doping-Induced Absorption Bands in P3HT: Polarons and Bipolarons. ChemPhysChem 17, 3836–3844 (2016). 67. Nowak, M. J., Spiegel, D., Hotta, S., Heeger, A. J. & Pincus, P. A. Charge storage on a conducting polymer in solution. 22, 10 (1989). 68. Reeves, B. D. et al. Spray Coatable Electrochromic Dioxythiophene Polymers with High Coloration Efficiencies. Macromolecules 37, 7559–7569 (2004). 69. Kumar, A. et al. Conducting Poly(3,4-alkylenedioxythiophene) Derivatives as Fast Electrochromics with High-Contrast Ratios. Chem. Mater. 10, 896–902 (1998). 70. Welsh, D. M. et al. Regiosymmetric Dibutyl-Substituted Poly(3,4-propylenedioxythiophene)s as Highly Electron-Rich Electroactive and Luminescent Polymers. Macromolecules 35, 6517–6525 (2002). 71. Lin, C., Endo, T., Takase, M., Iyoda, M. & Nishinaga, T. Structural, Optical, and Electronic Properties of a Series of 3,4-Propylenedioxythiophene Oligomers in Neutral and Various Oxidation States. J. Am. Chem. Soc. 133, 11339–11350 (2011). 227 72. Yamamoto, J. & Furukawa, Y. Electronic and Vibrational Spectra of Positive Polarons and Bipolarons in Regioregular Poly(3-hexylthiophene) Doped with Ferric Chloride. J. Phys. Chem. B 119, 4788– 4794 (2015). 73. Zayat, B., Das, P., Thompson, B. C. & Narayan, S. R. In Situ Measurement of Ionic and Electronic Conductivities of Conductive Polymers as a Function of Electrochemical Doping in Battery Electrolytes. J. Phys. Chem. C 125, 7533–7541 (2021). 74. Sahalianov, I. et al. UV-to-IR Absorption of Molecularly p-Doped Polythiophenes with Alkyl and Oligoether Side Chains: Experiment and Interpretation Based on Density Functional Theory. J. Phys. Chem. B 124, 11280–11293 (2020). 75. Lall-Ramnarine, S. I. et al. Connecting Structural and Transport Properties of Ionic Liquids with Cationic Oligoether Chains. J. Electrochem. Soc. 164, H5247–H5262 (2017). 76. Duong, D. T., Wang, C., Antono, E., Toney, M. F. & Salleo, A. The chemical and structural origin of efficient p-type doping in P3HT. Org. Electron. 14, 1330–1336 (2013). 77. Müller-Buschbaum, P. The Active Layer Morphology of Organic Solar Cells Probed with Grazing Incidence Scattering Techniques. Adv. Mater. 26, 7692–7709 (2014). 78. Yee, P. Y., Scholes, D. T., Schwartz, B. J. & Tolbert, S. H. Dopant-Induced Ordering of Amorphous Regions in Regiorandom P3HT. J. Phys. Chem. Lett. 10, 4929–4934 (2019). 79. Scholes, D. T. et al. The Effects of Crystallinity on Charge Transport and the Structure of Sequentially Processed F 4 TCNQ‐Doped Conjugated Polymer Films. Adv. Funct. Mater. 27, 1702654 (2017). 80. Hamidi-Sakr, A. et al. A Versatile Method to Fabricate Highly In-Plane Aligned Conducting Polymer Films with Anisotropic Charge Transport and Thermoelectric Properties: The Key Role of Alkyl Side Chain Layers on the Doping Mechanism. Adv. Funct. Mater. 27, 1700173 (2017). 81. Stanfield, D. A., Wu, Y., Tolbert, S. H. & Schwartz, B. J. Controlling the Formation of Charge Transfer Complexes in Chemically Doped Semiconducting Polymers. Chem. Mater. 33, 2343–2356 (2021). 82. Bennington, P. et al. Role of solvation site segmental dynamics on ion transport in ethylene-oxide based side-chain polymer electrolytes. J. Mater. Chem. A 9, 9937–9951 (2021). 83. Deng, C. et al. Role of Molecular Architecture on Ion Transport in Ethylene oxide-Based Polymer Electrolytes. Macromolecules 54, 2266–2276 (2021). 84. Song, J. & Bazant, M. Z. Effects of Nanoparticle Geometry and Size Distribution on Diffusion Impedance of Battery Electrodes. J. Electrochem. Soc. 160, A15–A24 (2013). 85. Lasia, A. Electrochemical Impedance Spectroscopy and its Applications. (Springer New York, 2014). doi:10.1007/978-1-4614-8933-7. 86. Montoro, L. A.; Rosolen, J. M. The Role of Structural and Electronic Alterations on the Lithium Diffusion in LixCo0.5Ni0.5O2. Electrochimica Acta 2004, 49 (19), 3243–3249. https://doi.org/10.1016/j.electacta.2004.03.001. 87. Song, J. & Bazant, M. Z. Effects of Nanoparticle Geometry and Size Distribution on Diffusion Impedance of Battery Electrodes. J. Electrochem. Soc. 160, A15–A24 (2013). 228 88. Li, J. et al. Dependence of Cell Failure on Cut-Off Voltage Ranges and Observation of Kinetic Hindrance in LiNi 0.8 Co 0.15 Al 0.05 O 2. J. Electrochem. Soc. 165, A2682–A2695 (2018). 89. Zhu, X.-J. et al. Preparation and characterization of LiNi0.80Co0.20– x Al x O2 as cathode materials for lithium ion batteries. J. Electroceramics 17, 645–649 (2006). 90. Nam, G. W. et al. Capacity Fading of Ni-Rich NCA Cathodes: Effect of Microcracking Extent. ACS Energy Lett. 4, 2995–3001 (2019). 91. Park, K.-J. et al. Degradation Mechanism of Ni-Enriched NCA Cathode for Lithium Batteries: Are Microcracks Really Critical? ACS Energy Lett. 4, 1394–1400 (2019). 92. Ryu, H.-H., Park, K.-J., Yoon, C. S. & Sun, Y.-K. Capacity Fading of Ni-Rich Li[Ni x Co y Mn 1– x – y ]O 2 (0.6 ≤ x ≤ 0.95) Cathodes for High-Energy-Density Lithium-Ion Batteries: Bulk or Surface Degradation? Chem. Mater. 30, 1155–1163 (2018). 93. Natarajan, S., Moodakare, S. B., Haridoss, P. & Gopalan, R. Concentration Gradient-Driven Aluminum Diffusion in a Single-Step Coprecipitation of a Compositionally Graded Precursor for LiNi 0.8 Co 0.135 Al 0.065 O 2 with Mitigated Irreversibility of H2 ↔ H3 Phase Transition. ACS Appl. Mater. Interfaces 12, 34959–34970 (2020). 229 Chapter 6 - NaNb 13 O 33 as an anode material for lithium-ion Batteries Abstract Sodium niobium oxide (NaNb 13O 33 ), a Wadsley-Roth phase stands out as a promising anode for lithium-ion batteries due to its high specific capacity of 220 mAh g -1 at C/15 and its lower nominal voltage of approximately 1.55 V vs Li + /Li. Rate capability, cycle-life, incremental capacity analysis (ICA), electrochemical impedance spectroscopy (EIS), and determination of the apparent diffusion coefficient was carried to understand the electrochemical proprieties and performance of NaNb 13O 33 in a half-cell. ICA in conjunction with EIS analysis provides a clear understanding of the reversible phase transitions that NaNb 13O 33 undergoes during the lithiation and delithiation. The data shows that NaNb 13O 3 has an impressive rate capability of 80 mAh g -1 at 20C and a remarkable capacity retention of 80% after 600 cycles at 1C. Moreover, NaNb 13O 33 was paired against NCA-PVDF and NCA-PProDOT-Hx 2 to build a Full-cell and to characterize its performance as an anode. Finally, we have built a pouch-cell to demonstrate the feasibility of building a practical lithium-ion battery with NaNb 13O 3 anode. Incremental capacity analysis Figure 6.1a,b shows the galvanostatic profile of NaNb 13O 33 during lithiation and delithiation in the potential range of 3 to 1 V vs Li + /Li. Both profiles exhibited at least eight different regions (Labeled with Roman numerals) indicating the sequence of transitions between various single phases and solid-solution compositions corresponding to the sloping regions of the charge/discharge curve. To understand the phase transitions that NaNb 13O 33 undergoes during charge and discharge incremental capacity analysis (ICA) was utilized. Figure 6.1c shows five 230 peaks and three valleys that have been identified by ICA. Each peak is related to a phase transition that the active material experiences during lithiation/delithiation. 1–3 In accordance to the operando XRD characterization ( carried out by Prof. Ram Seshadri’s group at UC Santa Barbara) and ICA analysis, region I corresponds to a monoclinic phase M1, region II (or first peak) at 1.66 V vs Li + /Li is related to the phase transition of M1 to an hexagonal H1 phase (❶M1+H1), the second peak at 1.64 V a transition from H1 to H2 (❷H1+H2), the third peak at 1.58 V the transition to H3 (❸H2+H3), the fifth region H3, the peak at 1.38 V to H4 (❹H3+H4), and finally the peak at 1.14 V is associated to a fifth hexagonal H5 phase transition (❺H4+H5). However, during delithiation peak ❷ is not present attributed to the merging with peak ❶ suggesting the co-existence of phases at same potential during charge (Figure 6.1c). 231 Figure 6.1 Li xNaNb 13O 33||Li cell galvanostatic charge/discharge (GCD) profile and differential capacity analysis. a) Galvanostatic profile during lithiation and b) delithiation between 1 and 3 V vs Li + /Li at C/15. c) Corresponding dQ dV -1 vs V plot for the GCD curves in a) and b). regions are labeled with Roman numerals and the phase transitions (peaks) with Arabic numerals. 232 Rate Capability and incremental capacity analysis. Specific capacity as a function of C-rate and cycle number was characterized for the NaNb 13O 33 in a cell with a lithium anode (Figure 6.2a). At 6C and 20C the cell delivered an impressive specific capacity of 108 mAh g -1 and 80 mAh g -1 , respectively, based on the active material loading, confirming the fast Li + injection and removal capability of NaNb 13O 33. Furthermore, at higher mass loadings of the active material (5.5 mg cm -2 ) the cells exhibited a remarkable specific capacity of 95 mAh g -1 at 10C (Figure 6.2a). We note that at low and moderate rates up to 2C the GCD curves (Figure 6.3a,b) show various plateaus but at higher rates the plateaus were not observed suggesting that at higher currents the distinct phase transitions are not observable. Thus, to understand the mechanisms that affect the electrode, ICA was carried for each rate. Figure 6.2b shows a three-dimensional representation of an ICA analysis for GCD curves in Figure 6.3a (2D plot Figure 6.4). 233 Figure 6.2 Rate Capability and incremental capacity analysis of a Li xNaNb 13O 33||Li cell. a) Specific capacity as function of cycle number with increments in C-rate. b) Three-dimensional ICA plot from the GCD curves from Figure 6.3. 234 Figure 6.3 Representative galvanostatic charge-discharge curves at various C-rates for Li xNaNb 13O 33||Li cells with an areal active material mass loading of a) 2 mg cm -2 and b) 5.5 mg cm -2 . Figure 6.4 Incremental capacity analysis (ICA) as a function C-rate (two-dimensional plot) from the GCD curves in Figure 6.3 Identified peaks are denote with numeric labels. 235 To characterize the ICA profiles we conducted a detailed peak analysis. The peaks were identified using the analysis tool “Peak Analyzer” from OriginPro-2021b. The peak position and height are defined at the voltage at which the local maximum absolute dQ/dV value occurs. Peak area is calculated using local maximum method across the width of the peak centered at the peak position where the peak width corresponds to the full width at half maximum (FWHM). Figure 6.5a,e shows the position of the identified peaks as a function or rate during lithiation and delithiation respectively. At C/15 all peaks are present confirming that the active material undergoes all its phase transitions. As the rate is increased, during lithiation peaks ❺, ❷, ❹ and ❸ are suppressed at C/2, 2C, and 6C respectively. During delithiation at C/2, 2C and 2C, the same was observed. The use of large currents can lead to non-equilibrium distribution of lithium through the active material resulting in the co-existence of phases as indicated by the disappearance of peaks and width increase (Figure 6.5d,h). All peaks shift toward higher polarization during charge and discharge. This shift results in the charge and discharge cut-off voltages being prematurely reached without achieving full removal or insertion of lithium. The shifts in polarization have similar magnitudes during lithiation and delithiation suggesting that the peak changes from ohmic and faradaic losses are similar during charge and discharge. 1 Furthermore, the analysis shows an inverse relation between C-rate and peak height (Figure 6.5b,f) indicating the lack of full equilibration of lithium distribution at higher currents. Finally, the increase in peak area and width at higher rates are attributed to the co-existence of phases and higher polarization leading to a proportional increase of capacity as the peaks merge and extend (Figure 6.5c,d,g,h). 236 Figure 6.5. Li xNaNb 13O 33||Li cell peak analysis as function of C-rate during lithiation a) Peak position, b) Peak height, c) peak area, d) peak width. Long-term cycling and EIS characterization. Long-term cycling of the NaNb 13O 33 lithium metal cell after two formation cycles at C/10 was conducted at a charge/discharge rate of 2C in the potential window of 1 to 3 V vs Li + /Li. After 600 cycles the cells had an impressive capacity retention of 80% (Figure 6.6a). As the number of cycles is increased the features of the GCD curves (Figure 6.7) are not as pronounced compared to the first few cycles suggesting material transformation upon cycling. To understand the degradation mechanisms occurring in the electrode, an ICA was performed in conjunction with EIS characterization as function of potential during the 1 st and 600 th cycles. Figure 6.6b shows a 237 three-dimensional representation of the ICA plots as function of cycle number from the GCD curves in Figure 6.7 (2D plot Figure 6.8). Figure 6.6 Li xNaNb 13O 33||Li cell specific capacity as a function of cycle number and differential capacity analysis. a) long-term cycling between 1 and 3 V vs Li + /Li at 2C. b) Three-dimensional ICA plot from the GCD curves in Figure 6.7 238 Figure 6.7 Representative galvanostatic charge-discharge curves as function of cycle number at 2C for a Li xNaNb 13O 33||Li cell with an areal active material mass loading of 2.2 mg cm -2 . Figure 6.8 Incremental capacity analysis (ICA) as a function of cycle number (two-dimensional plot) from the GCD curves in Figure 6.7 Identified peaks are denote with numeric labels. 239 The analysis shows that during the first cycle all phase transitions are present, but after 100 th cycles peaks ❷ and ❺ vanish indicating their non-reversibility upon cycling. In terms of peak position, a significant shift in polarization is present from the 1 st to the 100 th cycle, but no change is observed from the 100 th to 500 th (Figure 6.9a,e) suggesting that the main degradation occurs during the first 100 th cycles. Figure 6.9 Li xNaNb 13O 33||Li cell peak analysis as function of C-rate during lithiation a) Peak position, b) Peak height, c) peak area, d) peak width. The ohmic and faradaic polarization losses are symmetric between the charge and discharge. Peak height for both lithiation and delithiation shows a significant reduction from the 1 st to the 200 th cycle except for peak ❹ that remains constant (Figure 6.9b,f) indicating that the main loss 240 of active material occurs in phase transitions ❶ and ❸ during the first 200 th cycles consistent with the increase of peak area and width (Figure 6.9c,g and 6.9d,h). EIS was measured as function of potential through all the discharge profile at open circuit voltage (OCV) (Figure 6.10). Figure 6.11a and 6.11 shows a three-dimensional representation of the Nyquist plots for the 1 st and 600 th cycle respectively. Overall, during the first cycle the complex impedance plots are composed by one semicircle and a diffusion tail at low frequencies. The main changes through lithiation are seen in the diffusional tail between 1.7 to 1.2 V vs Li + /Li where there is an overall decrease in magnitude. Nevertheless, after 600 th cycles the Nyquist plots showed an increase in the extent of the diffusional tail when reaching lower potentials. To interpret the impedance response a Randles-type equivalent electrical circuit (EEC) was proposed (Figure 6.12a). Figure 6.10 Galvanostatic lithiation curve for the Li xNaNb 13O 33||Li cell at C/20. EIS measurements at open circuit voltage are denoted with yellow circles. 241 Figure 6.11 Electrochemical impedance characterization as a function of potential vs Li + /Li during lithiation plotted in 3-D Nyquist plots for a) the 1 st cycle after formation, and b) after 600 th cycles 242 Figure 6.12 a) Equivalent electrical circuit used for fitting the impedance data in Figure #. Change of the b) ohmic electrolyte resistance R1, c) the constant phase element Q1, d) a value, e) charge transfer resistance R2, and f) Warburg coefficient S1 as function of potential vs Li+/Li for the 1st (denoted with green dots) and 600th cycle (denoted with purple dots). Relative fitting errors are indicated with error bars. 243 The electrical equivalent circuit (EEC) is composed by the electrode and electrolyte resistance (R1), a constant phase element associated to the double layer capacitance (Q1) in parallel with a charge-transfer resistance (R2), and by a semi-infinite Warburg element (W) associated to the intercalation of lithium ions. The fitting parameters and their corresponding fitting errors are shown in Tables 6.1 and 6.2. After the 600 th cycle, the differences in the ohmic resistance are not as drastic (2 ohms) when compared to the 1 st (Figure 6.12b) indicating the degradation mechanism does not arise from electrolyte decomposition. Figure 6.12c shows the change for Q1 as function of potential where the exponent a is associated with the level of dispersion that the impedance response exhibits. An ideal capacitor has an a value equal to 1. During the 1 st cycle a had a value of approximately 0.8 indicating that there is not significant deviation from an ideal capacitor (Figure 6.12d). Two orders of higher magnitude for Q1 is observed when compared to the 600 th cycle where the a value ranges between 0.51 to 0.37 suggesting a prominent deviation from the ideal capacitor behavior attributable to changes of morphology and loss of inter- particulate contact electrode during cycling. Figure 6.12e shows that R2 did not exhibit a significant change during lithiation for the 1 st cycle, but it is 20 ohms higher on average than in the 600 th cycle that had a slight decrease in magnitude bellow 1.4 V vs Li + /Li. This increase in charge-transfer resistance with cycling is consistent with the changes in morphology suggested by changes in Q1 value and the reduced material utilization after long-term cycling (Figure 6.12a,e). Figure 6.12f shows that impedance associated with the Warburg element increases at higher levels of lithiation and cycle number. At lower potentials there is a limited number of lattice sites for the Li + diffusion into the active material. The analysis indicates that below 1.4 V vs Li + /Li the lithium-ion insertion is hindered after long cycling consistent with the disappearance 244 of peaks at low potentials in the ICA analysis (Figure 6.6b). Imposing a discharge cut-off up to 1.4 V vs Li + /Li could in principle increase the capacity retention during long-term cycling although almost ¼ of the capacity would be sacrificed. Therefore, the changes to the mass-transfer process (observed at low frequencies) appear to explain the performance decrease of NaNb 13O 33 after 600 th cycles. In other words, after long-term cycling irreversible changes to the active material composition and morphology below 1.4 V vs Li + /Li that hinders the full lithiation and occurrence of the phase transitions. Table 6.1 Fitting parameters and their corresponding relative error for the 1 st cycle of the Li xNaNb 13O 33||Li cell as a function of potential vs Li + /Li. POTENTIAL (V) R1 (OHM) DEV R1 Q1 (F.S^(A - 1)) DEV Q1 A1 () R2 (OHM) DEV R2 S1 (OHM.S^- 1/2) DEV S1 1.81107 3.40746 0.40332 1.96E-05 3.94E-06 0.814 67.51326 0.65237 5.61947 0.28515 1.70167 3.37415 0.39627 1.92E-05 3.83E-06 0.81612 68.06262 0.65226 3.99679 0.28489 1.67678 3.32364 0.40198 1.96E-05 3.92E-06 0.81426 67.60691 0.65265 4.50805 0.28516 1.67425 3.30896 0.38292 1.96E-05 3.83E-06 0.81393 68.98033 0.66471 5.02297 0.28692 1.66106 3.29866 0.35461 1.95E-05 3.69E-06 0.8146 70.7789 0.67893 5.25402 0.28906 1.63017 3.25928 0.38811 1.96E-05 3.89E-06 0.81446 68.16614 0.66058 5.05974 0.28626 1.60354 3.24825 0.35201 1.94E-05 3.68E-06 0.8158 70.60339 0.67787 4.8114 0.28879 1.60118 3.18295 0.36747 1.98E-05 3.83E-06 0.81374 69.45968 0.67377 5.60888 0.28831 1.59263 3.17966 0.3398 1.96E-05 3.68E-06 0.81474 71.22752 0.68619 6.35644 0.29026 1.51372 3.12262 0.33912 1.94E-05 3.67E-06 0.81632 70.87341 0.68341 7.00843 0.28968 1.43903 2.97698 0.32281 2.01E-05 3.72E-06 0.81289 71.75595 0.69803 7.54655 0.29245 1.40507 2.9153 0.32687 2.06E-05 3.87E-06 0.81061 70.98104 0.70067 8.38287 0.29304 1.37982 2.87869 0.31657 2.09E-05 3.89E-06 0.80963 71.33279 0.70736 9.95553 0.2943 1.34334 2.86732 0.31221 2.09E-05 3.86E-06 0.80946 71.75 0.70963 11.81652 0.29474 1.29469 2.83332 0.32126 2.10E-05 3.94E-06 0.80944 70.85195 0.7056 13.43812 0.29399 1.22965 2.86286 0.19445 2.00E-05 3.28E-06 0.81373 79.271 0.74288 18.185 0.30214 245 1.18454 2.80492 0.32634 2.05E-05 3.92E-06 0.81331 69.71036 0.69555 22.40561 0.29198 1.14633 2.89097 0.33391 1.93E-05 3.74E-06 0.8207 69.03864 0.67746 29.43201 0.28836 1.0976 2.98091 0.38255 1.79E-05 3.73E-06 0.82969 65.40832 0.63328 36.01122 0.28066 Table 6.2 Fitting parameters and their corresponding relative error for the 600 th cycle of the Li xNaNb 13O 33||Li cell as a function of potential vs Li + /Li. POTENTIAL (V) R1 (OHM) DEV R1 Q1 (F.S^(A - 1)) DEV Q1 A1 () R2 (OHM) DEV R2 S1 (OHM.S^- 1/2) DEV S1 1.72066 7.90932 0.88697 9.42E-04 3.15E-04 0.51516 33.97314 2.32019 5.85532 0.68005 1.68032 7.78676 0.90736 9.68E-04 3.25E-04 0.51072 34.01954 2.37201 6.05933 0.68615 1.68012 7.62174 0.95601 0.00106 3.50E-04 0.49906 34.88001 2.52055 6.79313 0.71979 1.66957 7.68009 0.9591 0.0011 3.46E-04 0.49637 36.11083 2.5599 7.28988 0.74309 1.63419 7.65562 0.94993 0.00106 3.49E-04 0.50002 34.85757 2.50836 6.8842 0.72057 1.61217 7.34727 1.05524 0.00128 3.94E-04 0.47596 37.47846 2.86464 7.90457 0.80318 1.60795 7.00562 1.15948 0.00146 4.48E-04 0.45737 38.48084 3.194 9.37708 0.85896 1.59393 6.94281 1.19236 0.00152 4.69E-04 0.45264 38.37802 3.28554 10.67591 0.869 1.50147 6.72663 1.27384 0.00162 5.14E-04 0.44277 37.86998 3.48693 13.05517 0.88122 1.43312 6.38013 1.41022 0.00189 5.78E-04 0.42206 40.30181 4.00415 12.96903 0.97642 1.40034 5.86734 1.63212 0.00227 6.83E-04 0.39616 42.77848 4.82977 15.8514 1.09597 1.36721 5.17624 1.97582 0.00265 8.56E-04 0.37169 42.23481 5.80778 21.97804 1.15437 1.32329 5.85836 1.67658 0.0021 7.43E-04 0.40449 36.24436 4.48072 31.77103 0.9426 1.26368 6.99884 1.17739 0.00132 5.54E-04 0.46985 28.92415 2.94616 43.17079 0.69492 1.19269 7.56738 0.91316 9.56E-04 4.68E-04 0.51383 24.66879 2.3194 52.62984 0.57557 1.15476 7.858 0.74506 7.69E-04 4.17E-04 0.5428 22.3987 2.00234 59.659 0.5154 1.11769 8.1576 0.59613 6.32E-04 3.72E-04 0.56819 20.83237 1.76764 66.17982 0.47247 246 Apparent diffusion coefficient by EIS. The impedance of the Warburg element is inversely related to the diffusion coefficient. Assuming semi- infinite conditions and using the geometrical electrode area, the apparent diffusion coefficient (D Li) can be calculated using the following equation 4,5 : 𝐷 𝐿𝑖 + = 1 2 [ 𝑉 𝑀 𝐴𝐹 𝜎 𝑤 ( 𝛿𝐸 𝛿𝑥 ) ] 2 Where 𝛿𝐸 𝛿𝑥 is the slope of the open circuit voltage vs Li + concentration, F is the Faraday constant (96490 C mol -1 ), V M is the molar volume (382.34 cm 3 mol -1 ), A is the geometrical electrode area, and 𝜎 𝑤 is the Warburg coefficient obtained from the slope of Zre vs. ω -0.5 or -Zim vs. ω -0.5 plot (Figure 6.13), where ω corresponds to the angular frequency. However, 𝛿𝐸 𝛿𝑥 is close to zero in any two-phase regions (plateaus) thus the D Li is undefined in the two plateaus that the GCD profile exhibits at 1.66 V and 1.59 V vs Li + /Li. Figure 6.14 shows the D Li value as function of potential without the undefined regions. The D Li decreases at lower potentials from 1.4×10 -11 to 1. 1 ×10 -12 cm 2 s -1 at 2.1 V and 1.2 V vs Li + /Li respectively, consistent with the observed changes in W (Figure 6.12f). Furthermore after 600 cycles there is a decrease in the D Li value being 3.6×10 -13 cm 2 s -1 at 1.2 V vs Li + /Li following the same trend as the material gets lithiated. The differences in the apparent diffusion coefficient after 600 cycles can be attributed to changes in the interparticulate contact consistent with the impedance and ICA analysis. 247 Figure 6.13 Real part of the impedance response Zre in the Warburg region as a function of the inverse square of the angular frequency ω -0.5 and their linear fitting for each potential vs Li + /Li during lithiation. Figure 6.14 Apparent diffusion coefficient by EIS during lithiation for the 1 st and 600 th cycle assuming semi-infinite conditions. 248 Full-Cell configuration: NaNb13O33||NCA and NaNb13O33||NCA-PProDOT-Hx2 NaNb 13O 33 was paired against NCA to assemble a Full-cell. The following parameters need to be stablished beforehand: The negative to positive electrode capacity ratio (N/P), reversible capacity of both electrodes at the same current density, defining C-Rate, determination of the cell’s potential window, and stablishing a SEI formation protocol. Specifically, the N/P ratio needs to be higher than 1 to prevent lithium plating on the negative electrode. Due to the slight excess of anode material (NaNb 13O 33) the cell’s overall capacity is limited by the positive electrode. The N/P ratio can be calculated utilizing the theoretical capacity of the active materials, but this value is often inadequate for a Full-cell because pragmatic utilization of the electrodes is not the same at all discharge/charge rates and it diminishes after cycling. Therefore, “reversible” capacity of the electrodes is used instead. Reversible capacity is defined as the maximum achievable capacity at 1C based on experimental data for each of the electrodes in a half-cell configuration (against metallic lithium). The half-cells need to have the same electrode composition, active material loading (thickness), and electrolyte as in a Full-cell. Capacity of a half-cell is mainly affected by the electrode composition, loading, electrode porosity/tortuosity, temperature, the electrolyte conductivity/stability, and SEI conductivity. The potential window of a Full-cell can be calculated by utilizing the experimental data from both active materials half-cells by subtracting the observed potential of the negative electrode from the positive electrode during lithiation/delithiation and delithiation/lithiation. Nevertheless, the experimental GCD profiles of a Full-cell tend to differ from the calculated values at moderate and higher rates due to the differences in the electrode’s rate capability. Therefore, it is suggested to utilize window at extremely low rates, or let the cell relax at OCV during an intermittent galvanostatic titration protocol to find the thermodynamic potential at each state of charge. The potential cut-off is stablished at the end of the maximum achievable capacity from the 249 limiting electrode (capacity known from the half-cell studies). The experimental potential window of the resulting NCA||NaNb 13O 33 Full-cell goes from 3 to 1.15 V (Figure 6.15). Table 6.3 shows the resumes the electrode composition and active material loading utilized in the Full-cell and half-cell studies. Figure 6.15 Galvanostatic profiles for NCA||Li (yellow), NaNb 13O 33||Li (red) and Full-cell paring of both materials NCA|| NaNb 13O 33 (purple). Table 6.3 Electrodes and electrolyte composition utilized in the NCA|| NaNb 13O 33 Full-cell. Component Composition Negative Electrode 80% Li XNaNb 13O 33 10% PVDF 10% Super P carbon clack Active material loading ≈ 5 mg cm -2 20 µm Cu foil (current collector) Positive Electrode 90% LiNi 0.8Co 0.15Al 0.05O 2 (NCA) 3% MWCNTs 3% Super P 4% Binder (PVDF or PProDOT-Hx 2) Active material loading ≈ 4 mg cm -2 15 µm Al foil (current collector) Electrolyte 1M LiTFSI in EC:DMC (1:1 v/v) 250 Rate capability studies of the half-cells showed a maximum experimental capacity of approximately 180 mAh g -1 for NCA-PProDOT-Hx 2 and 200 mAh g -1 for NaNb 13O 33 at low currents. At 1C the cells delivered a reversible capacity of 165 and 160 mAh g -1 for NCA-PProDOT-Hx 2 and for NaNb 13O 33 respectively. C- rate Is based on the reversible capacity of the limiting electrode (NCA), therefore, 1C = 165 mA g -1 (Figure 6.16). Figure 6.16 Rate Capability testing and their correspond galvanostatic profiles for (a-b) NCA-PProDOT- Hx 2||Li and (c-d) NaNb 13O 33||Li cells. Rate Capability of NCA-PProDOT-Hx 2||NaNb 13O 33 (Full-cell) PProDOT-Hx 2 was utilized as a conducting binder for the NCA electrodes (Chapter 3-5) and PVDF for the NaNb 13O 33 electrodes (Table 6.3). Figure 6.17 shows the rate capability of the NCA-PProDOT- 251 Hx 2||NaNb 13O 33 cell. Three cycles at C/10 were stablished for the SEI formation at room temperature, with a N/P ratio of 1.3 and a NCA areal loading of approximately 4 mg cm -2 . At 4C an impressive specific capacity of 75 mAh g -1 was obtained (Figure 6.18). The high rate capability of the Full-cell arises from the fast lithium insertion/extraction that the active materials can provide and the enhanced ionic and electronic conductivity of PProDOT-Hx 2, demonstrating the feasibility of building a high-power battery. Figure 6.17 a) Rate capability and b) corresponding GCD curves for the NCA-PProDOT-Hx 2||NaNb 13O 33 cell with a N/P ratio of 1.3 an areal loading of 4 mg cm -2 . At 1C cycle-life of the NCA-PProDOT-Hx 2||NaNb 13O 33 cell was characterized and compared against a NCA-PVDF||NaNb 13O 33 cell. After 70 cycles the NCA-PProDOT-Hx 2 cell had a specific capacity of 80 mAh g -1 with a capacity retention of 60%, in contrast, the NCA-PVDF cell showed a specific capacity of 40 mAh g -1 and a 51% capacity retention (Figure 6.18). The difference in performance is attributed to the NCA electrode because it limits the cell’s capacity (N/P = 1.3) and rate capability (Figure 6.16). Utilizing PProDOT-Hx 2 as binder is advantageous because it leads to a better conducting electrode and SEI than a non-conducting conventional binder as PVDF (see Chapter 3-5). 252 Figure 6.18 a) Cycle-life comparison and characterization of the NCA||NaNb 13O 33 full-cells at 1C with an NCA areal loading of 4 mg cm -2 and a N/P ratio of 1.3. Corresponding GCD curves from 6.18a for the b) NCA-PProDOT-Hx 2||NaNb 13O 33 and c) NCA-PVDF||NaNb 13O 33 cells. Pouch cell design. Scaling-up from a coin-cell to a pouch-cell constitutes a significant milestone for establishing the proof of concept of building a practical cell. Figure 6.19a shows a schematic representation of the pouch-cell electrodes design with their dimensions, their respective aluminum and nickel tabs, and a scheme (Figure 6.19b) of how the electrodes are placed against each other with a polypropylene separator between them. Figure 6.20 shows a photograph of the actual electrodes that were made according to design showed in Figure 6.19. 253 Figure 6.19 a) Schematics of the NCA and NaNb 13O 33 pouch cell electrodes b) representation of the electrode assembling were a polypropylene seperator is placed between the electrodes to prevent a short-circuit. The assembled pouch cell had a N/P ratio of ≈ 2 and was cycled in the previously established reversible cell potential window from 1.15 V to 3 V. The NCA-PProDOT-Hx 2||NaNb 13O 33 pouch cell constitutes the first SCALAR-EFRC 6 pouch cell and full-cell (Figure 6. 21). At C/27 the pouch cell was able to deliver a remarkable capacity of 10 mAh (Figure 6.22). Figure 6.20 Photograph of the NCA-PProDOT-Hx 2 and the NaNb 13O 33 pouch-cell electrodes coated onto Al foil with an Al tab and onto Cu foil with a Ni tab respectively. 254 Figure 6.21 Photograph of the assembled SCALAR-EFRC first NCA-PProDOT-Hx 2||NaNb 13O 33 pouch-cell. Figure 6.22 First Cycle GCD for the SCALAR-EFRC pouch-cell NCA-PProDOT-Hx 2||NaNb 13O 33 with an approximate mass loading of 4.7 mg cm -2 and an N/P ratio = 2 cycled at C/27. 255 References 1. Dubarry, M., Svoboda, V., Hwu, R. & Yann Liaw, B. Incremental Capacity Analysis and Close-to-Equilibrium OCV Measurements to Quantify Capacity Fade in Commercial Rechargeable Lithium Batteries. Electrochem. Solid- State Lett. 9, A454 (2006). 2. Bloom, I. et al. Differential voltage analyses of high-power lithium-ion cells. 4. Cells containing NMC. Journal of Power Sources 195, 877–882 (2010). 3. Fly, A. & Chen, R. Rate dependency of incremental capacity analysis (dQ/dV) as a diagnostic tool for lithium-ion batteries. Journal of Energy Storage 29, 101329 (2020). 4. Rui, X. H., Yesibolati, N., Li, S. R., Yuan, C. C. & Chen, C. H. Determination of the chemical diffusion coefficient of Li+ in intercalation-type Li3V2(PO4)3 anode material. Solid State Ionics 187, 58–63 (2011). 5. Ho, C., Raistrick, I. D. & Huggins, R. A. Application of A‐C Techniques to the Study of Lithium Diffusion in Tungsten Trioxide Thin Films. J. Electrochem. Soc. 127, 343–350 (1980). 6. Synthetic Control Across Length-scales for Advancing Rechargeables (SCALAR) A US Department of Energy Energy Frontier Research Center Class: 2018 — 2022 http://www.chem.ucla.edu/SCALAR/ (2022). 256 257 Chapter 7 - Understanding the Polarization Behavior of Ketjen black-based Lithium-Sulfur Battery Cathodes Abstract The performance of the sulfur-carbon composite cathode in a lithium-sulfur battery is chiefly determined by the sulfur content and the properties of the carbon materials. However, the role of these major constituents in achieving optimal cathode performance is often poorly understood. The present study focuses on understanding the rate limitation observed in a Ketjenblack-carbon-based sulfur cathode, with the goal of uncovering the factors contributing to the internal resistance resulting from cathode formulation and processing. Despite its enormous surface area, Ketjenblack-carbon exhibited surprisingly poor discharge rate capability, providing no more than 180 mAh g -1 at the C/20 rate and delivering a high specific capacity of 1200 mAh g -1 only at a very low discharge rate of C/50. Our studies uncovered a large contribution to the polarization resistance from inter-particulate contacts and charge- transfer processes occurred close to 30% depth of discharge. Sulfur formed thick insulating sheets on the Ketjenblack particles necessitating an additional electron-percolation pathway for reducing the polarization. By the addition of a low-surface area carbon, Super-P  , such an electron conduction pathway was provided, and the internal resistance to the Ketjenblack cathode decreased by four times and the cell was able to deliver an impressive and stable capacity of 950 mAh g -1 at C/5 rate. We anticipate the learnings from this study to provide the insight needed for improved formulations of the sulfur cathode. 258 Introduction The lithium-sulfur battery is promising as the next generation high-energy density storage battery because of its high theoretical specific energy of 2500 Wh kg -1 and volumetric energy density of 2800 Wh L -1 . These values are about five times larger than that of lithium-ion batteries. 1,2 Furthermore, the cathode active material, sulfur, is substantially more abundant and less expensive compared to materials like nickel and cobalt. 3 Yet, the commercialization of lithium-sulfur batteries has been hindered largely by poor discharge rate capability, high rates of self-discharge, low Coulombic efficiency, and limited cycle life. 4 These performance limitations arise mainly from the inherent properties of sulfur, namely: (i) electrically insulating nature, (ii) electro-reduction of sulfur to long-chain polysulfides that dissolve into the electrolyte and redistribute to various parts of the cell, and (iii) changes to the volume of the electrode to the extent of 80% during discharge of the cathode. 5 In the last two decades, various strategies have aimed at resolving the foregoing technical issues. Among the most successful of these efforts is adding high surface area carbon to the formulation of the sulfur cathode. Besides the carbon content of the cathode, the physical/ chemical nature of the carbon material and its spatial distribution have a notable effect on the performance of the sulfur cathode. 1–5 Different types of commercially-available carbon materials such as acetylene black, carbon nanotubes (CNT), graphene, Super-P  etc., have been added to the sulfur electrode mixture to the extent of 20-40 wt%. This mixture is made into a slurry or ink that is coated onto an aluminum substrate. 6 Sulfur is often impregnated into the carbon materials by melt-diffusion, sublimation or by dissolving in solvents such as carbon disulfide. Various types of heteroatom-doped carbons and hierarchical porous structures have also been proposed to improve the performance. 7,8 The surface area and pore size of the carbon significantly affected the performance of the sulfur electrode. 9 Larger pores allowed a higher loading of 259 sulfur but the reduced area of interfacial contact between sulfur and carbon lead to the lower utilization of sulfur. Thus, carbon materials with a large surface area are attractive for increasing sulfur loading, utilization, and rate capability. To this end, Ketjenblack (KB) carbon is an ideal choice owing to its enormous surface area that exceeds 1000 m 2 g -1 compared to acetylene black, Super-P  and CNT that have surface area values of less than 100 m 2 g -1 . Surprisingly, there are only a few reports available on the exclusive use of KB carbon as an additive to the sulfur electrode. 10,11 In these reports, a large fraction of the electrode mass is KB (close to 40 wt %) and the electrode preparation method is elaborate to ensure uniform dispersion of sulfur. These preparation methods involve high-energy ball milling for several hours 10 or solution infiltration of sulfur dissolved in carbon disulfide followed by melt-infusion at 155 °C. 11 KB-based electrodes also exhibited large cracks and high overpotential values during charging although these observations have not been fully understood. 10 Instead of using high-surface-area KB as such, many of the studies reported in the literature propose a composite of KB with metal oxides 12 or other carbon materials. 13,14 The results from these previous studies have prompted us to investigate the performance of lithium-sulfur batteries with KB as the sole carbon additive, and uncover the reasons for the beneficial impact of addition of other types of carbon materials. We expect that the results and analysis presented here will guide the formulation of improved carbon-sulfur composites to achieve high levels of electrical performance. The surface area, porosity and the distribution of sulfur and carbon together determine the observed electrical performance of the sulfur cathode. Firstly, since sulfur is electrically insulating, the added carbon reduces the ohmic polarization losses to raise the electrochemical utilization at high discharge and charge rates. Secondly, the large interfacial contact area provided by the porous carbon structure also lowers the polarization losses resulting from the sluggish charge-transfer process of the electro- 260 reduction of sulfur. Next, the distribution of sulfur and electrolyte within the porous carbon electrode would influence the transport of active species (soluble polysulfides and lithium ions) resulting in voltage losses attributed to a mass transport polarization resistance. 15 At low rates of discharge, we could expect charge-transfer overpotential losses to dominate, while at high rates of discharge, mass transport and ohmic losses are likely to be preponderant. The foregoing voltage losses are expected to vary with the state-of-charge of the cathode and the spatio-temporal distribution of active materials as the discharge/charge proceeds. Electrochemical impedance spectroscopy (EIS) is a versatile and non-intrusive measurement technique that permits the separation of the polarization resistance and charge storage contributions arising from various processes occurring over a wide range of timescales. 16 Previous EIS studies include, investigating the effect of interlayers to stop the polysulfide shuttle, the role of polymer coatings on the cathode materials, the effect of lithium nitrate additive in the electrolyte, and analysis of the capacity fade in solid and liquid electrolyte cells. 2,17–20 A few attempts have been made to study the variation of cell impedance with the state-of-charge. 20,21 For instance, the study by Natalia et al., 20 related capacity fade to the accumulation of non-conductive solid products. The circuit elements in the EIS fitting were related to processes occurring at the anode, cathode, and in the electrolyte. However, this study used low surface area Super-P  carbon and LiPF 6 in TEGDME as the electrolyte while most of the literature reports are based on LiTfSI in DOL-DME electrolyte. Similarly, Zhaofeng et al., 21 reported analysis of impedance variations as functions of states of charge and temperature. This study suggested that the semicircle in the high frequency was due to interphase contact whereas the semicircle in the middle frequency could be attributed to charge transfer process. In addition, the increase in charge transfer resistance was identified as a key factor in the capacity fading. The EIS technique was also used to study reaction 261 mechanisms, 22 influence of electrolyte and electrode compositions, 23 effect of electrode pressing, 24 ionic conductivity, 25 role of temperature, 26 capacity fade, 27 cathode loading effect, 28 etc. However, we noticed that all these studies use sulfur electrodes prepared with low or medium surface area carbons, and there are no studies focusing on the use of high surface area KB carbon and the resulting internal resistance of the Li-S cell. Thus, in the present study, we use EIS to understand the impact of KB carbon as a conductive additive on the performance of lithium-sulfur battery. Specifically, we study the underlying polarization phenomena that govern the rate capability of the sulfur electrode and explain the beneficial impact of adding a second kind of carbon, Super-P  on mitigating the polarization losses in a KB- containing electrode. Results and Discussion Low-rate performance of sulfur electrodes with KB-300 and KB-600 Sulfur electrodes were prepared using two different types of Ketjenblack carbon from Akzo, namely EC-300 J and EC-600 JD, designated here as KB-300 and KB-600, respectively. These two carbon materials differed markedly in surface area and pore volume. KB-600 has higher surface area (1270 m 2 g - 1 ) and larger pore volume (2.28 cm 3 g -1 ) compared to KB-300 with 800 m 2 g -1 and 1.05 cm 3 g -1 , respectively. Lithium-sulfur cells were assembled with both these materials and tested under identical conditions. The open circuit voltage was close to 2.3 V for both the cells after three hours of equilibration. Then the cells were discharged at a very low rate of C/50 (C-rate is based on the theoretical capacity of sulfur, 1.675 Ah g -1 ). The periodic voltage spikes in the charge-discharge curves (Figure 8.1) were caused by the interruptions at various states of charge for EIS measurements. For both types of cells, the initial discharge curves exhibited a short voltage plateau at 2.3 V, a sloping region before a long voltage plateau at 2.1 V followed by a sharp drop towards the end of the discharge. The voltage plateau at 2.3 V was 262 attributed to the conversion of elemental sulfur (S8) to soluble higher-order polysulfides whereas the voltage plateau at 2.1 V was governed by the precipitation of insoluble lower-order sulfides such as Li2S and Li2S2. The sloping region at the end of the discharge process is linked to a solid phase conversion of Li2S2 to Li2S. 29,30 The discharge capacity of the sulfur electrodes with KB-300 and KB-600 was close to 925 mAh g -1 and 1200 mAh g -1 , respectively (Figure 8.1). The high utilization of the active material even with 20 wt % of conductive additive was attributed to large surface area of Ketjenblack (~1000 m 2 g -1 ) compared to other commercially available carbon black materials such as acetylene black and Super-P  that yielded 800 mAh g -1 under the same conditions. The higher utilization observed with KB-600 was attributed to the higher surface area and larger pore volume compared to KB-300. 31 Thus, high surface area and large pore volume were necessary for achieving high utilization in the sulfur electrode. During charging, both the KB-300 and KB-600 electrodes showed comparable voltage profiles with no distinct regions for the conversion of various polysulfides unlike the two voltage plateaus observed in the discharge curve. Coulombic efficiency is the ratio of the discharge output divided by the charge input, expressed as a percentage. The Coulombic efficiency of less than 100% is the result of the parasitic loss via the inter-electrode shuttling of the polysulfides. The sulfur electrode prepared with KB-600 exhibited a Coulombic efficiency of 90 %, significantly higher than 82% for the KB-300 electrode. The higher Coulombic efficiency suggested a lower shuttling rate in the case of KB-600 most likely due to higher degree of adsorption of the soluble polysulfides on the high surface area carbon surface of KB-600. Since the KB-600 electrode exhibited higher initial discharge capacity and relatively high Coulombic efficiency, all further studies focused on KB-600-based sulfur electrodes. 263 Figure. 8.1 The first charge and discharge curves (C/50 charge and discharge rate) for the sulfur electrodes prepared using KB-300 and KB-600. Discharge rate capability test The sulfur electrode with KB-600 delivered an impressive specific discharge capacity of 1200 mAh g -1 at C/50 rate in the first cycle (Figure 8.1). However, when the cell was tested at a higher discharge rate of C/20, the discharge capacity was 180 mAh g -1 , just 15 % of the utilization at C/50 (Figure 8.2a(i) and (ii)). This drop in utilization upon going from C/50 to C/20 was surprising as KB has very high surface area and is expected to provide good electrical connectivity between the sulfur particles in the electrode. However, as the discharge proceeds, the soluble polysulfides in the electrolyte when electro-reduced further could precipitate on the surface of KB particles and block electrolyte access to the electrode surface and reduce the interfacial area available for reaction. Alternately, the precipitated material could become electrically isolated from the carbon particles. The precipitation is expected only at the depth of discharge of about 20-30%. This spatio-temporal redistribution of the discharged products and the saturation levels preceding precipitation will depend on the rate of discharge. Concomitant with the saturation and precipitation process we can expect the internal resistance of the cell to undergo 264 significant changes with the depth of discharge including an increase followed by a return to lower values. An increase of internal resistance would cause the voltage of the cell during discharge to dip below the cut-off voltage leading to the early termination of discharge. We suspected that this would be case with our cells. Thus, if the cut-off voltage were to be lowered, we would be able to pursue the discharge. To test this hypothesis, we lowered the voltage cut-off for C/20 discharge from the previous value of 1.2 V to 0.0 V. As anticipated, the cell voltage suffered a dip at about 30% DOD but stayed above the cut-off value and then gradually recovered to a higher value finally delivering a high discharge capacity of 1000 mAh g -1 even at C/20 rate (Figure 8.2b). The observed utilization was almost 83% of that observed at C/50 (Figure 8.2a(i)) and almost seven times higher than the capacity that observed when the voltage cut-off was set at 1.2 V (Figure 8.2a(ii)). Furthermore, the large voltage polarization causing the drop of the cell voltage below 1.2 V was seen between 170 to 320 mAh g -1 (15-30 % DOD) consistent with the depth-of-discharge values at which precipitation of the disulfide is expected. Following the forced discharge, 85 % of the capacity was still above the 1.2 V. Thus, our finding was that the rate capability in a lithium-sulfur cell with KB is drastically reduced by the sharp rise of the polarization resistance at ~30 % DOD. However, with persistent forced discharge at the same rate, the subsequent decrease of the polarization losses allowed for a more complete utilization of the sulfur electrode. 265 Figure. 8.2 (a) The first discharge curves with 1.2 V voltage cut-off at (i) C/50 rate and (ii) C/20 rate, and (b) the discharge curve at C/20 rate with 0 V voltage cut-off. EIS studies on sulfur electrode prepared using KB-600 Nyquist plots at different depths of discharge (DOD) during the first discharge showed that the total impedance of the cell increased with DOD reaching a maximum value around 30 % DOD and then decreased (Figure 8.3a). This trend in impedance values was consistent with the increased polarization observed in the discharge curve at ~30 % DOD (Figure 8.2b). The Nyquist plots below and above 33% DOD had characteristic differences (Figure 8.3b and 8.3c). Below 33 % DOD, two partially overlapping and depressed semicircles were observed in the range of 100 kHz to 0.2 kHz and 200 Hz-1 Hz (Figure 8.3b). The beginning of a sloping line was noted around 1 Hz but this feature did not contribute much to the overall impedance. An intercept of comparatively small magnitude of 5-10 Ω was observed above100 kHz. Thus, below 33 % DOD the Nyquist plots could be modeled by two sets of parallel R and C elements connected serially along with another resistance (Figure 8.3b inset). In practice, a constant phase element (CPE) substituted an ideal capacitor (C) in consideration of the distributed nature of the capacitive elements of the porous electrode. The impedance of a generalized CPE is defined as, 32 266 Z CPE = 1/[(Q 0(jω) n ] (1) where ω is the angular frequency and n is a constant number between -1 and 1 and j = (-1) 1/2 . Thus, when n=1 the element becomes identical to a capacitor, when n=0 it becomes a resistor, and when n= -1 it turns into an inductor. In our case, the n value was close to 0.8 for CPE1 and CPE2 (Table 1). At DOD values greater than 33%, only one semi-circle with a sloping line at low frequencies was observed (Figure 8.3c). Thus, for the Nyquist plots above 35 % DOD, we used an equivalent circuit as shown in Figure 8.3c inset. For the inclined line in the low frequency region, we chose to model by a CPE as in a previous literature report. 29 The attempt to model the line with a Warburg element or by finite diffusion elements produced large errors during fitting. In fact, CPE represents the Warburg element when n= 0.5 according to Eq.1. In our case, the n value for CPE3 was close to 0.3 (Table 1) which was much lesser than 0.8 observed for CPE1 and CPE2, suggesting the involvement of diffusional mass transport processes. By using the equivalent circuits shown in the inset of Figure 8.3b and c, the fitted impedance plots agreed well with the actual impedance spectra, within 5% error (Table 8.1). The physical significance of each resistance element in the context of lithium-sulfur battery is described below in detail. Table 7.1. Representative impedance fit parameters for 20 % and 72 % DOD 267 Figure. 8.3 EIS analysis during first discharge: Here (a) Nyquist plots at different DODs, and representative Nyquist plots for the cells discharged (b) below and (c) above 35 % DOD. Inset shows corresponding equivalent circuits. Physical significance of resistance elements (R0, R1 and R2) In the equivalent circuits (inset in Figure 8.3b and 8.3c), the series resistance, R0 represents pure ohmic resistance which is the sum of the contribution from the electrolyte resistance, the resistance of the current collectors, cell casing and electrical connections. However, the electrolyte dominates the value of R0 in lithium-sulfur batteries due to the poor ionic conductivity of organic electrolytes compared to electronic conductivity of carbon, the stainless-steel current collector, and cell casing. During the discharge process, higher-order polysulfides dissolve in the electrolyte to increase the viscosity reducing the mobility of ions, leading to an increase in electrolyte resistance. 268 There has been controversy regarding the assignment of the semicircle seen at high frequencies (HFS), typically above 200 Hz. 20, 29, 34-37 This semicircle is seen at all states of charge. In the lithium-sulfur battery, the insoluble discharge products are generated only at the lower voltage plateau and the electrolyte is electrochemically stable on the surface of cathode. Hence, this semicircle cannot be ascribed to the formation of insoluble discharge products or a solid-electrolyte interface (SEI) on the cathode surface. 29 Canas et al., attributed HFS to the charge- transfer resistance at the anode surface since this semicircle appears at all states of charge and is considerably smaller than the low frequency semicircle (LFS) arising from the fast kinetics of lithium stripping and plating compared to sulfur electrode reaction. 20 However, it is to be noted lithium metal on exposure to organic electrolyte quickly reacts and forms a layer of complex composition called the solid electrolyte interphase (SEI). Such an SEI usually has an inner layer formed by species such as lithium hydroxide (LiOH), lithium fluoride (LiF) and lithium nitride(Li 3N) and an outer organic layer composed of alkoxide (ROLi) like moieties. 33 The dissolved polysulfide ion in the electrolyte can diffuse to the anode and react with lithium to further modify the composition and thickness of this SEI layer during cycling. Hence, considerable variation in the value of R1 during cycling is to be expected. Similarly, Qiu et al., 29 reported that R1 originates from the Schottky junction formed at the interface of conducting carbon and insulating sulfur. According to thermionic emission diffusion theory, the Schottky contact resistance is described as, 29 ln R1 = ln C - qE/nkT (2) where C = (I 0q/nkT) -1 ; q is the electronic charge, I 0 is the saturation current, E is the applied voltage, n is the ideality factor, T is the absolute temperature and k is the Boltzmann constant. When the contact media do not change, C is constant and thus a plot of ln R1 versus E will be a straight line with q/nkT as the slope and ln C as the y-intercept. In the case of lithium-sulfur battery, contact media can change 269 owing to dissolution of polysulfide at higher voltage plateau region followed by re-precipitation in the lower plateau. Thus, several straight lines with different slopes are expected at different voltages. 29 Similarly, in the lithium-ion battery literature, HFS was also interpreted as an interparticle contact resistance in the electrode bulk. 34 Here, the dissolution of insulating sulfur will improve interphase contact between conducting particles, and the production of insulating discharge products could increase the contact resistance and enlarge the HFS. Thus, as described in the literature, charge transfer at the anode SEI, the Schottky barrier resistance, and interparticle contact resistance could contribute to the net value of R1 since these processes with similar characteristic frequencies are plausible in the lithium-sulfur battery. 34-37 The low-frequency semicircle (LFS) is associated with the electrochemical reaction at the sulfur cathode. This assignment is supported by previous studies on the temperature dependence of the diameter of this semicircle. 21 Here, the resistance element represents the charge-transfer resistance at the cathode (ie. R2=R ct,cathode) and the CPE element represents the capacitive properties of the double layer (CPE2~C dl). The charge-transfer resistance is inversely related to the rate constant of the reaction, electrochemically active surface area, and concentration of reactants near the reaction site. In the case of lithium-sulfur battery, multiple elementary reactions with different rate constant values occur as the discharge proceeds. 35 Also, dissolution of active material and subsequent shuttling produce variation in the surface area and concentration of polysulfides which will cause notable variation in the charge transfer resistance (R2). Variation in R0, R1 and R2 during 1 st discharge In the previous section we have considered the plausible processes occurring during the operation of lithium-sulfur battery that can be represented by the equivalent circuit elements in Figure 8.3b and 8.3c. 270 A plot of open circuit voltage vs. DOD is annotated with descriptions of the various processes occurring at different stages of discharge (Figure 8.4a) to support the explanation of the variation of the various impedance parameters. 271 Figure. 8.4 (a) Open circuit voltage variation with DOD, (b) series resistance (R0) changes with DOD, (c) resistance from HFS (R1) variation with DOD, (d) a plot of lnR1 vs. OCV, (e) effect of lithium polysulfides on impedance of Li/Li symmetric cells, and (f) variation of charge transfer resistance (R2) with DOD. 272 The series resistance, R0 was close to 6 Ω at the beginning of discharge and increased as the cell was discharged and attained a maximum value of 10 Ω at around 30 % DOD (Figure 8.4b). After this point the value of R0 decreased gradually returning to the initial value at the end of discharge. Conversion of S 8 to soluble Li 2S 4 delivers 25 % of the theoretical capacity. Hence, the initial increase in R0 until 30 % DOD can be attributed to the increase in the viscosity of electrolyte caused by the dissolution of the higher- order polysulfides between 2.6 and 2.1 V (Region-I, Figure 8.4a). Beyond 30 % DOD, the soluble discharge products are converted in to insoluble Li 2S 2 and Li 2S over Region II. These species begin to precipitate as their solubility is not high, reducing the viscosity again leading to higher electrolyte conductivity and lower series resistance over Region-II of Figure 8.4a. 36 Interestingly, formation of insulating solid products of Li 2S 2 or Li 2S at the end-of-discharge process did not increase the R0 value suggesting a more uniform distribution of these insulating products on the surface of carbon particles (Region III of Figure 8.4a) The variation in the resistance from HFS (R1) followed almost the same trend as that of R0 except that the absolute values of R1 at any DOD were almost 40 times higher than R0 (Figure 8.4c). The initial value of R1 that was around 200 Ω gradually increased to reach a maximum value of 400 Ω at 30 % DOD before the value started to drop. There was again a small increase in the value towards the end of DOD. To test the validity of resistance contribution from Schotkky junction, a plot of lnR1 against the cell voltage was analyzed (Figure 8.4d) which should produce a straight line. As shown in Figure 8.4d, four straight lines with different slopes were found in the different potential regions. The straight lines with various slopes were attributed to the differences in the surrounding medium that caused changes in the saturation current (in Eq. 2). 29 This analysis of R1 is in support of the process of charge-transfer through a Schottky junction between carbon and insulating sulfur. However, slight deviation from the straight lines in the 273 Figure 8.4d suggested that the SEI on lithium and interparticle resistance could also contribute to the observed value of R1. At 0 % DOD, there would be no soluble polysulfide and hence the initial contribution to the resistance from anode is by the pristine SEI formed by the reaction of lithium with the ether solvents and LiTfSI salts. However, upon discharge, soluble polysulfides are formed which diffuse to the anode and react with lithium to modify initial SEI. If such reactions had produced a thick insulating layer on the surface of lithium, resistance, R1 wouldn’t have dropped back to the initial value after 30 % DOD (Figure 8.4c). In addition, studies on Li-Li symmetric cells ( Figure 8.4e) further confirmed that the lithium anode in the presence of polysulfide contributed less than 35 Ω or only 8 % of the maximum value of R1. Similarly, as shown in Figure 8.4c, there was a small increase in R1 especially after 80 % DOD. This increase was in the region-III in Figure 8.4a where solid Li 2S is formed by the reduction of Li 2S 2. The formation of electronically and ionically insulating Li 2S accompanied by volume change could reduce the inter-particle contact in the bulk leading to a higher resistance value. Thus, our analysis confirmed that the HFS could be attributed to multiple processes such formation of anodic SEI, a Schottky junction and inter-particle contacts. The contribution of R1 to the total cell impedance was significantly larger than R0. The value of R2 was as high as 400 Ω even at 0 % DOD (Figure 8.4f). The contribution of R2 was 80 times larger than R0, and twice the value of R1 even in the beginning of discharge suggesting that the overall internal resistance of the cell was dominated by the charge transfer process. Since sulfur is an insulator, the electrons are transported through the conductive network of KB to the active material followed by electron transfer at the reaction site. Thus, the charge transfer resistance depends on the proximity of the Ketjenblack particles and the sulfur particles. The high value of R2 even at 0 % DOD suggested poor distribution of such a conductive network around sulfur (vide infra). The variation in R2 with DOD 274 followed a similar trend as R0 and R1 initially (Figure 8.4f), but its values dropped significantly after 30 % DOD marked by the complete disappearance of LFS and development of an inclined line in the Nyquist plots. The initial increase in R2 up to 20 % DOD was attributed to depletion in the reactant concentration due to diffusion of soluble polysulfides away from cathode, and reduction in the rate constants (or exchange current) as short chain polysulfides 35 were formed consistent with Eqs. 3 and 4, R ct (or R2) = RT/nFj o (3) j o = nFk s(C o) α (C R) β (4) where R is gas constant, T absolute temperature, j 0 is the exchange current density, k s is the standard rate constant, C o and C R are concentration of oxidized and reduced species, with α and β as the cathodic and anodic transfer coefficients. At the end of the higher voltage plateau region (Region-I, Figure 8.4a), most of the active material is dissolved in the electrolyte providing an open network of conductive carbon. The availability of such a large surface area was consistent with the large drop in the charge-transfer resistance value in the lower plateau region (Region-II, Figure 8.4a). However, it is to be noted that actual magnitude of R2 might depend on pore length and porosity as determined by a porous electrode model 37 which is beyond the scope of the present study. Nonetheless, the qualitative assessment of R2 value suggested that the charge-transfer resistance dominated mainly in the upper voltage plateau region (Region-I, Figure 8.4a) where the soluble polysulfides were formed, and the mass transfer effects were more pronounced in the lower plateau region as well as in the sloping region at the end of the discharge (Region-II and III, Figure 8.4a) when insoluble discharge products obstructed ion movement. Then again, a sum of all three resistances R0, R1 and R2 reached as high as 600 Ω near 30 % DOD, resulting in the huge polarization and limited discharge rate capability (Figure 275 8.2). Among these three resistors, the charge transfer resistance (R2) contributed significantly followed by R1 and R0. Thus, we emphasize that rate capability in lithium-sulfur cells with KB carbon was limited by the large polarization (Figure 8.2b) caused by the high charge-transfer resistance and contact resistance at the cathode (Figure 8.4c and f). Hence, we were convinced that the cathode structure must be modified to ensure uniform dispersion of sulfur to provide more reaction sites. This modification would lead to the reduction of both the contact resistance and charge transfer resistance values, and hence an increase in the rate capability of the KB-based sulfur electrode. Effect of melt infusion of sulfur on the internal resistance Since the charge-transfer resistance (R2) was identified as limiting the rate capability, we decided to modify the cathode structure by commonly-used techniques. The charge transfer in the sulfur electrode occurs at the triple phase boundary of sulfur, carbon, and the electrolyte, and hence R2 is expected to depend directly on the distribution of sulfur on the carbon surface. The melt infusion process is a commonly-used strategy to ensure uniform distribution of sulfur. 1,2 Hence, we decided to employ this technique for incorporating sulfur into KB. This was achieved by melting sulfur with KB at 155 ᵒC for 12 h as reported with other carbons in the literature. 7 α-sulfur melts at 115 ᵒC and displays a distinct viscosity-temperature behavior in the molten state. On heating, the viscosity of the sulfur slowly decreases, followed by a notable increase beyond 160 ᵒC as a result polymerization. At 190 ᵒC the sulfur starts depolymerizing and the viscosity again decreases. 38 Therefore, the viscosity of molten sulfur being the lowest between 135 o C and 160 ᵒC is a suitable temperature range for impregnating sulfur into the KB matrix. The electrodes were prepared and tested under identical conditions to compare the “melt-infused” electrode with the electrode prepared by simple mixing, referred to here as “mixed-in” electrodes. The 276 melt-infused sulfur electrode yielded 1200 mAh g -1 in the first discharge at C/50 rate comparable to the value obtained for the regular electrode (Figure 8.5a). The Nyquist plots showed similar trends for both melt-infused and mixed-in electrodes (Figure 8.5b). The overall cell impedance of the cell was not affected significantly, and again there was a large increase in impedance near 30 % DOD. The Nyquist plots could be fit with the same equivalent circuits as in Figure 8.3 to extract R0, R1 and R2 values. The variation of these parameters with DOD are shown in Figures 8.5c to 5e. The series resistance, R0 and resistance from the HFS, R1 were not affected. R0 and R1 in the melt-infused electrodes exhibited similar trends to the mixed-in electrodes, increasing initially and then reaching their maximum value at 30 % and then decreasing again. The charge-transfer resistance, R2 was slightly reduced in the melt-infused electrodes. For instance, the maximum value R2 was 600 Ω for the mixed-in sulfur electrode whereas it was close to 450 Ω for the melt-infused electrode. The slight decrease in the value of R2 suggested better distribution of sulfur active material on the surface of the KB. Although the melt infusion process reduced the charge transfer resistance by 150 Ω at the peak, the absolute values were still significantly higher than the reported values on other carbons. 29 This high value of R2 suggested that melt infusion may not be sufficient to efficiently distribute sulfur on the surface of the KB particle. TEM images and elemental mapping of sulfur on the surface of KB (Figure 8.6) supported this hypothesis. The TEM image showed roughly spherical particles of KB surrounded by sulfur on the outer surface. Only a small fraction of the sulfur was noticed in the interior of the particles although KB had a void fraction of 80 %. Thus, most of this inner particle surface area was unutilized. Since sulfur is insulating, electron transfer cannot occur at the interface of a thick outer sulfur layer and electrolyte. This observation was consistent with the large charge transfer resistance observed for the sulfur electrode. 277 Figure. 8.5 (a). First discharge curves at C/50 for the sulfur electrodes prepared by (i) mixing and (ii) melt infusion processes and (b) Nyquist plots for the lithium-sulfur cell with sulfur electrode prepared using melt infusion technique. The variation in the (c) R0, (d) R1 and (e) R2 with DOD 278 Figure. 8.6. (a) TEM image of KB particle infused with sulfur, and element mapping of (b) carbon and (c) sulfur. Effect of mixing with Super-P  The TEM images confirmed that sulfur formed a thick insulating layer on the surface of the KB particles even after the melt-infusion process. Thus, it was necessary to provide an electron-conduction pathway to this layer for the electrochemical utilization of the active materials. We proposed that such an improvement could be achieved by adding another low-surface area carbon such as Super-P  on the top of the melted-infused sulfur-KB composite. For the preparation of this “hybrid” electrode, sulfur was first melt-infused into the KB particles at 155 ᵒC for 12 h in air. Separately, the PVDF binder dispersed in NMP was mixed with Super-P  followed by the addition of the melt-infused sulfur-KB. Then the 0 . 1 µ m 0 . 1 µ m 0. 1 µm [a] [b] [c] 279 electrodes were processed as before and assembled with a lithium anode into coin cells. The performance of the cells with the hybrid electrodes were compared with cells containing mixed-in and simple melt-infused electrodes. Although, these kinds of protocol for electrode fabrication have been routinely used in the literature, a systematic study on the role of the added carbon and its impact on internal resistance of the cell have not been presented and analyzed. 39,40 The hybrid electrodes provided 1200 mAhg -1 in the first discharge at C/50 rate comparable to the value observed for the melt-infused electrodes (Figure 8.7a). However, the voltage curves suggested a lesser degree of polarization especially in the upper plateau region( inset of Figure 8.7a). For instance, the discharge voltage was close to 2.04 V at 400 mAhg -1 for the electrode prepared without adding Super- P  whereas it was 2.1 V for the hybrid electrode. Thus, there was 60 mV difference between the discharge voltages of these two electrodes even at a discharge rate as low as C/50 suggesting a significantly lower internal resistance for the hybrid electrode. This difference was confirmed by EIS measurements at different values of DOD. The total impedance for the hybrid electrode showed almost a four-fold reduction at all DODs ( Figure 8.7b) with a maximum value of 190 Ω compared to ~ 800 Ω for both the mixed-in (Figure 8.3a) and the simple melt-infused electrodes (Figure 8.5b). The series resistance, R0 for the hybrid electrode followed the same trend as before with a maximum value 10 Ω near 30 % DOD which was only slightly lower than the values obtained for mixed-in and melt-infused electrodes. Since R0 value was mainly determined by the electrolyte resistance, we did not expect the cathode modification to alter this value. However, the electrolyte resistance R0 was affected by the viscosity changes due to the dissolution of the higher order polysulfides passing through a maximum value just as in the case of the mixed-in and melt-infused electrodes. There was a dramatic decrease in the resistance value of HFS, R1 after incorporating Super-P  (Figure 8.7d). The resistance was as low as 280 50 Ω at 0 % DOD, reached maximum of 100 Ω near 30 % DOD, and then dropped to the original value of 50 Ω. This was in stark contrast to 400 Ω observed for the electrodes prepared without Super-P  . Such a sharp decrease in the value of R1 by a simple cathode modification further confirmed that the inter- particle contact resistance might be determining the absolute value of R1, and that the contribution of the SEI on the lithium electrode to the value R1 was far less significant. It is also to be expected that the calendaring process that is commonly used in the industrial battery electrode fabrication can improve the contact and reduce the value of R1. Similar to R1, the charge-transfer resistance (R2) was also reduced significantly in the case of the hybrid electrode (Figure 8.7e). The maximum value of R2 was 125 Ω which was four and five times smaller than that of the mixed-in and simple melt-infused sulfur electrodes, respectively. Thus, the added Super-P  were able to provide an electron percolation pathway to the insulating layer of sulfur on the surface of the KB particles. With such an electron transport pathway, the sulfur surrounded with Super-P  can participate in the battery reaction even when discharged at a high rate. Thus, the dramatic reduction in the values of R1 and R2 after adding Super-P  would improve the discharge rate capability of the sulfur electrode. 281 Figure. 8.7 (a) First discharge curves. Inset shows the enlarged region near 30 % DOD. (b) Nyquist plots at different DOD for a cell with hybrid sulfur electrode. The variation in (c) R0, (d) R1 and (e) R2 with the depth of discharge. 282 Rate capability and Long-term cycling test The combined approach of melt infusion of sulfur to KB and the addition of another layer of Super-P  carbon significantly reduced the charge transfer resistance and inter-particle contact resistance. Hence, we expected the hybrid electrode to yield a higher capacity at high discharge current compared to the mixed-in and melt-infused sulfur electrodes. The cell with the hybrid electrode delivered a discharge capacity of 1200 mAh g -1 at C/50 rate with 99 % coulombic efficiency (Figure 8.8a). It was also able to deliver reasonable capacities of 1000 and 950 mAh g -1 at C/20 and C/5, respectively. This is stark contrast to less than 200 mAh g -1 observed for the regular sulfur electrode prepared by the mixing process (Figure 8.2) due to large polarization. The cell with the hybrid electrode could also be charged and discharged multiple times at C/5 rate with a stable capacity close to 750 mAh g -1 for 50 cycles (Figure 8.8b). Figure. 8.8 (a) Rate capability test for the lithium-sulfur cell with hybrid electrode and (b) long-term cycling at C/5 charge and C/5 discharge rates. It is to be noted that such high-rate performance and remarkable stability of discharge capacity for several hundreds of cycles has been reported in the literature for the lithium-sulfur cell using specialized cathode structures. 1-5,40 Also, in these studies, the electrolyte usually contains special additives and/or the separator has several types of blocking interlayers to stop the polysulfide shuttle. However, the 283 present study has provided the fundamental understanding of rate capability limitations in the lithium- sulfur cell with sulfur electrodes prepared with KB carbon and shown how the internal resistance and sulfur utilization is modified by the addition of Super-P  . EIS was particularly useful in gaining insight into the factors affecting the high current performance in such cells, and show how simple modifications to the cathode fabrication processes improved the interfacial resistance to charge transfer. Conclusion In this work, we systematically evaluated the polarization characteristics of lithium-sulfur cells fabricated with high surface area KB-based sulfur cathodes. Electrochemical impedance spectroscopy (EIS) technique was used to gain insight into the variations in internal resistance of the cell, and to understand the origin of the polarization losses during the discharge process. These results allowed us to understand the benefits of the melt-infusion approach vis-à-vis the addition of lower surface area carbon towards the lowering of cell impedance. The electrode prepared with KB EC 600 JD by regular mixing process provided a reasonable initial discharge capacity of 1200 mAh g -1 at C/50 discharge rate but yielded less than 200 mAh g -1 at C/20 rate, signifying poor discharge rate capability. Large polarization losses were noted with the melt-infused electrodes at the end of soluble region. Thus, the electrode utilization was low during high-current discharge. We have identified the charge-transfer resistance and inter-particle contact resistance as primarily responsible for this polarization. TEM studies revealed that sulfur formed a thick insulating layer on the outer surface of the carbon particles during electrode preparation even after employing the melt-infusion strategy. This internal structure of the melt-infused electrode demanded additional electron-conduction pathways to access the sulfur layer. Such a pathway was provided by a network of low-surface area Super-P  carbon. As a result, the internal 284 resistance was lowered by four times and the cell was able to deliver a remarkable discharge capacity of 950 mAh g -1 at C/5 rate and a stable capacity value of 750 mAh g -1 even after 50 cycles. Thus, this study identified the fundamental limitations to the discharge of the Ketjenblack-based sulfur cathode and showed how a simple modification of the cathode architecture by the addition of low surface area carbon dramatically reduced the internal resistance. We expect that the study presented a basis for the rational design for improving the cathode structure for high performance lithium-sulfur batteries. References 1. Manthiram, A., Fu, Y., Chung, S.-H., Zu, C. & Su, Y.-S. Rechargeable Lithium–Sulfur Batteries. Chem. Rev. 114, 11751–11787 (2014). 2. Elizalde-Segovia, R., Irshad, A., Zayat, B. & Narayanan, S. R. Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature. J. Electrochem. Soc. 167, 140529 (2020). 3. Sevilla, M., Carro-Rodríguez, J., Díez, N. & Fuertes, A. B. Straightforward synthesis of Sulfur/N,S-codoped carbon cathodes for Lithium-Sulfur batteries. Sci. Rep. 10, 4866 (2020). 4. Zhu, K. et al. How Far Away Are Lithium-Sulfur Batteries From Commercialization? Front. Energy Res. 7, 123 (2019). 5. Li, S., Jin, B., Zhai, X., Li, H. & Jiang, Q. Review of Carbon Materials for Lithium-Sulfur Batteries. ChemistrySelect 3, 2245–2260 (2018). 6. Eftekhari, A. & Kim, D.-W. Cathode materials for lithium–sulfur batteries: a practical perspective. J. Mater. Chem. A 5, 17734–17776 (2017). 7. Li, Q., Guo, J., Zhao, J., Wang, C. & Yan, F. Porous nitrogen-doped carbon nanofibers assembled with nickel nanoparticles for lithium–sulfur batteries. Nanoscale 11, 647–655 (2019). 8. Chulliyote, R., Hareendrakrishnakumar, H. & Joseph, M. G. Hierarchical Porous Carbon Material with Multifunctionalities Derived from Honeycomb as a Sulfur Host and Laminate on the Cathode for High-Performance Lithium–Sulfur Batteries. ACS Sustain. Chem. Eng. 7, 19344–19355 (2019). 9. Kang, N. et al. Cathode porosity is a missing key parameter to optimize lithium-sulfur battery energy density. Nat. Commun. 10, 4597 (2019). 10. Jozwiuk, A., Sommer, H., Janek, J. & Brezesinski, T. Fair performance comparison of different carbon blacks in lithium–sulfur batteries with practical mass loadings – Simple design competes with complex cathode architecture. J. Power Sources 296, 454–461 (2015). 11. Zheng, J. et al. Revisit Carbon/Sulfur Composite for Li-S Batteries. J. Electrochem. Soc. 160, A1624–A1628 (2013). 285 12. Qian, X. et al. CeO 2 nanodots decorated ketjen black for high performance lithium–sulfur batteries. RSC Adv. 6, 111190–111196 (2016). 13. Zhao, D. et al. Separator modified by Ketjen black for enhanced electrochemical performance of lithium– sulfur batteries. RSC Adv. 6, 13680–13685 (2016). 14. Yu, X. et al. Direct Observation of the Redistribution of Sulfur and Polysufides in Li-S Batteries During the First Cycle by In Situ X-Ray Fluorescence Microscopy. Adv. Energy Mater. 5, 1500072 (2015). 15. Zhang, K. et al. A novel class of functional additives for cyclability enhancement of the sulfur cathode in lithium sulfur batteries. Inorg. Chem. Front. 5, 2013–2017 (2018). 16. Single, F., Horstmann, B. & Latz, A. Theory of Impedance Spectroscopy for Lithium Batteries. J. Phys. Chem. C 123, 27327–27343 (2019). 17. Moy, D. & Narayanan, S. R. Mixed Conduction Membranes Suppress the Polysulfide Shuttle in Lithium- Sulfur Batteries. J. Electrochem. Soc. 164, A560–A566 (2017). 18. Fu, Y. & Manthiram, A. Orthorhombic Bipyramidal Sulfur Coated with Polypyrrole Nanolayers As a Cathode Material for Lithium–Sulfur Batteries. J. Phys. Chem. C 116, 8910–8915 (2012). 19. Liang, X. et al. Improved cycling performances of lithium sulfur batteries with LiNO3-modified electrolyte. J. Power Sources 196, 9839–9843 (2011). 20. Cañas, N. A. et al. Investigations of lithium–sulfur batteries using electrochemical impedance spectroscopy. Electrochimica Acta 97, 42–51 (2013). 21. Deng, Z. et al. Electrochemical Impedance Spectroscopy Study of a Lithium/Sulfur Battery: Modeling and Analysis of Capacity Fading. J. Electrochem. Soc. 160, A553–A558 (2013). 22. Waluś, S., Barchasz, C., Bouchet, R. & Alloin, F. Electrochemical impedance spectroscopy study of lithium– sulfur batteries: Useful technique to reveal the Li/S electrochemical mechanism. Electrochimica Acta 359, 136944 (2020). 23. Conder, J. et al. Electrochemical impedance spectroscopy of a Li–S battery: Part 1. Influence of the electrode and electrolyte compositions on the impedance of symmetric cells. Electrochimica Acta 244, 61–68 (2017). 24. Barchasz, C., Leprêtre, J.-C., Alloin, F. & Patoux, S. New insights into the limiting parameters of the Li/S rechargeable cell. J. Power Sources 199, 322–330 (2012). 25. Xue, W. et al. Manipulating Sulfur Mobility Enables Advanced Li-S Batteries. Matter 1, 1047–1060 (2019). 26. Stroe, D. I. et al. An Electrochemical Impedance Spectroscopy Study on a Lithium Sulfur Pouch Cell. ECS Trans. 72, 13–22 (2016). 27. Risse, S. et al. Correlation of capacity fading processes and electrochemical impedance spectra in lithium/sulfur cells. J. Power Sources 323, 107–114 (2016). 28. Sun, K., Liu, H. & Gan, H. Cathode Loading Effect on Sulfur Utilization in Lithium–Sulfur Battery. J. Electrochem. Energy Convers. Storage 13, 021002 (2016). 29. Qiu, X., Hua, Q., Zheng, L. & Dai, Z. Study of the discharge/charge process of lithium–sulfur batteries by electrochemical impedance spectroscopy. RSC Adv. 10, 5283–5293 (2020). 286 30. Moy, D., Manivannan, A. & Narayanan, S. R. Direct Measurement of Polysulfide Shuttle Current: A Window into Understanding the Performance of Lithium-Sulfur Cells. J. Electrochem. Soc. 162, A1–A7 (2015). 31. Choi, J.-Y., Hsu, R. S. & Chen, Z. Highly Active Porous Carbon-Supported Nonprecious Metal−N Electrocatalyst for Oxygen Reduction Reaction in PEM Fuel Cells. J. Phys. Chem. C 114, 8048–8053 (2010). 32. Córdoba-Torres, P. et al. On the intrinsic coupling between constant-phase element parameters α and Q in electrochemical impedance spectroscopy. Electrochimica Acta 72, 172–178 (2012). 33. Yu, X. & Manthiram, A. Electrode–electrolyte interfaces in lithium-based batteries. Energy Environ. Sci. 11, 527–543 (2018). 34. Ju, H., Wu, J. & Xu, Y. Revisiting the electrochemical impedance behaviour of the LiFePO4/C cathode. J. Chem. Sci. 125, 687–693 (2013). 35. Wang, D. R. et al. Rate Constants of Electrochemical Reactions in a Lithium-Sulfur Cell Determined by Operando X-ray Absorption Spectroscopy. J. Electrochem. Soc. 165, A3487–A3495 (2018). 36. Yuan, L., Qiu, X., Chen, L. & Zhu, W. New insight into the discharge process of sulfur cathode by electrochemical impedance spectroscopy. J. Power Sources 189, 127–132 (2009). 37. Lasia, A. Impedance of Porous Electrodes. ECS Trans. 13, 1–18 (2019). 38. Ferreira, A. G. M. & Lobo, L. Q. The low-pressure phase diagram of sulfur. J. Chem. Thermodyn. 43, 95– 104 (2011). 39. School of Metallurgy and Environment, Central South University, Changsha, Hunan 410083, China & Li, T. Effects of Carbon Hosts on Electrochemical Properties of Lithium-Sulfur Batteries. Int. J. Electrochem. Sci. 5731– 5741 (2017) doi:10.20964/2017.06.41. 40. Ji, X., Lee, K. T. & Nazar, L. F. A highly ordered nanostructured carbon–sulphur cathode for lithium–sulphur batteries. Nat. Mater. 8, 500–506 (2009). 287 Chapter 8 - Solid-State Lithium-Sulfur Battery Based on Composite Electrode and Bi-layer Solid Electrolyte Operable at Room Temperature Abstract We report a unique solid-state lithium-sulfur cell based on a bilayer electrolyte and composite solid- state cathode. The bilayer electrolyte that contains a layer of mixed conduction membrane and a layer of polymer electrolyte eliminates the use of organic flammable liquid electrolytes and separators. The sulfur electrode is also a unique composite of sulfur with ionically-conducting intercalating nano- particulate material. Unlike many other solid-state batteries, this cell can be cycled at room temperature to utilize 85% of the active material at the sulfur electrode. The low internal resistance of the cell is comparable to that of a liquid electrolyte based lithium-sulfur cell. Impedance studies indicate that the low internal resistance results from the high ionic conductivity of the intercalating nano-particulate materials and the thin layer of polymeric electrolyte. While the volume changes at the cathode result in loss of inter-particle contact with repeated cycling, the addition of alumina to the polymer layer improved the capacity retention. This unique solid-state cell configuration opens a new pathway towards a safer high-energy lithium battery. Introduction Lithium-Sulfur (Li-S) batteries have the potential to be the next-generation candidate energy storage systems to replace lithium-ion batteries due to the high theoretical specific capacity of the sulfur electrode (1672 mAh g -1 ), high theoretical specific energy of the cell (2600 Wh Kg -1 ), and the relatively low cost of the active materials. 1–6 Nevertheless, the widespread practical application of Li-S batteries has been deterred by technical issues that limit the realization of battery cycle life, specific energy, 288 energy density, and rate capability. These performance limitations arise from the low conductivity of the active materials, volume changes in the sulfur electrode, dendrite formation in the lithium electrode, flammability of the organic electrolyte, and the effects of internal shuttling of the soluble polysulfides. 7,8 Organic liquid electrolytes typically used in a Li-S battery such as 1,3-dioxolane (DOL), 1,2- dimethoxymethane (DME), and tetrahydrofuran (THF) are flammable and potentially hazardous especially when the battery is operated at high temperatures, overcharged, or when any mechanical damage occurs. 9 In addition, the use of organic liquid electrolytes produces the well-known shuttling of the polysulfides involving the dissolution, re-distribution, and precipitation of the polysulfide active materials during cycling, affecting the reversibility and efficiency of the battery. 10–12 The use of solid-state electrolytes (SSEs) is an attractive approach to avoid flammable liquid electrolytes, suppress the shuttle effect, and inhibit the growth of lithium dendrites. 13–16 Accordingly, a variety of approaches have been reported utilizing SSEs for Li-S batteries. 17–23 SSEs can be either inorganic or polymeric materials. Nevertheless, some of the main issues with inorganic SSEs are their poor mechanical properties such as brittleness, poor interfacial contact between the solid electrode and the solid electrolyte, instability in air, and need for elevated temperature to achieve acceptable values of ionic condictivity. 24,25 On the other hand, polymeric SSE are generally stable in air and their inherent flexibility can lower the solid/solid interphase resistance, and accommodate the volume changes that occur in the cell. 26 Nevertheless, polymer SSEs generally lack mechanical strength to stop the dendrite growth and require the use of elevated temperatures to reach the desirable values of ionic conductivity. Further, the utilization of SSEs requires the incorporation of the ionically-conducting materials in the sulfur electrode. Adjusting the electrode structure and composition of the sulfur electrode is a challenge considering the insulating nature of sulfur and the volume expansion that occurs during discharge. These 289 limitations further impose a compromise between achieving high energy density and rate capability. 27– 30 Given the advantages and limitations of both types of SSEs, a hybrid composite of polymer and inorganic materials presents itself as a feasible strategy to overcome the problems encountered in the Li-S battery. 31 There are many reports related to the use of SSEs bilayers as bifunctional separators for improving Li-S performance. 32–39 For example, Zhu et al. reported an SSE bilayer composed by sulfur- carbon nanofibers and a Li 0.33La 0.557TiO 3 (LLTO)-poly(ethylene oxide) (PEO) composite electrolyte showing up to 400 mAh g -1 at C/20. 35 Also, Fu et al. reported a garnet bilayer SSE with a high sulfur loading of 7 mg cm -2 reaching up to 600 mAh g -1 . 36 In the research presented here, we demonstrate for a new mixed-conductor composite solid-state electrode/bilayer solid-state electrolyte configuration (CEBE) that aims at overcoming the challenges described above to realize a solid-state Li-S battery that can be operated at room temperature. The concept of a bilayer solid electrolyte and the composite intercalating sulfur electrode The cell consists of a composite sulfur electrode, a bilayer solid electrolyte, and metallic lithium negative electrode (Figure 8.1). The composite sulfur electrode is made of sulfur, conductive carbon, and a lithium ion-intercalating material such as lithiated cobalt oxide (LCO). The bilayer is made of a lithium ion- intercalating layer and a polymer electrolyte. The lithium-ion-intercalation material serves as a solid electrolyte allowing facile transport of lithium ions in the matrix of the active material. The incorporation of lithium intercalating materials into the sulfur electrode has been previously reported by Zhang and Ma. 40,41 They have found that the addition of LCO results in remarkable improvement in the rate performance and cycling stability, attributed to the lithium-ion fast transport and the suppression of the 290 shuttling. The concept of utilizing an intercalating material as an electrolyte is also based on our previous work with Li-S battery where we demonstrated an intercalating mixed conduction membrane (MCM) that exclusively allows the transport of lithium ions and suppresses the polysulfide shuttle. 42 Briefly, a lithium-ion intercalation material conducts lithium ions as the intercalation reaction occurs. Lithium ions diffuse through the bulk of the material functioning as an ion conductor. The electronic conductivity of the material enables the electrochemical reaction of intercalation and the concomitant diffusion of lithium ions to occur efficiently. Figure 8.1. Schematic representation of the solid-state composite-electrode/bi-layer electrolyte lithium- sulfur battery. We chose to employ LCO as the intercalating material because the equilibrium potential window for intercalation is positive with respect to the fully charged sulfur electrode (>2.7 V vs. Li/Li + ). Under these 291 conditions the sulfur electrode can be charged completely without the LCO taking up any of the charge. Furthermore, we chose “nano-particulate” LCO to enhance the interfacial reaction surface area with the sulfur particles in the cathode. It is important to note that no liquid electrolyte was added to the cell and we were relying on the ionic conductivity of LCO to transport lithium ions in and out of the cathode The “bilayer solid electrolyte” is composed of a 30-micron thick layer of LCO (also referred to herein later as the mixed conduction membrane or MCM) coated with a 10-30 micron layer of polymer electrolyte. The polymer electrolyte consists of polyethylene oxide mixed with lithium TFSI salt. The polymer electrolyte layer is placed in direct contact with the lithium electrode, where it serves as a lithium-ion conductor and a separator to prevent an internal electronic short. The bilayer electrolyte configuration takes advantage of the lithium-ion insertion and extraction mechanisms in an intercalating material of the MCM layer. Compared to using conventional solid electrolytes that have only ion conductivity, using a lithium-ion intercalating material as a solid electrolyte has three major advantages: 1) it has a relatively rapid lithium ion conductivity 2) it exclusively transports lithium ions and 3) it does not contribute significantly to the cell’s overall resistance. For example, lithiated cobalt oxide (LiCoO 2) has an electronic conductivity 43,44 of 1 × 10 -3 S cm -1 , and a lithium-ion self-diffusion coefficient 45 of 2 × 10 -9 cm 2 s -1 . Therefore, the facile diffusion of ions is attributed to the intercalation reaction and the ion self-diffusion with the ion transport process driven by the electric field across the MCM layer that is produced during charge and discharge. The MCM layer between the composite cathode and the polymer electrolyte is an important component of the cell as it establishes a pathway for the transport of ions between the cathode and the polymer electrolyte. In addition, MCM has good mechanical properties (strong and flexible) and thus serves as a substrate for the coating of a thin layer of the polymer electrolyte to form the bilayer solid electrolyte. The polymer 292 coating conforms to the MCM providing a good physical contact allowing thin layers of the polymer electrolyte be used. Using a thin layer of the polymer electrolyte is important for lowering the overall internal resistance of the cell, avoiding a short between the anode and cathode, and enhancing rate capability at room temperature. One of the key benefits of the proposed composite electrode-bilayer configuration is the avoidance of flammable organic electrolytes. Furthermore, liquid electrolyte Li-S cells present a self-discharge process that is mainly attributed to the loss of active material from the cathode due to the polysulfide shuttle effect where the sulfur dissolves into the electrolyte to produce high-order polysulfides that tend to diffuse towards the lithium electrode reducing the cell’s capacity. 46,47 With the use of a solid electrolyte the shuttle effect is avoided as the active materials are retained as solids in the composite electrode. In addition, the inclusion of a polymer electrolyte layer that is relatively stable over time, prevents any transport of materials. 48 Further, the bilayer could potentially block the growth of lithium dendrites during repeated cycling. 49,50 We show that the cells can be cycled at room temperature proving the benefit of employing intercalating cathode materials as ionic conductors. From a manufacturing standpoint the bilayer is flexible, air stable, and can be cut to any desired shape, allowing for the possibility of a flexible Li-S battery without liquid organic electrolyte. From the sulfur electrode perspective, using intercalating particles as a solid electrolyte addresses the problems associated with the electrolyte distribution and penetration that are predominant issues when thick electrodes are used with liquid electrolytes. This uniform solid electrolyte distribution and the absence of separators in the cell opens the possibility of building Li-S batteries with high energy density. Furthermore, inorganic fillers can be added to the polymer electrolyte layer to enhance its ionic conductivity and shear modulus. As an example, we have added nano-particulate alumina to the polymer electrolyte layer that improves the 293 ionic conductivity and reduces the polarization during operation. The components of the experimental version of the CEBE Li-S cell add up to give a specific energy of 70.7 Wh Kg -1 (Table 9.1). This value of energy density can be increased significantly with a five to six fold reduction in the thickness of the bilayers and a three times reduction of the fraction of nano-particulate material in the cathode. However, we have not focused on demonstrating these improvements in this work. Finally, the composite-bilayer configuration yields a low internal resistance compared to other solid-sate approaches because of the mixed conductivity of the intercalating materials and the use of very thin layers of the polymer electrolyte. Table 8.1. Energy density at the cell-level for the CEBE Li-S cell. Cathode Electrolyte Anode Weight of Cathode (mg) Loading of sulfur (mg) Weight of Solid Electrolyte (mg) Weight of Anode (mg) Capacity (mAh g -1 ) Average Voltage (V) Energy Density (Wh kg -1 ) Sulfur LiCoO 2/PE O-Al 2O 3 (Bilayer) Li 14.4 2.6 39.86 46 1481 1.84 70.7 Results and Discussion Bilayer electrolyte Figure 8.2a shows a photograph of the bilayer after coating the MCM with PEO-based electrolyte. In the micrograph of the cross-section (Figure 8.2b), we can observe the densely LCO particles with an 294 approximate thickness of 80 micrometers coated with a uniform thin PEO-LiTFSI layer of about 30 micrometers. Figure 8.2. (a) MCM coated with PEO, an example of the bilayer solid electrolyte. (b) Cross-section scanning electron microscopy image of the bilayer. Discharge profile and impedance response of CEBE Li-S cell The discharge profile at equilibrium of the CEBE Li-S cell presented a small plateau followed by a sloping region of relatively low capacity and a then a longer duration plateau (Figure 8.3a). The region between 2.6 V to 2.5 V was associated with the formation of high-order polysulfides and the second part commencing at 2.05 V vs Li/Li + to the conversion of sulfur to the low-order polysulfides. The presence of the first region is attributed to the phase separation between sulfur and the higher-order polysulfides. The second wider lower voltage plateau was associated with the reduction and phase separation into Li 2S 2 and Li 2S. The discharge profile observed is similar to that reported for other quasi-solid-state lithium-sulfur batteries. 51–53 We note that the voltage of the first part of the discharge is about 300 mV higher than that for a typical liquid electrolyte Li-S cell (Figure 8.3a). In addition, the second voltage 295 plateau of the CEBE Li-S cell at 2.1 V is lower than that for the liquid electrolyte cell by about 50-100 mV. We attribute these differences in the cell voltage to the activity and solubility differences of the redox materials in the liquid electrolyte cell vs. the solid-state cell. Figure 8.3. (a) Comparison of the discharge curve at C/60 between the CEBE Li-S cell and a liquid electrolyte Li-S cell during the first cycle. (b) A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of discharge for a liquid electrolyte Li-S cell and (c) for the CEBE Li-S cell. (d) Equivalent circuits proposed to describe the Liquid and CEBE Li-S cells. The impedance response (Nyquist plots) as a function of the state-of-charge (SOC) for the liquid cell (Figure 8.3b) can be de-convoluted into two semicircular arcs at the beginning of the discharge process. 296 Below 60% SOC, the semicircle at high frequency gives place to a single semicircle and an inclined line. This change in impedance response may be associated with the soluble polysulfides at high states of charge precipitating to form the lower-order polysulfides. This particular evolution of the impedance response with SOC is in accordance with previously reported EIS characterization of the Li-S liquid cell. 54,55 For the liquid cell we propose using two electrical equivalent circuits to analyze the impedance as a function of SOC (Figure 8.3d). Circuit 1 is used for describing the region prior to the second plateau where the reduction of sulfur and dissolution of the high-order polysulfides occurs. Circuit 1 consists of R 0, the ohmic resistance of the electrode and electrolyte, R 2, the interphase contact resistance related to the electronic conduction between the current collector (carbon, aluminum) and the active material bulk (according to Deng et al. 54 ), CPE 2, the capacitance of the sulfur electrode bulk, R 1, the charge transfer resistance of the electrochemical reaction, CPE 1, the electrode double-layer capacitance, and W o the mass transport impedance. On the other hand, Circuit 2 consists of just R 0, the ohmic resistance, R 1, the charge transfer resistance, CPE 1, the electrode double-layer capacitance, and W o, the Warburg impedance. Circuit 1 transforms to Circuit 2 when the value of R 2 goes to zero. Upon analyzing the impedance data using Circuit 1, we note that R 2 becomes negligibly low below 60% SOC (Figure 8.4a). This change in the value of R 2 is associated with the complete transformation of the sulfur to the soluble polysulfides that diffuse out of the electrode. Thus, there is no longer a resistance associated with the interphase contact between the active material and the current collector. This transformation is notable when the discharge is conducted at the very low rate of C/60. Under these conditions, the polysulfides can diffuse freely into the bulk electrolyte and towards the anode, allowing for the complete dissolution of the polysulfides avoiding any precipitate formation in the cathode. 297 Figure 8.4. (a) Change of the internal resistances as function of the of the state of discharge for the Liquid Li-S cell. (b) Change of the internal resistances as function of the of the state of discharge for the CEBE Li-S cell. In the Nyquist plot of the CEBE Li-S cell we find one semicircle and a sloping line (Figure 8.3c). At the outset, we note that the overall resistance of the CEBE Li-S cell is lower than for the liquid cell. The impedance of the CEBE-Li-S cells may be modeled by Circuit 2 (Figure 8.3d) as discussed with the precipitation step of the liquid electrolyte cell. The absence of two semicircles throughout the discharge process in the CEBE Li-S cell suggests that the interphase resistance between the active material and the current collector is not significant, and that the nano-particles of intercalation materials in the cathode provide sufficient electrical conductivity between the various phases. The value of R 0 is comparable to that observed in the liquid electrolyte cell, confirming that the bilayer electrolyte and the cathode structure are of comparable conductivity as the liquid electrolyte cell. However, we do observe that the cell resistance decreases and increases to a small extent during the discharge process. Figure 8.4b shows the change of R 0 and R 1 as a function of SOC for the CEBE Li-S cell. The change in charge transfer resistance with SOC is consistent with previous reports suggesting that low-order polysulfides present a 298 higher charge transfer resistance compared to the higher order polysulfides. 56 We also observe that the magnitude of charge transfer resistance R 1 is lower for the CEBE Li-S cell compared to a liquid Li-S cell, attributable again to the large interfacial area provided by the nano-particulate intercalation material. In contrast to the CEBE Li-S cell the overall resistance of the liquid cell decreases with SOC (Figures 9.3b and 4a). The latter is attributed to the dissolution and diffusion of the higher order polysulfides towards the anode (especially at low rates). Thus, the impedance results confirm that the CEBE Li-S cell’s ability to operate at room temperature is due to the low interfacial resistance and high conductivity of the bilayer electrolyte. Effect of Galvanostatic cycling. The CEBE Li-S cells were cycled at room temperature between 2.8 V and 1.4 V at various rates. At C/20 the cell delivered a high specific capacity of 1445 mAh/g during the first discharge, demonstrating that the utilization of the sulfur was almost 85% (Figure 8.5a). 1C corresponded to a rate of 1675 mA/g. The remarkable specific capacity of the solid-state Li-S battery at room temperature is attributed to rapid ion transport in the cathode facilitated by the intercalation material with a high lithium-ion diffusion coefficient. The cell could be repeatedly cycled at C/20 rate, and the capacity decreased gradually with about 63% of the capacity (1050 mAh/g) retained even after the fourth cycle. The results of charge/discharge for C/20 and C/8 rate are shown in Figure 8.5a and 8.5b, respectively. When the cells were cycled at C/8 rate, about 50% of the capacity of the sulfur electrode (820 mAh/g) could be utilized, still a significant fraction of the capacity of the cathode. However, the decrease in capacity from C/20 to C/8 suggested that the diffusion of lithium ions was rate limiting. In comparison, the cells with liquid electrolyte when discharged at C/20 delivered only 54% of the theoretical capacity (906 mAh/g). The utilization in the cell with liquid electrolyte is dependent strongly on the rate, the sulfur to carbon ratio, 299 porosity, and accessible surface area due to the precipitation and nucleation of insoluble polysulfides. In contrast, CEBE cells are not affected as significantly by the precipitation and nucleation process due to the absence of soluble species. In addition, the nano-particles contribute to the accessible surface area, enhancing the porosity of the electrode and creating a hierarchical pore architecture with the Ketjen Black. Figure 8.5. (a) Galvanostatic charge/discharge curves of the CEBE Li-S battery for the first and fourth cycle at rate of C/20 (1C=1675 mA g -1 ). (b) Galvanostatic charge/discharge curves of the CEBE Li-S battery for the first and second cycle at C/8. We found that repeated cycling led to steady capacity loss in each cycle. After eight cycles at room temperature at C/20 rate, retained about 15% of the first cycle capacity (Figure 8.7a). We attribute this loss of capacity to arise from the 79% volume increase that occurs during the discharge of the cathode where sulfur is converted to lithium disulfide. The expansion and contraction that occurs during repeated cycling can cause loss of contact between the layers as well as among the particles in the cathode. Such volume changes would tend to affect all conventional solid electrolyte cells more than liquid electrolyte cells, because of the inability of conventional solid electrolyte cells with rigid structures 300 to accommodate the volume changes. If we can build some resilience in the cathode layer and electrolyte, longer cycle life will be achievable. The impedance data of the cell suggests that the inter- particulate contacts are being lost with cycling, leading to about 350 times increase of the charge- transfer resistance (Figure 8.6). However, with the composite bi-layer type of solid electrolyte cell pursued here, the polymer electrolyte layer can potentially buffer the changes through elastic deformation. Figure 8.6. Impedance response comparison of the CEBE Li-S cell before cycling (depicted with red circles) and after eight cycles (depicted with black squares). Beneficial Effects of Alumina Addition to the Polymer Electrolyte Layer. We recognized that the polymer electrolyte layer does not have the same mechanical strength as the MCM layer. Thus, the polymer electrolyte layer could be irreversibly deformed during the volume changes accompanying cycling. Literature reports suggested that adding alumina as an inorganic filler into PEO increased the mechanical strength and improved the ionic conductivity 32,57,58 The strengthening was caused by the adhesion of the filler to the macromolecular chains that led to immobilization of the polymer chains. 57,59 . It has been reported that solid-state polymer electrolytes in contact with lithium tend to form stable interfaces. 60,61 Moreover, the utilization of inorganic fillers in the polymer electrolyte 301 layer as Al 2O 3 decreases the interfacial resistance at the lithium electrode. 62,63 In addition, it has been reported that 10 wt% of Al 2O 3 can enhance the ionic conductivity 64 from 1 × 10-8 up to 1 × 10-5 S cm- 1 at 30 °C. Therefore, we dispersed alumina nanoparticles (10 wt %) in the polymer electrolyte slurry and then cast it onto the MCM layer. The CEBE Li-S cells with a MCM/PEO-LiTFSI-Al 2O 3 bilayer were cycled at C/20 between 2.8 V and 1.4 V at room temperature. The CEBE-Al2O3 Li-S cell when discharged at C/20 rate, achieved almost 88% (1481 mAh/g) of the theoretical capacity (Figure 8.7a), much like the cells that did not have alumina in the polymer electrolyte. However, even after eight cycles (Figure 8.7b) the cell with the alumina containing PEO retained 60% of the first cycle capacity (901 mAh/g) about 5 times larger capacity compared to the cells without alumina. These results suggested that the polymer electrolyte layer could be strengthened to mitigate some of the capacity loss. Figure 8.7. (a) Specific Capacity against cycle number plot comparison between the CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI at C/20 (depicted with squares) and a CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI-Al 2O 3 at C/20 (depicted with circles) (b) Galvanostatic charge/discharge curves for the CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI-Al 2O 3 for the first and eight cycle at C/20. 302 Conclusions We have introduced a unique solid-state lithium sulfur cell based on a flexible bilayer made of a lithium- ion intercalation material and a polymer electrolyte. Albeit being a solid-state cell, it can be charged and discharged at room temperature at C/20 to yield as much as 85% of the theoretical capacity. The solid- state cell and the bilayer eliminate the use of flammable electrolytes and obviates the issue of transport of polysulfides across the cell. The bilayer is flexible, easy to fabricate, and can be shaped into any cell geometry unlike many of the conventional solid-state cells. The solid-state bi-layer configuration results in a lower resistance compared to the liquid electrolyte cell because of the highly conducting mixed conduction membrane and the thin layer of polymer electrolyte. The addition of nanoparticles of intercalation materials into the sulfur electrode ensures good utilization of the active materials even at C/8 rate yielding >50% of the theoretical capacity. The ability to sustain rates as high as C/8 is attributed to the rapid lithium ion transport and the large interfacial area offered by the nanoparticulate intercalating material. The repeated cycling of these cells leads to capacity loss because of the inability of the cell to accommodate the large volume changes accompanying the discharge at the sulfur electrode that leads to loss of inter-particle contact. The addition of inorganic fillers like alumina into the polymer electrolyte layer reduces the rate of capacity loss, proving that the electrolyte layer also needs to be mechanically robust. Electrode composites that can elastically deform to accommodate the changes in volume of the solid-state cathode would be needed to address the capacity fade. What is presented here is a proof-of-concept of a new approach and additional effort will be needed to optimize the sulfur electrode composition as increasing the areal loading of the active material for achieving higher energy density cells. Alternate mixed conductors that would not only provide fast ion transport but also cathode capacity would be beneficial in retaining the high energy density of the lithium-sulfur 303 cells. The results presented here demonstrate that the composite electrode/bilayer electrolyte architecture is a new pathway for the further development of lithium-sulfur solid-state batteries. 304 A mass-efficient CEBE Li-S cell (TAG proposal) Abstract To develop a superior and practical CEBE Li-S cell it is necessary to improve its energy density and cycle- life. Building a high-energy cell requires a decrease in content of all the cell’s inactive materials and the increase of sulfur content in the electrode. In addition, an economically viable cell would require the reduction of the intercalating nanoparticles in the composite sulfur electrode. Therefore, we have revisited the original design of CEBE Li-S cell reported in 2020, and improved its energy density by implementing a mass-efficient cell design. Moreover, we have verified the suppression of the polysulfide shuttle by utilizing our electrochemical in-house method known as the polysulfide shuttle current 10 , and measured the impedance response as a function of SOC to understand the phenomena occurring in cell. Finally, GCD cycling was tested to verify the mass-efficient cell’s capacity retention and energy density. At the cell-level (excluding the cell casing) the mass-efficient CEBE Li-S cell achieved a remarkable specific energy of 220 Wh kg -1 at C/9 being a significant improvement from the previous cell design and becoming a new milestone in the development of a practical solid-state lithium-sulfur batteries. Cell optimization For building a mass-efficient cell the percentage of sulfur in the composite electrode was increased from 43 to 54 wt%. In addition, the intercalating LCO nanoparticles were reduced from 33 to 18%. For the bi- layer solid-electrolyte, the mixed conduction solid-electrolyte (LCO) layer thickness was significantly decreased from 60 to 10 μm and from 30 to 10 μm for the PEO:Al 2O 3 polymer electrolyte layer, being 72% lower in weight from the previous bi-layer. Finally, the lithium electrode mass was decreased to more than 90% by utilizing an extremely thin lithium electrode (<10 μm). 305 Shuttle current measurement for the CEBE Li-S cell To verify the suppression of the polysulfide shuttling in the CEBE Li-S cell, we utilized our previous in- house method denoted as the polysulfide shuttle current 10 . Briefly, at a certain state of charge of the cell when the soluble polysulfides are present (liquid cell) under open-circuit conditions, the potential of the cathode decreases steadily due to the diffusion of the low order polysulfides that are reduced at the lithium electrode towards the cathode. Maintaining the electrode potential of the sulfur electrode at a constant value, will maintain the diffusion fluxes involved in the shuttling process constant. The constant potential is reached by the passage of electric current that will cause the re-oxidation of the reduced polysulfides arriving at the cathode. Therefore, the required steady-state current to maintain the potential is a direct measure of the shuttling rate process, we call this current the polysulfide shuttle- current or simply the shuttle current. A shuttle current closer to 0 indicates the absence of the polysulfide shuttling in the Li-S cell. In a solid- state no polysulfide shuttling is expected to be present do the lack of liquid electrolyte. Figure 8.8a shows the potentiosatic observed shuttle current as function of time at 2.7 V vs Li + /Li for the mass-efficient CEBE Li-S cell. The shuttle current tends to 0 after 1 hour of holding the voltage confirming the suppression of the polysulfide shuttling at that potential. Figure 8.8b shows the steady-state specific shuttle current after holding for 1 hour the potential at various potentials vs Li + /Li during the discharge of the mass efficient CEBE Li-S cell. The shuttle current values tend to 0 through all the discharge potential window verifying that the solid-state Li-S cell does not exhibit any polysulfide shuttling. 306 Figure 8.8 Shuttle current measurement of the mass-efficient CEBE Li-S cell (TAG). a) Potentiostatic specific current as a function time at 2.7 V vs Li + /Li. b) Steady-state specific shuttle current as a function of potential in the mass-efficient CEBE Li-S cell. EIS characterization of the mass-efficient CEBE Li-S cell The impedance response of the mass-efficient CEBE Li-S cell was measured as function of SOC. Figure 8.9 shows the GCD profiles and the steps where EIS was measured during charge and discharge (denoted with yellow circles). A three-dimensional representation of the Nyquist plots as function of %SOC during discharge and charge is shown in Figures 8.10a and 8.10b respectively. During discharge at 75% SOC and high frequencies the real part of the impedance response (Zre) exhibits a resistance of 120 ohms cm -2 that then decreases to 95 and 72 ohm cm -2 at 50% and 18% SOC respectively (Figure 8.11a). The decrease in resistance at high frequencies during discharge is associated to the decrease in ohmic resistance. Assuming, that the solid-state layers do not exhibit a significant change in conductivity, the data suggests that the interlayer mechanical contact is enhanced due to volume expansion the occurs at the sulfur electrode during discharge. Furthermore, the overall resistance decreases as the cell is discharged, particularly at low frequencies where the main differences in magnitude are present, linked to the slow transport processes occurring in the cell. The resistance decrease at low frequencies, suggests a better 307 interconnection of the insulting active material with the intercalating LCO nanoparticles and conducting carbon due to the morphological changes from the aggressive volume expansion of more than 79% that the cathode undergoes during discharge. Moreover, during charge the positive electrodes contracts, where the Li 2S gets oxidized into elemental sulfur. Due to the additional pores/voids created during the first expansion of the sulfur electrode, the cell exhibits a significant lower Zre during charge that tends to reach a constant value beyond 50% SOC (Figure 8.11b), suggesting that the cell reaches an “equilibrium” morphology. Figure 8.9 Galvanostatic charge-discharge profiles for the mass-efficient CEBE Li-S cell at C/20 with a charging cut-off of 2.7 V vs Li + /Li. EIS was measured as function of SOC at OCV (denoted with yellow circles). 308 Figure 8.10 A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of discharge the mass-efficient CEBE Li-S cell during a) discharge and b) charge. Figure 8.11 Real part of the impedance response Zre against frequency for the mass-efficient CEBE Li-S cell during the at 18%, 50%, and 75% SOC during a) discharge and b) charge. Galvanostatic cycling of the mass-efficient CEBE Li-S cell. GCD cycling was carried for the mass-efficient CEBE Li-S cell. Figure 8.12a and 8.12b shows the specific capacity as function of cycle number and their corresponding GCD curves cycled at a constant rate of 309 C/9 where 1C = 1675 mA g -1 . Although the mass fraction of sulfur was increased to 54% and the LCO nanoparticles where reduced to 18%, the cell still exhibited a remarkable specific capacity of 1211 mAh g -1 at room temperature, and after 25 cycles the cell reached a constant capacity value of approximately 400 mAh g -1 . The capacity fade during cycling is attributed to the loss of mechanical contact of all the layers caused by constant expansion and contraction occurring in the composite sulfur electrode. Additional optimization for the cell construction design is required to reduce the capacity fade. Nevertheless, the mass-efficient CEBE Li-S cell exhibited an impressive specific energy of 219.7 Wh kg -1 at the cell-level. Table 9.2 resumes and compares all the cell components and performance of the 2020 CEBE Li-S cell and the new mass-efficient CEBE Li-S cell. The optimization of the CEBE Li-S cell lead to a 280% energy enhancement from the previous cell reported in 2020. Figure 8.13 shows an energy comparison of the first discharge curves between the mass-efficient CEBE Li-S cell and the previous CEBE Li-S cell. Figure 8.12 (a) Specific capacity against cycle number plot for the mass-efficient CEBE Li-S cell (b) Corresponding galvanostatic charge/discharge curves for the mass-efficient CEBE Li-S cell in at C/9. Table 9.2 Cell parameters and energy performance comparison of the CEBE Li-S cell baseline from 2020 against the mass-efficient 2022 TAG cell. 310 2020 Baseline 2022 TAG proposal Fraction of Active Material (Sulfur) 43% 54% Fraction of LCO nano-particles 33% 18% Sulfur loading [mg cm -2 ] 1.3 1.05 Lithium loading [mg] 46 5 Bilayer LCO-PEO:Al 2O 3 Weight [mg] 39.86 11 Cell Weight [g] 0.1002 0.025 Negative/positive areal capacity ratio [N/P] 30.94 3.9 Specific Capacity [mAh g -1 ] 1481 1211 Nominal avg Voltage [V vs Li + /Li] 2 2 Specific Energy at the cell-level [Wh kg -1 ] 77.5 219.7 Figure 8.13 Galvanostatic discharge curve (potential vs specific energy at the cell-level) comparison of the CEBE Li-S cell baseline from 2020 against the mass-efficient 2022 TAG cell. 311 References 1. Huang, L. et al. Electrode Design for Lithium–Sulfur Batteries: Problems and Solutions. Adv. Funct. Mater. 30, 1910375 (2020). 2. Chen, X., Hou, T., Persson, K. A. & Zhang, Q. Combining theory and experiment in lithium–sulfur batteries: Current progress and future perspectives. Materials Today 22, 142–158 (2019). 3. Zhao, H. et al. A review on anode for lithium-sulfur batteries: Progress and prospects. Chemical Engineering Journal 347, 343–365 (2018). 4. Kang, W. et al. A review of recent developments in rechargeable lithium–sulfur batteries. Nanoscale 8, 16541–16588 (2016). 5. Zhang, X. et al. Advances in lithium—sulfur batteries. Materials Science and Engineering: R: Reports 121, 1–29 (2017). 6. Rauh, R. D., Abraham, K. M., Pearson, J. K., Surprenant, J. K. & Brummer, S. B. A Lithium/Dissolved Sulfur Battery with an Organic Electrolyte. J. Electrochem. Soc. 126, 523 (1979). 7. Fotouhi, A., Auger, D., O’Neill, L., Cleaver, T. & Walus, S. Lithium-Sulfur Battery Technology Readiness and Applications—A Review. Energies 10, 1937 (2017). 8. Rong, G. et al. Liquid-Phase Electrochemical Scanning Electron Microscopy for In Situ Investigation of Lithium Dendrite Growth and Dissolution. Adv. Mater. 29, 1606187 (2017). 9. Scheers, J., Fantini, S. & Johansson, P. A review of electrolytes for lithium–sulphur batteries. Journal of Power Sources 255, 204–218 (2014). 10. Moy, D., Manivannan, A. & Narayanan, S. R. Direct Measurement of Polysulfide Shuttle Current: A Window into Understanding the Performance of Lithium-Sulfur Cells. J. Electrochem. Soc. 162, A1–A7 (2015). 11. Wild, M. et al. Lithium sulfur batteries, a mechanistic review. Energy Environ. Sci. 8, 3477–3494 (2015). 12. Yamin, H. & Peled, E. Electrochemistry of a nonaqueous lithium/sulfur cell. Journal of Power Sources 9, 281–287 (1983). 13. Judez, X. et al. Review—Solid Electrolytes for Safe and High Energy Density Lithium-Sulfur Batteries: Promises and Challenges. J. Electrochem. Soc. 165, A6008–A6016 (2018). 14. Manthiram, A., Yu, X. & Wang, S. Lithium battery chemistries enabled by solid-state electrolytes. Nat Rev Mater 2, 16103 (2017). 15. Lin, Z. & Liang, C. Lithium–sulfur batteries: from liquid to solid cells. J. Mater. Chem. A 3, 936–958 (2015). 16. Umeshbabu, E., Zheng, B. & Yang, Y. Recent Progress in All-Solid-State Lithium−Sulfur Batteries Using High Li-Ion Conductive Solid Electrolytes. Electrochem. Energ. Rev. 2, 199–230 (2019). 17. Xu, R. et al. Cathode-Supported All-Solid-State Lithium–Sulfur Batteries with High Cell-Level Energy Density. ACS Energy Lett. 4, 1073–1079 (2019). 18. Yamada, T. et al. All Solid-State Lithium–Sulfur Battery Using a Glass-Type P 2 S 5 –Li 2 S Electrolyte: Benefits on Anode Kinetics. J. Electrochem. Soc. 162, A646–A651 (2015). 312 19. Zhang, Q. et al. Rational design of multi-channel continuous electronic/ionic conductive networks for room temperature vanadium tetrasulfide-based all-solid-state lithium-sulfur batteries. Nano Energy 57, 771–782 (2019). 20. Suzuki, K. et al. High Cycle Capability of All-Solid-State Lithium–Sulfur Batteries Using Composite Electrodes by Liquid-Phase and Mechanical Mixing. ACS Appl. Energy Mater. 1, 2373–2377 (2018). 21. Suzuki, K. et al. Composite Sulfur Electrode Prepared by High-Temperature Mechanical Milling for use in an All-Solid-State Lithium–Sulfur Battery with a Li3.25Ge0.25P0.75S4 Electrolyte. Electrochimica Acta 258, 110– 115 (2017). 22. Han, F. et al. High-Performance All-Solid-State Lithium–Sulfur Battery Enabled by a Mixed-Conductive Li 2 S Nanocomposite. Nano Lett. 16, 4521–4527 (2016). 23. Gao, Y. et al. Polymer–inorganic solid–electrolyte interphase for stable lithium metal batteries under lean electrolyte conditions. Nat. Mater. 18, 384–389 (2019). 24. Wang, Y., Sahadeo, E., Rubloff, G., Lin, C.-F. & Lee, S. B. High-capacity lithium sulfur battery and beyond: a review of metal anode protection layers and perspective of solid-state electrolytes. J Mater Sci 54, 3671–3693 (2019). 25. Sun, Y.-Z., Huang, J.-Q., Zhao, C.-Z. & Zhang, Q. A review of solid electrolytes for safe lithium-sulfur batteries. Sci. China Chem. 60, 1508–1526 (2017). 26. Zhao, Y. et al. Polymer Electrolytes for Lithium/Sulfur Batteries. Membranes 2, 553–564 (2012). 27. Peng, H.-J., Huang, J.-Q., Cheng, X.-B. & Zhang, Q. Review on High-Loading and High-Energy Lithium-Sulfur Batteries. Adv. Energy Mater. 7, 1700260 (2017). 28. Rana, M. et al. Review on areal capacities and long-term cycling performances of lithium sulfur battery at high sulfur loading. Energy Storage Materials 18, 289–310 (2019). 29. Kim, C.-S. et al. Importance of open pore structures with mechanical integrity in designing the cathode electrode for lithium–sulfur batteries. Journal of Power Sources 241, 554–559 (2013). 30. Song, J. et al. Flexible freestanding sandwich-structured sulfur cathode with superior performance for lithium–sulfur batteries. J. Mater. Chem. A 2, 8623–8627 (2014). 31. Dirican, M., Yan, C., Zhu, P. & Zhang, X. Composite solid electrolytes for all-solid-state lithium batteries. Materials Science and Engineering: R: Reports 136, 27–46 (2019). 32. Das, S. & Ghosh, A. Ion conduction and relaxation in PEO-LiTFSI-Al 2 O 3 polymer nanocomposite electrolytes. Journal of Applied Physics 117, 174103 (2015). 33. Song, C. et al. Stable and Fast Lithium–Sulfur Battery Achieved by Rational Design of Multifunctional Separator. Energy Environ. Mater. 2, 216–224 (2019). 34. Kim, S., Kim, J., Cho, S. & Lee, S. All‐Solid‐State Printed Bipolar Li–S Batteries. Adv. Energy Mater. 9, 1901841 (2019). 35. Zhu, P. et al. Flexible electrolyte-cathode bilayer framework with stabilized interface for room- temperature all-solid-state lithium-sulfur batteries. Energy Storage Materials 17, 220–225 (2019). 36. Fu, K. (Kelvin) et al. Three-dimensional bilayer garnet solid electrolyte based high energy density lithium metal–sulfur batteries. Energy Environ. Sci. 10, 1568–1575 (2017). 313 37. Tao, X. et al. Solid-State Lithium–Sulfur Batteries Operated at 37 °C with Composites of Nanostructured Li 7 La 3 Zr 2 O 12 /Carbon Foam and Polymer. Nano Lett. 17, 2967–2972 (2017). 38. Zhu, Y., Li, J. & Liu, J. A bifunctional ion-electron conducting interlayer for high energy density all-solid- state lithium-sulfur battery. Journal of Power Sources 351, 17–25 (2017). 39. Nagata, H. & Chikusa, Y. An all-solid-state lithium–sulfur battery using two solid electrolytes having different functions. Journal of Power Sources 329, 268–272 (2016). 40. Zhang, K. et al. A novel class of functional additives for cyclability enhancement of the sulfur cathode in lithium sulfur batteries. Inorg. Chem. Front. 5, 2013–2017 (2018). 41. Ma, W. & Xu, Q. Lithium cobaltate: a novel host material enables high-rate and stable lithium–sulfur batteries. Rare Met. 37, 929–935 (2018). 42. Moy, D. & Narayanan, S. R. Mixed Conduction Membranes Suppress the Polysulfide Shuttle in Lithium- Sulfur Batteries. J. Electrochem. Soc. 164, A560–A566 (2017). 43. Wei, G., Hass, T. E. & Goldner, R. B. Thin films of lithium cobalt oxide. Solid State Ionics 58, 115–122 (1992). 44. Tukamoto, H. & West, A. R. Technical Papers Solid-State Science and Technology. J. Electrochem. Soc. 144, 6 (1997). 45. Julien, C., Camacho-Lopez, M. A., Escobar-Alarcon, L. & Haro-Poniatowski, E. Fabrication of LiCoO2 thin- film cathodes for rechargeable lithium microbatteries. Materials Chemistry and Physics 68, 210–216 (2001). 46. Knap, V., Stroe, D.-I., Swierczynski, M., Teodorescu, R. & Schaltz, E. Investigation of the Self-Discharge Behavior of Lithium-Sulfur Batteries. J. Electrochem. Soc. 163, A911–A916 (2016). 47. Barchasz, C. et al. Lithium/Sulfur Cell Discharge Mechanism: An Original Approach for Intermediate Species Identification. Anal. Chem. 84, 3973–3980 (2012). 48. Zhang, J.-G., Xu, W. & Henderson, W. A. Lithium Metal Anodes and Rechargeable Lithium Metal Batteries. vol. 249 (Springer International Publishing, 2017). 49. Wang, C. et al. Suppression of Lithium Dendrite Formation by Using LAGP-PEO (LiTFSI) Composite Solid Electrolyte and Lithium Metal Anode Modified by PEO (LiTFSI) in All-Solid-State Lithium Batteries. ACS Appl. Mater. Interfaces 9, 13694–13702 (2017). 50. Chen, L. et al. PEO/garnet composite electrolytes for solid-state lithium batteries: From “ceramic-in- polymer” to “polymer-in-ceramic”. Nano Energy 46, 176–184 (2018). 51. Han, D.-D. et al. Lithiophilic gel polymer electrolyte to stabilize the lithium anode for a quasi-solid-state lithium–sulfur battery. J. Mater. Chem. A 6, 18627–18634 (2018). 52. Pang, Q. et al. Tuning the electrolyte network structure to invoke quasi-solid state sulfur conversion and suppress lithium dendrite formation in Li–S batteries. Nat Energy 3, 783–791 (2018). 53. Zhong, H., Wang, C., Xu, Z., Ding, F. & Liu, X. A novel quasi-solid state electrolyte with highly effective polysulfide diffusion inhibition for lithium-sulfur batteries. Sci Rep 6, 25484 (2016). 54. Deng, Z. et al. Electrochemical Impedance Spectroscopy Study of a Lithium/Sulfur Battery: Modeling and Analysis of Capacity Fading. J. Electrochem. Soc. 160, A553–A558 (2013). 314 55. Ahn, W., Kim, K.-B., Jung, K.-N., Shin, K.-H. & Jin, C.-S. Synthesis and electrochemical properties of a sulfur- multi walled carbon nanotubes composite as a cathode material for lithium sulfur batteries. Journal of Power Sources 202, 394–399 (2012). 56. Ji, X., Lee, K. T. & Nazar, L. F. A highly ordered nanostructured carbon–sulphur cathode for lithium–sulphur batteries. Nature Mater 8, 500–506 (2009). 57. Long, L., Wang, S., Xiao, M. & Meng, Y. Polymer electrolytes for lithium polymer batteries. J. Mater. Chem. A 4, 10038–10069 (2016). 58. Yap, Y. L., You, A. H., Teo, L. L. & Hanapei, H. Inorganic Filler Sizes Effect on Ionic Conductivity in Polyethylene Oxide (PEO) Composite Polymer Electrolyte. Int. J. Electrochem. Sci. 8, 10 (2013). 59. Fan, L. Effect of modified SiO2 on the properties of PEO-based polymer electrolytes. Solid State Ionics 164, 81–86 (2003). 60. Peled, E., Golodnitsky, D., Ardel, G. & Eshkenazy, V. THE SEI MODEL-APPLICATION TO LITHIUM-POLYMER ELECTROLYTE BATTERIES. Electrochimica Acta 40, 2197–2204 (1995). 61. Appetecchi, G. B., Croce, F., Persi, L., Ronci, F. & Scrosati, B. Transport and interfacial properties of composite polymer electrolytes. Electrochimica Acta 45, 1481–1490 (2000). 62. Peled, E. & Menkin, S. Review—SEI: Past, Present and Future. J. Electrochem. Soc. 164, A1703–A1719 (2017). 63. Appetecchi, G. B., Alessandrini, F., Duan, R. G., Arzu, A. & Passerini, S. Electrochemical testing of industrially produced PEO-based polymer electrolytes. Journal of Power Sources 101, 42–46 (2001). 64. Croce, F., Appetecchi, G. B., Persi, L. & Scrosati, B. Nanocomposite polymer electrolytes for lithium batteries. Nature 394, 456–458 (1998). 315 Chapter 9 – The Role of Functionalized Conducting Polymer Binders in Addressing the Technical Challenges of Lithium-Sulfur Batteries Abstract Lithium-Sulfur batteries are promising for the next-generation of low-cost, high-energy rechargeable batteries. However, the commercial adoption of this battery has been limited by poor rate capability and cycle life. To improve the performance of Li-S cells, we have investigated the role of two n-dopable conducting polymers, N2200 and N2200-OE, as binders for the sulfur cathode. The electron and ion transport properties of these polymer binders, along with their strong affinity for polysulfide, results in significantly enhanced performance compared to cells with the traditional insulating PVDF binder. The role of these conducting polymer binders in enhancing the performance of the cells is analyzed using impedance spectroscopy, shuttle current measurements, conductivity measurements, UV-Vis absorption spectroscopy, and GIWAXS studies. The conducting polymer binders reduce the cell impedance by a factor of four leading to a marked improvement in rate capability. The shuttling of the polysulfides is notably curtailed by the strong interaction of the polysulfides with the backbone of the polymer binders resulting in an impressive capacity retention of 82% after 500 cycles. These studies demonstrate the benefit of using tailored polymer binders at the sulfur electrode in addressing the limitations of lithium-sulfur batteries. Introduction The lithium-sulfur (Li-S) battery is a promising next generation energy storage technology to replace the current lithium-ion battery (LIB) due to the high theoretical specific capacity of the sulfur electrode at 1672 mAh g −1 and the high gravimetric energy density of the Li-S cell at 2500 Wh kg −1 . 1, 2 In addition, 316 sulfur is substantially more abundant and less expensive than the cathode materials currently used in current commercial LiBs, such as nickel and cobalt. 3-5 However, the commercial adoption of Li-S batteries has been hindered by several issues, including poor cycle life, low discharge rate capability, and low coulombic efficiency. 6, 7 These performance limitations arise largely from the insulating nature of sulfur, the formation and shuttling of soluble polysulfides, and the large volume expansion of the sulfur electrode. 8 In the last decade, a range of material strategies have been explored to address these shortcomings with varying degrees of success. 9-13 Among these strategies include the use of various types of carbon materials, 14-17 unique architectures and coatings, 18-22 and tailored barrier layers 23-26 to trap the soluble polysulfides at the sulfur electrode. Mesoporous carbon has been reported as a strategy to decrease the polysulfide shuttle due to the small pore diameter. 27, 28 However, the large surface area of the carbon often allows the polysulfides to escape into solution. To address this problem, composites of the carbon materials with polymers have also been widely investigated. 29 Unlike the commonly used polyvinylidene fluoride (PVDF) binder, polymer materials such as polyvinylpyrrolidone can serve not only as a binder but also as a trap for polysulfides. 30 If the chemical groups in the polymer can interact strongly with the polysulfides, the shuttling can be reduced. In this regard, carbonyl-containing functional groups such as esters and amides have been found to sequester the polysulfides through a Li-O interaction. 31 Thus, polymers with good mechanical properties such as catechol-conjugated chitosan sulfate and gum arabic that have an affinity to the polysulfides have been used as binders leading to enhanced cycle life in Li-S batteries. 32, 33 Although such polymers are effective as binders and can trap the polysulfides, their non-conductive nature impedes electron transport through the composite structure. Thus, the addition of such non- 317 conductive polymer coatings, although beneficial with respect to trapping the polysulfide, can also increase the resistance of the carbon composites and reduce sulfur utilization and rate capability.30 Consequently, significant research has focused on using electronically conductive polymers as binders for the sulfur electrode. To this end, several p-dopable polymers (i.e., polymers capable of producing free charge carriers upon oxidation) such as polythiophene and its derivatives, PEDOT, polyaniline, and polypyrrole have been utilized with the goal of simultaneously enhancing electron transport in the sulfur electrode and reducing the polysulfide shuttling.34-40 Furthermore, there were attempts to develop dual ionic and electronically conducting polymer binder for sulfur electrodes by copolymerizing poly(3,4- ethylenedioxythiophene) (PEDOT) with polyethylene glycol (PEG).41 While these polymers can offer electronic conductivity and stability on p-doping in the potential range of 3.1 V to 4.3 V vs. Li + /Li, they remain insulating below 3.1 V vs. Li + /Li. 42, 43 This voltage window of electronic conductivity and stability of these p-dopable conducting polymers are well matched with the cycling window of cathode materials such as nickel cobalt oxide. 42, 44-46 However, as the voltage window for the cycling of the sulfur electrode is between 1.7 and 3.0 V vs. Li + /Li, it is important to choose a polymer that is an electronic conductor in this potential range. Although p-dopable conductive polymers are more common, there is no known p- dopable conjugated polymer that will remain in the p-doped state below ~2.3 V vs. Li + /Li. 47-49 Any p- dopable polymer with a lower oxidation potential would be remarkably unstable in air and profoundly difficult to synthesize and handle. Therefore, n-dopable polymers are selected as the primary candidates in this study. There are a number of n-dopable polymers that will exist in the conducting state between 2.6 – 1.6 V vs. Li + /Li. A well-known example is N2200 (or P(NDI2OD-T2)); poly{[N,N′-bis(2-octyldodecyl)-naphthalene- 1,4,5,8-bis(dicarboximide)-2,6-diyl]-alt-5,5′-(2,2′-bithiophene)} which has an onset of reduction at ~2.6 318 V vs. Li + /Li and remains conductive down to potentials below 1.6 V. 50, 51 Furthermore, the carbonyl groups in the polymer backbone should increase the interaction of the polymer with the polysulfides. Incorporation of oligoether groups into the N2200 structure is envisioned to further enhance lithium-ion conduction, and potentially further enhance the interaction with polysulfide. We hypothesized that such key structural features could improve both the discharge capacity and cycle life. Therefore, in the present study we focus on characterizing the properties and benefits of a well-known polymer, N2200. This polymer is n-dopable (i.e., capable of producing polaronic charge carriers by reduction) and can provide several anticipated benefits to the sulfur electrode, such as: 52 (1) electrically connecting the sulfur particles by virtue of its electronical conducting and electrochemically stable in the potential range between 1.7 V and 3.0 V vs Li + /Li, (2) providing strong affinity of the polymer to polysulfides through several carbonyl groups on the polymer backbone (Fig. 9.1), (3) possessing ionic transport properties that can be tuned by producing copolymers of N2200 that contain oligoether side-chains to form N2200-OE (or PNDI (OD:OE)), (structures shown in Fig. 9.1). 44 We have synthesized N2200 and a copolymer of N2200 (PNDI (OD:OE)) in an 80:20 ratio and characterized the structural, electron-transport and ion-transport properties of these polymers, and measured the interaction of the polymers with the polysulfides. By comparing the cycling performance and rate capability of lithium-sulfur cells with N2200, N2200-OE, and PVDF, we could verify the multiple anticipated benefits of these functionalized electronically conducting polymers as binders in the sulfur electrode. 319 Results and Discussion Figure 9.1. Polymers investigated for binders in Lithium-Sulfur (Li-S) batteries. Electrochemical Cycling of Polymer Thin films Thin films of N2200 and N2200-OE were electrochemically characterized in a three-electrode cell with Li-S battery electrolyte in an argon glovebox between 3.0 and 1.7 V vs Li + /Li, a potential range relevant for the sulfur electrode. Initial thin film cycling was performed at 100 mV s −1 (Fig. 9.2 a,b). In the first cycle, the cyclic voltammogram of N2200 showed a reduction peak at 2.08 V and two oxidation peaks at 2.33 and 2.55 V vs Li + /Li. The capacity of N2200 increased significantly with subsequent cycles from 5.4 nAh in the first cycle to 92 nAh after 20 cycles. The cyclic voltammogram on the 25th cycle showed that the reduction peak shifted slightly to 1.97 V and the oxidation peaks shifted to 2.39 and 2.53 V vs Li + /Li. The CV data was stable after the 20th cycle with no peak shifts or further change in 320 capacity. The slow increase in capacity is likely associated with solvent swelling of the polymer network, which facilitates ion transport through the film. The CV data for N2200-OE, on the other hand, showed a small reduction peak at 2.55 V and a major reduction peak at 2.15 V vs Li + /Li (Fig. 9.2b). A small oxidation peak appears at 2.37 V and a major oxidation peak appears at 2.54 V vs Li + /Li. No major changes in the CV data were observed in the subsequent cycles. The slight shifts to slightly more positive potentials in the major cathodic peak for both polymers in subsequent cycles is attributed to the enhanced conductivity of the polymers after electrochemical doping that led to lowering of the overpotential for reduction. Unlike N2200, the capacity in the first cycle for N2200-OE was 298 nAh and slightly decreased to 287 and 283 nAh in the second and fifth cycles, respectively. This is likely due to the ability of N2200-OE to swell significantly more compared to N2200, allowing electrolyte to enter the structure and electrochemically dope the polymer film much faster than with N2200. To ensure that electrochemical doping/de-doping of the polymer binders does not limit cathode cycling, we also examined the kinetics of the electrochemical doping based on a series of CV measurements at various scan rates (Fig. 9.2 c,d) from 20 to 100 mV s −1 . The minimal shifts of the redox peaks with increasing scan rates indicate rapid reaction rates. Finally, we examined the long-term electrochemical stability of the polymer thin films by cycling them between 1.7 and 3.0 V vs Li + /Li at 10 mV s −1 for 100 cycles (Fig. 9.2 e,f). Both polymers exhibited stable redox properties in the potential range of 1.7 to 3.0 V vs Li + /Li over multiple cyclic voltammogram scans. 321 Figure 9.2. Cyclic Voltammograms of N2200 and N2200-OE. Initial CV data (a and b), CV data as a function of scan rate (c and d), and long-term cycling (e and f) for N2200 and N2200-OE, respectively. Electronic and Ionic Conductivity of Polymer Thin films. To probe the electron and ion transport properties of N2200 and N2200-OE, we measured the electronic and ionic conductivities of the polymer thin films in 1 M LiTFSI in DOL/DME as a function potential of electrochemical doping. Cycling the polymers at a slow scan rate of 5 mV s −1 (Fig. 9.3a) revealed a cathodic peak at 2.16 V and two anodic peaks at 2.31 V and 2.50 V vs Li + /Li for N2200. In the cathodic region for N2200-OE, a small shoulder appears at 2.64 V and a main peak appears at 2.30 V. For the anodic region, we observed a shoulder at 2.36 V and a main oxidation peak at 2.51 V vs Li + /Li. Electronic conductivity data (Fig. 9.3b) shows that the electronic conductivity of undoped N2200 is 2×10 −6 S cm −1 above 2.5 V vs Li + /Li, consistent with the absence of any significant electrochemical doping 2.0 2.5 3.0 -50 0 50 2.0 2.5 3.0 -100 0 100 2.0 2.5 3.0 -50 0 50 2.0 2.5 3.0 -100 0 100 2.0 2.5 3.0 -10 0 10 2.0 2.5 3.0 -10 0 10 20 f e d c b Current ( A) Potential (V vs Li + /Li) Cycle 1,10 15, 20, 25 100 mV s -1 a 100 mV s -1 Current ( A) Potential (V vs Li + /Li) Cycle 1, 2 3, 4, 5 N2200-OE Current ( A) Potential (V vs Li + /Li) 100, 80, 60, 40, 20 mV s -1 N2200 Current ( A) Potential (V vs Li + /Li) 100, 80, 60, 40, 20 mV s -1 2 nd 60 th 20 th 80 th 40 th 100 th Current ( A) Potential (V vs Li + /Li) 2 nd 60 th 20 th 80 th 40 th 100 th 10 mV s -1 10 mV s -1 Current ( A) Potential (V vs Li + /Li) 322 in the cyclic voltammograms (Fig. 9.3a). The electronic conductivity of N2200 starts to increase below 2.5 V vs Li + /Li which aligns with the onset of electrochemical n-doping of the polymer (Fig. 9.3a) and levels off at 2×10 −4 S cm −1 around 2.2 V. Thus, we observed two orders of magnitude increase in electronic conductivity upon doping of N2200. N2200-OE, on the other hand, has a slightly lower electronic conductivity of 8×10 −7 S cm −1 when undoped at potentials greater than 2.65 V vs Li + /Li. Its electronic conductivity starts to increase at potential below 2.65 V, the potential aligning with the onset of n-doping as observed in Fig. 9.3a, to reach a maximum of 1×10 −4 S cm −1 just below 2.4 V. The electronic conductivity remains relatively constant throughout the potential window tested. Thus, the electronic conductivity of N2200-OE reaches a lower maximum conductivity compared to N2200, but the onset of the conductivity increase occurs at a higher potential compared to N2200. The introduction of oligoether side chains seems to lower electronic conductivity which has been observed in our previous work. 44 The ionic conductivity data (Fig. 9.3c) revealed that N2200 has ionic conductivity of 4×10 −10 S cm −1 when undoped. Upon doping, the ionic conductivity increases to 7×10 −9 S cm −1 at 1.7 V vs Li + /Li. Fig. 9.3c shows that N2200-OE has an ionic conductivity of 6×10 −9 S cm −1 at 2.8 V (when undoped) and 2×10 −7 S cm −1 at 1.7 V vs Li + /Li. Thus, the addition of the oligoether side chains appears to contribute to over one-order of magnitude improvement in ionic conductivity across the tested potential range. This increases in ionic conductivity is consistent with the structural changes that occur during n-doping of the polymer, as observed in the GIWAXS studies, discussed below. Furthermore, the increased ionic conductivity of N2200-OE compared to N2200 is well supported by the swelling behavior of these polymers. Almost all organic polymer swell to some degree when placed in organic solvents like the DOL:DME electrolyte used here, and the OE side chains significantly enhance that process. The addition of the oligoether sidechains increases the % mass on swelling in the DOL:DME solvent. We observed that neutral N2200 323 had a 100% mass uptake, while N2200-OE had a 220% mass uptake. The solvent swollen binder has many pathways for ionic motion through the polymer network and thus should show enhanced ionic conductivity. Solvent swelling upon electrochemical doping is likely also responsible for the increase in ionic conductivity upon doping. Figure 9.3. (a) Cyclic voltammograms of N2200 and N2200-OE at 5 mV s −1 . (b) Electronic and (c) ionic conductivity of N2200 and N2200-OE as a function of potential. 2.0 2.5 3.0 -10 0 10 1.6 2.0 2.4 2.8 10 -6 10 -5 10 -4 1.6 2.0 2.4 2.8 10 -9 10 -7 N2200 N2200-OE N2200 N2200-OE Potential (V vs Li + /Li) Current ( A) Potential (V vs Li + /Li) N2200 N2200-OE 5 mV s -1 c b  electronic (S cm -1 ) a  ionic (S cm -1 ) 324 Structural Characterization GIWAXS was used to understand how the structures of N2200 and N2200-OE evolve during the cycling process. GIWAXS provides information about lattice constants, scattering intensity, and the orientation of the polymer with respect to the substrate. N2200 is known to typically be a face-on polymer, with the π-stacks of the large aromatic rings nucleating on the flat substrate during film casting and growing perpendicular to this substrate. This produces crystallites with the lamellar stacking (corresponding to the distance between polymers across the side chains) in the in-plane direction and the π -stacking in the out-of-pane direction. 54 Interestingly, while the N2200 polymer used in this study has a largely face- on population, it also has a small edge-on population (Fig. 9.4a). Side-chain influence on packing is easily observed in these two polymers. As monomers with oligoether side-chains are added to N2200, the population of edge-on polymer slightly increases, in agreement with literature observations (Fig. 9.4b). 55 325 Figure 9.4. 2D diffractograms of N2200 (left column) and N2200:OE (87:13) (right column) at the neutral, undoped state (top row), doped at the 2.4 V vs. Li/Li+, and 2.1 V vs. Li/Li + . For the pristine N2200 polymer, shown in Fig. 9.5a in black, the lamellar (100) peak, predominantly observed in the in-plane direction, is found at 0.25 Å −1 , corresponding to a d-spacing of 2.6 nm. The (020) π-stacking peak, predominantly found in the out-of-plane direction, is at 1.6 Å −1 , corresponding to a d- spacing of 0.39 nm. Pristine N2200-OE, shown in Fig. 9.5b in black, has similar d-spacings. The lamellar (100) peak is located at 0.27 Å −1 and the π-stacking (020) peak at 1.6 Å −1 , corresponding to a d-spacing 326 of 2.3 nm, and 0.39 nm, respectively. The smaller lamellar spacing for N2200-OE is due to the curling of the oligoether sidechains. 44, 56, 57 In the configuration that these polymers are used, they are n-type polymers, where the polymers are reduced to induce electronic conductivity. Thus, 2.3 V vs. Li + /Li corresponds to partial reduction and 2.3 V vs. Li + /Li to higher doping. When the polymers are reduced, an n-polaron, which carries a negative charge is delocalized along the polymer chain and Li + is the counter-ion. For both polymers, electrochemical doping does not change the orientation of the polymers, but it does change the crystallinity, and the extent of change of crystallinity is different. While the crystallinity of the N2200-OE polymer decreases somewhat upon doping, the overall polymer structure remains relatively unchanged (Fig 9.5b). Aside from the decrease in intensity of the principal (100) lamellar peak upon doping, all other peaks retain similar intensity and the peak positions are constant, regardless of the doping potential. This indicates that the small Li + cation is likely incorporated in the polar oligoether side-chains and does not significantly affect the overall backbone stacking. On the other hand, the crystallinity of N2200 (Fig. 9.5a) is more significantly reduced by electrochemical doping, particularly at lower potentials. This is likely because there is no low-energy place for the Li + to sit in the lamellar side chain region, and so the Li + associates more closely with the polar polymer backbone. 327 Figure 9.5. Integrated GIWAXS diffractograms. Panel (a) shows the full integration for N2200 plotted on a linear scale, and the inset shows the in-plane integration plotted on a log scale. Panel (b) shows the full integration of N2200-OE plotted on a linear scale, and the inset shows the in-plane integration plotted on a log scale. Interaction of the polymers with polysulfides. To examine the interaction between the conducting polymer binders and lithium polysulfides (LiPSs), UV-Vis absorption spectra of LiPS solutions were collected after exposing the solutions to polymer electrodes to allow for LiPS uptake. Fig. 9.6 shows the UV-Vis spectra of LiPS solutions after being exposed to PVDF, N2200, and N2200-OE over a period of three hours. No significant change in the absorption spectrum of the LiPS solution is observed when it is exposed to the PVDF electrode. However, a significant reduction in the absorption intensity is observed when LiPS solution is exposed to either N2200 or N2200-OE. This was also visually confirmed by taking photographs of the LiPS solutions (Fig. 328 9.7). Fig. 9.6d shows that LiPS absorption at 470 nm, which is dominated by absorption from Li 2S 3, 58 rapidly decreased when exposed to either N2200 and N2200-OE electrodes, although a slightly lower absorbance was observed for the LiPS solution exposed to N2200-OE. This suggested that both conducting polymers exhibit effective binding to the polysulfides. Lithium association with the OE side chains may be responsible for the slight enhanced binding in the N2200-OE. Figure 9.6. UV-Vis spectra of LiPS exposed to (a) PVDF, (b) N2200, and (c) N2200-OE. (d) Absorbance at 470 nm of LiPS solution over 3 hours of polymer exposure. 400 500 600 700 0 1 2 3 400 500 600 700 0 1 2 3 400 500 600 700 0 1 2 3 0 50 100 150 200 0.0 0.2 0.4 0.6 0.8 Absorbance (a.u.) Wavelength (nm) 0 minutes 60 minutes 120 minutes 180 minutes Absorbance (a.u.) Wavelength (nm) 0 minutes 60 minutes 120 minutes 180 minutes Absorbance (a.u.) Wavelength (nm) 0 minutes 60 minutes 120 minutes 180 minutes d b c Abs at 470 nm (a.u.) Time (minutes) PVDF N2200 N2200-OE a 329 Figure 9.7. Photographs of LiPS (a) before exposure to polymers and after 3 hours of exposure to (b) PVDF, (c) N2200, and (d) N2200-OE. Polysulfide Shuttle Current Measurement The polysulfide shuttle current or simply the shuttle current measurement quantifies the degree of polysulfide shuttling occurring between the electrodes in a Li-S cell using a methodology previously reported by our team. 59 Briefly, the measurement is carried after “cell formation” where the cell is discharged to a target state-of-charge (SOC), following which the cell is placed in open circuit and the voltage relaxes to a steady-state value. The cell voltage is then held constant at this steady-state value. The current that is required to maintain this cell voltage (potentiostatic conditions) is known as the “polysulfide shuttle current”. This current changes as a function of time, exhibiting a transient region at the beginning followed by a steady-state or plateau region (Fig. 9.8). 59 In the potential range where the soluble polysulfides are present, the observed steady-state current is attributed to the polysulfide shuttle. Li-S cells were assembled without additives to determine the baseline shuttle current. At approximately 2.3 V vs Li + /Li, the Li-S-N2200 and Li-S-N2200-OE cells exhibited four times lower shuttle current than the Li-S-PVDF cells (Fig. 9.9). Fig. 9.9b shows the observed shuttle current as function of potential in the polysulfide soluble potential region that goes from 2.4 to 2.15 V vs Li + /Li. Thus, the 330 conducting polymer (CP) cells had a notably lower shuttle current than those with PVDF, consistent with the interaction between the N2200 and N2200-OE and polysulfides observed by UV-Vis spectroscopy (Fig. 9.6). This reduction in shuttle current is observed in the absence of additives such as lithium nitrate that are often added to the electrolyte to reduce the effects of shuttling by passivating the Li electrode against LiPS reduction. To determine the sustained benefit of the polymer binder during cycling we measured the shuttle current in the 35 th cycle (Fig. 9.10). The shuttle current remained lower in the Li- S-N2200 cell compared to the Li-S-PVDF cell showing that the interaction between the CP polymers and LiPS was persistent even after continuous cycling. 59 Figure 9.8 (a) Potentiostatic intermittent shuttle current measurement curve of a Li-Sulfur-N2200 cell. (b) Voltage profile prior to and during shuttle current measurements on a Li-Sulfur-N2200 cell. 331 Figure 9.9. Polysulfide shuttle current measurement. (a) Observed current as a function of time at approximately 2.3 V vs Li + /Li after an OCV relaxation step and (b) steady-state current or shuttle current as a function of potential during discharge for sulfur-N2200 (denoted with green circles), sulfur-N2200- OE (denoted with pink triangles), and sulfur-PVDF electrodes (denoted with yellow squares). 332 Figure 9.10. Steady-state current or shuttle current as a function of potential after 35 cycles during discharge for a Li-Sulfur-N2200 (denoted with red circles) and Li-Sulfur-PVDF cells (denoted with black circles). Electrochemical Impedance Spectroscopy Electrochemical Impedance Spectroscopy (EIS) was measured as function of SOC at various points of the galvanostatic charge-discharge (GCD) curves for the Li-S-N2200 and Li-S-N2200-OE cells (Fig. 9.11a and 9.12). The real part of the impedance response (Z re) normalized to the electrode active mass loading (ohm mg −1 ) as a function of frequency was analyzed at 2.28, 2.15, and 2.1 V vs Li + /Li (Fig. 11b). For all SOCs, at high frequencies, Z re had the almost the same value indicating that the ohmic resistance of the electrolyte is similar for all cells. However, at mid and low frequencies, the cells with the conducting polymers (CPs) had approximately 4 to 5 times lower impedance than the PVDF cells. The reduction in impedance of the CP cells is attributed to the enhanced charge-transfer kinetics and lower mass transport resistance arising from the increased electronic and ionic conductivity of electrodes with conducting polymer binders that leads to improved utilization of the sulfur electrode. 333 Figure 9.11. (a) GCD profiles for Li-sulfur-N2200 and Li-sulfur-N2200-OE cells at C/20. Potentiostatic impedance response was measured at OCV as a function of SOC during charge and discharge (denoted with yellow circles). (b) Real part of the impedance response against frequency for the Li-sulfur-N2200, Li-sulfur-N2200-OE, and Li-Sulfur-PVDF cells at approximately 2.28 V, 2.15 V, and 2.1 V vs Li + /Li. 334 Figure 9.12. Three-dimensional chart of the impedance response as function of SOC. Nyquist plots as function of SOC for (a) sulfur-N2200, (b) sulfur-N2200-OE, and (c) sulfur-PVDF during charge. (d) sulfur- N2200, (e) sulfur-N2200-OE, and (f) sulfur-PVDF during discharge. Full Cell Electrochemical Cycling Rate capability of Li-S-N2200, Li-S-N2200-OE, and Li-S-PVDF cells was tested at moderate (1.79 mg cm −2 ) and high sulfur electrode loadings (3.4 mg cm −2 ). These tests were conducted following 24 hours of equilibration after assembly to allow adequate time for polymer swelling. Following equilibration, we carried out two “formation” cycles at C/15 rate before conducting the rate capability tests. At moderate loadings, the CP cells showed a higher rate capability than the PVDF cells (Fig. 9.13a and S10a-c) consistent with the results of EIS measurements (Fig. 7b), confirming the benefit of the CPs in reducing the impedance to electron and ion transport. At 2C rate, the Li-S-N2200-OE cells had a 335 remarkable specific capacity of more than 400 mAh g −1 attributed to the polymer’s enhanced electronic and ionic conductivity that leads to a higher accessible surface area for reaction and ion transport in the sulfur electrode. Even more significant, for the higher loading electrodes, the N2200-OE and N2200 cells showed even higher rate capability compared to the PVDF cells, indicating that the higher ionic and electronic conductivity of the CP binders is even more important in thicker electrodes, where Li + transport is hampered due to the higher electrode tortuosity and lower accessible surface area (Fig. 9.13b and 9.14d-f). 60 336 Figure 9.13. Rate Capability for Li-sulfur-n-dopable polymer cells under moderate and high loadings. Specific capacity as a function of cycle number with increments in the discharge rate and a constant charging rate at C/10 for sulfur electrode with (a) moderate loadings and (b) high loadings. (c) Specific capacity as a function of cycle number at a discharge/charge rate of C/2 for sulfur-N2200, sulfur-N2200- OE and Sulfur-PVDF. Impedance response as function of cycle number at approximately 2.2 V vs Li + /Li for the same (d) Li-sulfur-N2200, (e) Li-sulfur-N2200-OE, and (f) Li-sulfur-PVDF cells for the 15 th , 200 th , and 495 th cycles. GCD cycling and EIS as a function of cycle number was measured to test the effectiveness of the CPs in enhancing the cycle-life of the cells. After 500 cycles, the Li-S-N2200 and Li-S-N2200-OE cells had an 337 impressive capacity retention of approximately 67 and 82%, respectively (Fig. 9.13c and 9.15). In contrast, the Li-S-PVDF cell showed significant capacity decay in the first 100 cycles with a capacity retention of only 18% after 200 cycles. Fig. 9.13d-f shows the Nyquist plots at the 15 th , 200 th , and 495 th cycles. Compared to the CP cells after 200 th cycles the PVDF cells showed 5 and 4 times higher resistance at mid-low and high frequencies respectively, indicating an increase in the faradaic, mass-transport, and ohmic resistance after cycling, attributed to the effects of polysulfide shuttling in the cell. Thus, n- dopable CPs can lower the cell’s overall impedance and reduce the polysulfide shuttling to extend the battery cycle-life. Furthermore, in comparing the CP containing cells after the 495 th cycle, the N2200-OE cell had approximately 3 times lower impedance than the N2200 cell, suggesting higher effectiveness of the N2200-OE binder (Fig. 9.13 d,e). We expect the mechanical elasticity induced by the increased degree of swelling in the N2200-OE binder to benefit the expansion and contraction of the sulfur electrode. Thus, the shuttle current, impedance values as function of SOC and cycle number, the LiPS binding evidenced in changes in UV-Vis absorption, and the swelling measurements provide the basis for interpreting the observed beneficial effects of the CPs on rate capability and cycle life. 338 Figure 9.14 Corresponding galvanostatic charge-discharge curves from Figure 9.13a; (a) Li-Sulfur-N2200, (b) Li-Suflur-N2200-OE, and (c) Li-Sulfur-PVDF. Corresponding galvanostatic charge-discharge curves from Figure 9.13b; (d) Li-Sulfur-N2200, (e) Li-Suflur-N2200-OE, and (f) Li-Sulfur-PVDF. Figure 9.15. Corresponding galvanostatic charge-discharge curves from Figure 9.13c; (a) Li-Sulfur-N2200 and (b) Li-Suflur-N2200-OE. 339 Conclusions We have demonstrated that n-dopable conducting polymers N2200 and N2200-OE when used as binders at the sulfur electrode can significantly enhance the capacity output and cycle life of lithium- sulfur batteries compared to standard insulating binders such as PVDF. The benefit of these conducting polymer binders resulted from the electronic and ionic conductivity of the polymers over the range of potential relevant to the operation of the sulfur electrode. Furthermore, the strong affinity of the polysulfides to the functional groups on the polymer curtailed the degree of polysulfide shuttling in the electrolyte leading to markedly enhanced capacity retention as high as 82% after 500 cycles. The modification of the N2200 polymer with oligoether side chains increases the swelling of the polymer and results in enhanced ionic conductivity. The electronic and ionic conductivity of the conducting polymer binders led to a significant reduction in impedance of the lithium-sulfur cell compared to the baseline PVDF cells, and manifested as a notable increase in rate capability. These benefits are observed both at moderate and high loadings of the sulfur electrode. These results emphasize the advantages of using conducting polymer binders that can support electron and ion transport in the potential range of electrode operation, and the presence of functional groups in the polymer that interact strongly with the soluble polysulfides in addressing the limitations of lithium-sulfur cells. 340 References 1. A. Manthiram, Y. Fu, S.-H. Chung, C. Zu, and Y.-S. Su, Chem. Rev., 114, 11751 (2014). 2. M. Winter and J. O. Besenhard, Handbook of battery materials, 383 (1998). 3. G. Li, Z. Chen, and J. Lu, Chem, 4, 3 (2018). 4. C. Yang, P. Li, J. Yu, L.-D. Zhao, and L. Kong, Energy, 201, 117718 (2020). 5. X. Zeng, M. Li, D. Abd El ‐Hady, W. Alshitari, A. S. Al ‐Bogami, J. Lu, and K. Amine, Adv. Energy Mater., 9, 1900161 (2019). 6. Z. Han, S. Li, Y. Wu, C. Yu, S. Cheng, and J. Xie, J. Mater. Chem. A, 9, 24215 (2021). 7. A. Manthiram, Y. Fu, and Y.-S. Su, Acc. Chem. Res., 46, 1125 (2013). 8. K. Zhu, C. Wang, Z. Chi, F. Ke, Y. Yang, A. Wang, W. Wang, and L. Miao, Frontiers in Energy Research, 123 (2019). 9. C. Deng, Z. Wang, S. Wang, and J. Yu, J. Mater. Chem. A, 7, 12381 (2019). 10. X. Fan, W. Sun, F. Meng, A. Xing, and J. Liu, Green Energy & Environment, 3, 2 (2018). 11. J. Guo, H. Pei, Y. Dou, S. Zhao, G. Shao, and J. Liu, Adv. Funct. Mater., 31, 2010499 (2021). 12. Y. He, Z. Chang, S. Wu, and H. Zhou, J. Mater. Chem. A, 6, 6155 (2018). 13. Y. Hu, W. Chen, T. Lei, Y. Jiao, J. Huang, A. Hu, C. Gong, C. Yan, X. Wang, and J. Xiong, Adv. Energy Mater., 10, 2000082 (2020). 14. L. Borchardt, M. Oschatz, and S. Kaskel, Chemistry–A European Journal, 22, 7324 (2016). 15. N. Jayaprakash, J. Shen, S. S. Moganty, A. Corona, and L. A. Archer, Angew. Chem. Int. Ed., 50, 5904 (2011). 16. Z.-L. Xu, J.-K. Kim, and K. Kang, Nano Today, 19, 84 (2018). 17. C. Zhang, H. B. Wu, C. Yuan, Z. Guo, and X. W. Lou, Angew. Chem. Int. Ed., 51, 9592 (2012). 18. J.-j. Chen, Q. Zhang, Y.-n. Shi, L.-l. Qin, Y. Cao, M.-s. Zheng, and Q.-f. Dong, Phys. Chem. Chem. Phys., 14, 5376 (2012). 19. J. Guo and J. Liu, Nanoscale Advances, 1, 2104 (2019). 20. G. Liu, D. Luo, R. Gao, Y. Hu, A. Yu, and Z. Chen, Small, 16, 2001089 (2020). 21. L. Luo and A. Manthiram, ACS Energy Lett., 2, 2205 (2017). 22. Y. Zhao, Y. Ye, F. Wu, Y. Li, L. Li, and R. Chen, Adv. Mater., 31, 1806532 (2019). 23. W. Fan, L. Zhang, and T. Liu, Materials Chemistry Frontiers, 2, 235 (2018). 24. Y. Xiang, J. Li, J. Lei, D. Liu, Z. Xie, D. Qu, K. Li, T. Deng, and H. Tang, ChemSusChem, 9, 3023 (2016). 25. T. Z. Zhuang, J. Q. Huang, H. J. Peng, L. Y. He, X. B. Cheng, C. M. Chen, and Q. Zhang, Small, 12, 381 (2016). 26. D. Moy and S. Narayanan, J. Electrochem. Soc., 164, A560 (2017). 341 27. X. Ji, K. T. Lee, and L. F. Nazar, Nat. Mater., 8, 500 (2009). 28. C. Liang, N. J. Dudney, and J. Y. Howe, Chem. Mater., 21, 4724 (2009). 29. K. Liu, H. Zhao, D. Ye, and J. Zhang, Chem. Eng. J., 417, 129309 (2021). 30. H. Xiang, N. Deng, H. Zhao, X. Wang, L. Wei, M. Wang, B. Cheng, and W. Kang, Journal of Energy Chemistry, 58, 523 (2021). 31. Z. W. Seh, Q. Zhang, W. Li, G. Zheng, H. Yao, and Y. Cui, Chemical Science, 4, 3673 (2013). 32. S. Tu, X. Chen, X. Zhao, M. Cheng, P. Xiong, Y. He, Q. Zhang, and Y. Xu, Adv. Mater., 30, 1804581 (2018). 33. H. Yi, T. Lan, Y. Yang, Z. Lei, H. Zeng, T. Tang, C. Wang, and Y. Deng, J. Mater. Chem. A, 6, 18660 (2018). 34. H. Chen, W. Dong, J. Ge, C. Wang, X. Wu, W. Lu, and L. Chen, Sci. Rep., 3, 1 (2013). 35. M. Kazazi, Ionics, 22, 1103 (2016). 36. M. Kazazi, M. Vaezi, and A. Kazemzadeh, Ionics, 20, 635 (2014). 37. L. Xiao, Y. Cao, J. Xiao, B. Schwenzer, M. H. Engelhard, L. V. Saraf, Z. Nie, G. J. Exarhos, and J. Liu, Adv. Mater., 24, 1176 (2012). 38. L. Xiao, Y. Cao, J. Xiao, B. Schwenzer, M. H. Engelhard, L. V. Saraf, Z. Nie, G. J. Exarhos, and J. Liu, J. Mater. Chem. A, 1, 9517 (2013). 39. J. Yan, B. Li, and X. Liu, Nano Energy, 18, 245 (2015). 40. Y. Yang, G. Yu, J. J. Cha, H. Wu, M. Vosgueritchian, Y. Yao, Z. Bao, and Y. Cui, ACS Nano, 5, 9187 (2011). 41. T.-G. Jeong, Y.-S. Lee, B. W. Cho, Y.-T. Kim, H.-G. Jung, and K. Y. Chung, J. Alloys Compd., 742, 868 (2018). 42. P. Das, B. Zayat, Q. Wei, C. Z. Salamat, I.-B. Magdău, R. Elizalde-Segovia, D. Rawlings, D. Lee, G. Pace, A. Irshad, L. Ye, A. Schmitt, R. A. Segalman, T. F. Miller, S. H. Tolbert, B. S. Dunn, S. R. Narayan, and B. C. Thompson, Chem. Mater., 32, 9176 (2020). 43. R. Elizalde-Segovia, P. Das, B. Zayat, A. Irshad, B. C. Thompson, and S. R. Narayanan, J. Electrochem. Soc., 168, 110541 (2021). 44. P. Das, R. Elizalde-Segovia, B. Zayat, C. Z. Salamat, G. Pace, K. Zhai, R. C. Vincent, B. S. Dunn, R. A. Segalman, S. H. Tolbert, S. R. Narayan, and B. C. Thompson, Chem. Mater., (2022). 45. S. N. Patel, A. E. Javier, and N. P. Balsara, ACS Nano, 7, 6056 (2013). 46. C. H. Lai, D. S. Ashby, T. C. Lin, J. Lau, A. Dawson, S. H. Tolbert, and B. S. Dunn, Chem. Mater., 30, 2589 (2018). 47. M. E. Abdelhamid, A. P. O'Mullane, and G. A. Snook, Rsc Advances, 5, 11611 (2015). 48. L. M. Santino, Y. Lu, Y. Diao, H. Wang, and J. M. D’Arcy, Conjugated Polymers, 561 (2019). 49. G. A. Snook, P. Kao, and A. S. Best, J. Power Sources, 196, 1 (2011). 50. D. Trefz, A. Ruff, R. Tkachov, M. Wieland, M. Goll, A. Kiriy, and S. Ludwigs, J. Phys. Chem. C, 119, 22760 (2015). 342 51. B. A. Wavhal, M. Ghosh, S. Sharma, S. Kurungot, and S. Asha, Nanoscale, 13, 12314 (2021). 52. B. Zayat, P. Das, B. C. Thompson, and S. R. Narayan, J. Phys. Chem. C, 125, 7533 (2021). 53. S. D. Oosterhout, V. Savikhin, J. Zhang, Y. Zhang, M. A. Burgers, S. R. Marder, G. C. Bazan, and M. F. Toney, Chem. Mater., 29, 3062 (2017). 54. D. Simatos, L. Spalek, U. Kraft, M. Nikolka, X. Jiao, C. McNeill, D. Venkateshvaran, and H. Sirringhaus, APL Materials, 9, 041113 (2021). 55. J. Liu, L. Qiu, R. Alessandri, X. Qiu, G. Portale, J. J. Dong, W. Talsma, G. Ye, A. A. Sengrian, P. C. T. Souza, M. A. Loi, R. C. Chiechi, S. J. Marrink, J. C. Hummelen, and L. J. A. Koster, Adv. Mater., 30, 1704630 (2018). 56. S. I. Lall-Ramnarine, M. Zhao, C. Rodriguez, R. Fernandez, N. Zmich, E. D. Fernandez, S. B. Dhiman, E. W. Castner, and J. F. Wishart, J. Electrochem. Soc., 164, H5247 (2017). 57. I. Sahalianov, J. Hynynen, S. Barlow, S. R. Marder, C. Mü ller, and I. Zozoulenko, J. Phys. Chem. B, 124, 11280 (2020). 58. N. Xu, T. Qian, X. Liu, J. Liu, Y. Chen, and C. Yan, Nano Lett., 17, 538 (2017). 59. D. Moy, A. Manivannan, and S. Narayanan, J. Electrochem. Soc., 162, A1 (2014). 60. J. Wu, Z. Rao, X. Liu, Y. Shen, C. Fang, L. Yuan, Z. Li, W. Zhang, X. Xie, and Y. Huang, Adv. Mater., 33, 2007428 (2021). 343 References Each chapter has its own list of references at the end of it. Refer to the corresponding chapter. 344 Appendix - Full list of Figures Figure 1.1. Schematic illustration of a LiCoO2||graphite lithium-ion battery. Figure 1.2. Scheme illustration of the strategies that can be pursued to increases the energy and power density at the active material level, electrode level, and cell level. Figure 1.3. Evaluation flowchart for building a high-energy and high-power lithium battery. Figure 1.4 Specific capacity and energy density comparison among the of Li-ion and Li-S active materials and resulting cell. Figure 1.5 Schematic representation of a Li-S battery showing a simplification of the electrochemical processes that occur during the cell’s discharge/charge. Figure 1.6 Charge and discharge profiles for a Li-S battery. 4 regions are present; I) correspond to the dissolution of elemental sulfur into the high order polysulfides, II) soluble polysulfides reduction, III) reduction of the soluble polysulfides to the non-soluble polysulfides, and IV) Li2S2 to Li2S. Figure 1.7. Schematic illustration of the polysulfide shuttling occurring in the Lithium-Sulfur Figure 1.8 Challenges and electrolyte classification of lithium-sulfur batteries. Scheme 2.1. Synthesis of Monomers (S3 and S4). Scheme 2.2. Polymerization of PProDOT-Hx2 using DArP Figure 2.1 GPC analysis in Chloroform with calibration to a polystyrene standard: Mn = 19,182 g mol-1; Mw =31,820 g mol-1; PDI = 1.6 Figure 2.2. (a) SEM image of the spin-coated PProDOT-Hx 2 on interdigitated gold microelectrodes. The light color shows the gold electrodes and the dark color shows the polymer imbedded between the electrodes. (b) The 2- electrode configuration used to measure the electronic conductivity of the polymer as a function of voltage. (c) Cyclic voltammetry plot of PProDOT-Hx 2 at various scan rates. (d) The 3-electrode configuration used to electrochemically dope the polymer and determine its ionic conductivity. Figure 2.3 The normalized, integrated EPR signal intensity for a series of DPPH calibration samples with known spin concentrations. Figure 2.4 Synthesis of Monomer (S5) and Polymers P3HT (P1), PProDOT-Hx 2 (P2) and (Hex:OE)(80:20) PProDOT Random Copolymer (P3). Figure 2.5. 1 H-NMR of S5 in CDCl 3 at 25 °C and 500 MHz. Figure 2.6. 1 H-NMR of P1 (P3HT) in CDCl 3 at 25 °C and 500 MHz. Figure 2.7. 1 H-NMR of P2 (PProDOT-Hx 2) in CDCl 3 at 25 °C and 500 MHz Figure 2.8. 1 H-NMR of P3 ((Hex:OE) (80:20) PProDOT Random Copolymer) in CDCl 3 at 25 °C and 500 MHz. Scheme 2.5 Synthesis of (Hex:OE) PProDOT Random Copolymers and PProDOT-OE 2. Figure 2.9 1 H-NMR of B in CDCl 3 at 25 °C and 500 MHz. Figure 2.10 1 H-NMR of A in CDCl 3 at 25 °C and 500 MHz. Figure 2.11 1 H-NMR of C in CDCl 3 at 25 °C and 500 MHz. Figure 2.12 1 H-NMR of D in CDCl 3 at 25 °C and 500 MHz. Figure 2.13 1 H-NMR of PProDOT-Hx 2 in CDCl 3 at 25 °C and 500 MHz. 345 Figure 2.14 1 H-NMR of (Hex:OE) (95:5) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 4.7 % as observed from NMR Figure 2.15 1 H-NMR of (Hex:OE) (85:15) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 14.7 % as observed from NMR Figure 2.16 1 H-NMR of (Hex:OE) (75:25) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 24.6 % as observed from NMR Figure 2.17 1 H-NMR of (Hex:OE) (65:35) PProDOT in CDCl 3 at 25 °C and 500 MHz. Oligoether (OE) percentage = 32.6 % as observed from NMR Figure 2.18 1 H-NMR of (Hex:OE) (50:50) PProDOT Alternating Copolymer in CDCl 3 at 25 °C and 500 MHz. Figure 2.19 1 H-NMR of (Hex:OE) (25:75) PProDOT Random Copolymer in CDCl 3 at 25 °C and 500 MHz.Oligoether (OE) percentage = 72.5 % as observed from NMR Figure 2.20 1 H-NMR of (Hex:OE) (15:85) PProDOT Random Copolymer in CDCl 3 at 25 °C and 500 MHz.Oligoether (OE) percentage = 80 % as observed from NMR Figure 2.21 1 H-NMR of PProDOT-OE 2 Random Copolymer in CDCl 3 at 25 °C and 500 MHz. Electrochemical Characterization and Cell Fabrication Figure 2.22 (a) A representative cyclic voltammogram for (75:25) PProDOT, acquired at 1 mV s-1 in 1 M LiPF 6 in EC:DMC (1:1 v/v). (b) Polaron character of the (75:25) PProDOT measured the same day as electrochemical doping (red squares) and after 1 week (blue circles). Figure 2.23 Swelling set-up with glass vial and propylene carbonate inside. Figure 2.24 GPC trace of PProDOT-Hx 2: M n = 17.4 kDa, Ð = 1.76 Figure 2.25 GPC trace of (95:5) PProDOT: M n = 17.6 kDa, Ð = 1.74 Figure 2.26 GPC trace of (85:15) PProDOT: M n = 13.9 kDa, Ð = 1.81 Figure 2.27 GPC trace of (75:25) PProDOT: M n = 16.3 kDa, Ð = 1.78 Figure 2.28 GPC trace of (65:35) PProDOT: M n = 13.7 kDa, Ð = 1.49 Figure 2.29. 1H-NMR of S1 in CDCl3 at 25 °C and 400 MHz. Figure 2.30. 1H-NMR of S2 in CDCl3 at 25 °C and 400 MHz. Figure 2.31. 1H-NMR of S4 in CDCl3 at 25 °C and 400 MHz. Figure 2.32. 1H-NMR of PNDI (OD:OE) or N2200-OE in CDCl3 at 25 °C and 400 MHz. Oligoether (OE) percentage = 20 % (Feed Ratio) Figure 3.1. (a) CV data for PProDOT-Hx 2 thin film in the potential range of 3-4.2 V at a scan rate of 10 mV s -1 for cycles 1-4. (b) CV data for PProDOT-Hx 2 thin film at various potential intervals, at a scan rate of 10 mV s -1 . (c) CV data for PProDOT-Hx 2 thin film as a function of various scan rates from 10 to 100 mV s -1 . (d) Log of the peak current (i) vs. log of the scan rate (v) for the data shown in Figure 3.1c. Figure 3.2 shows the CV curves of PProDOT-Hx 2 thin film in 3-4.2 V at 50 mV s -1 for 100 cycles. The redox peaks are still very evident, thus verifying the redox stability of the polymer. However, the capacity does decrease on cycling; the capacity retention is 60% after 100 cycles. Figure 3.3. (a) An exemplar cyclic voltammogram (CV) of PProDOT-Hx 2 thin film electrode in 1 M LiPF 6 in EC:DMC (50:50 v/v). Total number of charges of each sample were determined by CV of each sample. (b) Ratio of the number 346 of spins to the number of charges in the PProDOT-Hx 2 film as a function of the electrochemical potential determined from the normalized, integrated EPR signal intensity. Figure 3.4. (a) SEM image of the spin-coated PProDOT-Hx 2 on interdigitated gold microelectrodes. The light color shows the gold electrodes and the dark color shows the polymer imbedded between the electrodes. (b) The 2- electrode configuration used to measure the electronic conductivity of the polymer as a function of voltage. (c) Cyclic voltammetry plot of PProDOT-Hx 2 at various scan rates. (d) The 3-electrode configuration used to electrochemically dope the polymer and determine its ionic conductivity. Figure 3.5. (a) A representative Nyquist impedance plot obtained at 4V vs Li/Li + using the 2-electrode configuration and the fitting circuit (inset) used to obtain the electronic conductivity. (b) A representative Nyquist impedance plot obtained at 4V vs Li/Li + using the 3-electrode configuration to obtain the ionic conductivity. (c) The electronic and (d) ionic conductivities of P3HT and PProDOT-Hx 2. Figure 3.6. Representative 2-D diffractograms of PProDOT-Hx 2. The missing wedge of data in each image results from the grazing incidence geometry used in the experiment. Figure 3.7. Full integration of GIWAXS diffractograms for PProDOT-Hx 2 doped under different conditions. Part (a) shows the polymer doped at oxidation 1 using two common battery electrolytes, LiTFSI and LiClO 4, which provide TFSI - and ClO 4 - anions to the polymer film upon doping. Part (b) compares the structure of the polymer doped at oxidation peaks 1 and 2 using the most stable of our battery electrolyte, LiTFSI. Figure 3.8 TEM image of (a) 96-4% NCA-PProDOT-Hx 2 and (b) 90-3-3-4% NCA-SP-CNT-PProDOT-Hx 2 electrodes. Figure 3.9. Electrochemical properties for NCA-PProDOT-Hx2 and NCA-PVDF electrodes: (a) Rate capability of the NCA-PProDOT-Hx2 and NCA-PVDF. The corresponding galvanostatic charge-discharge (GCD) curves of the NCA- PProDOT-Hx2 (b) and NCA-PVDF (c) at various rates. (d) Long-term cycling for NCA-PProDOT-Hx2 and NCA-PVDF at a rate of 2C. The corresponding GCD curves of the NCA-PProDOT-Hx 2 (e) and NCA-PVDF (f) at different cycles. Figure 3.10 Rate capability comparison of the NCA-PVDF and NCA-CNT-Super P (No Binder) with an active material mass loading of 6 mg cm -2 (a) and 11 mg cm -2 (b). Figure 3.11 Rate capability comparison of the NCA-PProDOT-Hx 2, NCA-PVDF and NCA-CNT-SuperP (No Binder) with an active material mass loading of 11 mg cm -2 . Figure 3.12 Galvanostatic charge-discharge (GCD) curves at various rates of the NCA-CNT-Super P (No Binder) with an active material mass loading of 6 mg cm -2 (a) and 11 mg cm -2 (b). Figure 3.13 Galvanostatic charge-discharge (GCD) curves of the NCA-PProDOT-Hx 2 (a) and NCA-PVDF (b) at various rates with an active material mass loading of 11 mg cm -2 . Figure 3.14: The initial cycle charge and discharge profiles for NCA-PProDOT-Hx 2 and NCA-PVDF electrodes at a rate of C/5. Figure 3.15 (a) Comparison of the charge and discharge curve of the NCA-PProDOT-Hx 2 (a) and NCA-PVDF (b) at at C/10, EIS was measured as function of SOC (denoted with green circles). Figure 3.16 A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of charge for the NCA-PProDOT-Hx 2 (a) and NCA-PVDF (b) cell with a mass loading of 11 mg cm -2 . Figure 3.17 Real part of the impedance response agaisnt frequency (at low freqnuencies) as fucntion of SOC for the NCA-PProDOT-Hx 2 (black) and NCA-PVDF (red) cells. By maintaining the same electrode components ratios, we also made the comparison of PProDOT-Hx 2 and P3HT as electrode additives against commercial PVDF binder to confirm their comparable and superior performance at higher rates (Figure 3.18). Figure 3.18 Rate capability comparison of the NCA-PProDOT-Hx 2, NCA-PVDF, and NCA-P3HT with an active material mass loading of ~ 6 mg cm -2 . 347 Figure 4.1 CVs of (Hex:OE)(80:20) a) at various potential intervals at 50 mV s -1 , b) at various scan rates, and c) over 100 cycles at 10 mV s -1 . d) Electronic and Ionic Conductivities of P3HT, PProDOT-Hx 2 and (Hex:OE)(80:20). Figure 4.2 Specific energy at the component level of Li-NCA-π-conjugated polymer cells under mass-efficient and non- limited conditions. a) Mass distribution of cell components, b) Specific energy at the cell-level as a function of cycle number for Li-NCA-PProDOT-Hx 2, Li-NCA-(Hex:OE)(80:20), Li-NCA-P3HT, and Li-NCA-PVDF cells with an areal loading of 14 mg cm -2 under mass-efficient and non-limited conditions discharged at C/2 and charged at C/5. NCA electrodes were fabricated with a minimal amount of carbon additive (1 wt%). Two formation cycles at C/10 were conducted before cycling at 23 °C. Non-limited cells had an excess of lithium and electrolyte with an approximate N/P ratio of 45 and an E/C ratio of 17.2. In contrast, mass-efficient cells were assembled with a minimum amount of metallic lithium and electrolyte with an approximate N/P ratio of 3.3 (≈ 3 mg of Li/cell) and E/C ratio of 3.1 (≈ 9-13 µL/cell). Conventional carbonate/LiPF 6 electrolyte was utilized. c)-e) Corresponding galvanostatic charge-discharge curves for the mass-efficient cells in b). Figure 4.3. Galvanostatic charge-discharge curves as a function of cycling for the Li-NCA-PVDF cell under mass- efficient conditions charged at C/5 and discharged at C/2. Two formation cycles were carried at C/10 before cycling at 25 °C. Figure 4.4 Corresponding capacity as function of cycle number for the mass-efficient cells in Figure 1a. Specific capacity, and coulombic efficiency as function of cycle number for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li-NCA-P3HT, and d) Li-NCA-PVDF cells under mass-efficient conditions with an areal loading of approximately 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. Figure 4.5 Corresponding galvanostatic charge-discharge curves for the non-limited cells in Figure 1a. Specific energy at the cell level for a) Li-NCA-PProDOT-Hx 2 cell, b) Li-NCA-(Hex:OE)(80:20) cell, c) Li-NCA-P3HT cell, and d) Li-NCA- PVDF cell under non-limited conditions with an areal loading of 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. Figure 4.6 Corresponding capacity as function of cycle number for the non-limited cells in Figure 1a. Specific capacity and coulombic efficiency as function of cycle number for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li- NCA-P3HT, and d) Li-NCA-PVDF cells under non-limited conditions with an areal loading of approximately 14 mg cm -2 . Two formation cycles were carried at C/10 before cycling at 25 °C. Furthermore, GCD cycling of solely the CPs (without NCA) under lean electrolyte conditions over 100 cycles in the potential range of 2.7 V to 4.2 V vs Li + /Li exhibited a capacity retention of 98.3, 99.8, and 99.5% for PProDOT-Hx 2, (Hex:OE)(80:20), and P3HT respectively, confirming the CPs excellent cyclability and stability (Figure 4.7). Figure 4.7 Galvanostatic characterization of the conducting polymers over the potential window of 2.7 V to 4.2 V vs Li + /Li. a) Percentage capacity retention and coulombic efficiency as a function of cycle number for PProDOT-Hx 2, (Hex:OE)(80:20), and P3HT in a half-cell configuration with a conducting polymer loading of approximately 0.65 mg cm -2 under limited electrolyte conditions with an approximate electrolyte to polymer ratio of 9 µL per mg of polymer, discharged and charged at 0.26 mA cm 2 b)-c) corresponding galvanostatic charge-discharge curves for the cells in a. To prepare the conducting polymer electrodes a slurry composed of conducting polymer and carbon (Super P) in a weight ratio of 40:60 was prepared in OCDB. The slurry was coated onto Al foil using the doctor blading, followed by vacuum drying at 110 °C for two hours. The electrodes were punched into 1.539 cm 2 discs. Figure 4.8 Impedance characterization for the cells. Galvanostatic charge-discharge profiles for the a) Li-NCA- PProDOT-Hx 2 and b) Li-NCA-PVDF “mass-efficient” cells at C/20 with a charging cut-off of 4.2 V vs Li + /Li. EIS was measured as function of SOC at OCV (denoted with yellow circles). c) equivalent electrical circuit used for fitting the impedance response. Corresponding three-dimensional Nyquist plots of the experimental (denoted with color spheres) and fitted (denoted with black lines) impedance response as function of SOC during charge for the d) Li-NCA- PProDOT-Hx 2 cell and e) Li-NCA-PVDF cells in a and b. Figure 4.9. Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-PProDOT-Hx 2 cell under mass-efficient 348 conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.10Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-PProDOT-Hx 2 cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.11 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-(Hex:OE)(80:20) cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.12 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-(Hex:OE)(80:20) cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.13 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-P3HT cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.14 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-P3HT cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.15 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-PVDF cell under mass-efficient conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.16 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. a) Nyquist plots as function of SOC during charge c) discharge for the Li-NCA-PVDF cell under non-limited conditions for the first cycle, and b) for charge d) discharge after 26 cycles. With an approximate areal loading of 14 mg cm 2 . Figure 4.17 Impedance characterization for the Li-NCA-π-conjugated polymer cells under mass-efficient and non- limited conditions. Evolution of the a) electrolyte resistance R 1, b) SEI resistance R 2, c) charge transfer resistance R 3, d) SEI capacitance C 1 e), double layer capacitance C 2, and f) Warburg coefficient S as a function of SOC for the 1 st cycle and the 26 th cycle during charge and discharge for Li-NCA-PProDOT-Hx 2 (denoted by blue and turquoise connected dots), Li-NCA-(Hex:OE)(80:20) (olive and green), Li-NCA-P3HT (purple and magenta), and Li-NCA-PVDF (red and pink). Relative fitting errors are indicated with error bars. The equivalent electrical circuit utilized for the impedance fitting is shown in Figure 4.8c. Figure 4.18. SEM characterization of the NCA-π-conjugated polymer electrodes. SEM of the electrode composite films as prepared with a composition of 95 wt% NCA, 4 wt% conducting polymer and 1 wt% of carbon at various scales ranges from 100 microns to 500 nm for a) NCA-PProDOT-Hx 2, b) NCA-(Hex:OE)(80:20), c) NCA-P3HT and d) NCA-PVDF. Figure 4.19 Macro pores characterization of the NCA-PProDOT-Hx 2 electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. Figure 4.20 Meso pores characterization of the NCA-PProDOT-Hx 2 electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. 349 Figure 4.21 Macro pores characterization of the NCA-(Hex:OE)(80:20) electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. Figure 4.22 Meso pores characterization of the NCA-(Hex:OE)(80:20) electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. Figure 4.23 Macro pores characterization of the NCA-P3HT electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores, and c) bare outlines form the particle analysis. Figure 4.24 Meso pores characterization of the NCA-P3HT electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores, and c) bare outlines form the particle analysis. Figure 4.25 Macro pores characterization of the NCA-PVDF electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the macro pores and c) bare outlines form the particle analysis. Figure 4.26 Meso pores characterization of the NCA-PVDF electrodes. a) Original scanning electron microscopy image before the analysis, b) after the applied threshold for distinguishing the meso pores and c) bare outlines form the particle analysis. Figure 4.27 EDS Elemental mapping of NCA-PProDOT-Hx 2 electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% PProDOT-Hx 2, and 1% Carbon, b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. Figure 4.28 EDS spectrum of NCA-PProDOT-Hx 2 electrode surface. Figure 4.29 EDS Elemental mapping of NCA-(Hex:OE)(80:20) electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% (Hex:OE)(80:20), and 1% Carbon b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. Figure 4.30 EDS spectrum of NCA-(Hex:OE)(80:20) electrode surface. Figure 4.31 EDS Elemental mapping of NCA-P3HT electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% P3HT, and 1% Carbon, b) all combined elements mapping, c) carbon mapping, d) sulfur mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. Figure 4.32 EDS spectrum of NCA-P3HT electrode surface Figure 4.33 EDS Elemental mapping of NCA-PVDF electrodes. a) SEM image of the electrode, contains 95 wt% NCA, 4% PVDF, and 1% Carbon b) all combined elements mapping, c) carbon mapping, d) fluorine mapping, e) cobalt mapping, f) nickel mapping, g) oxygen mapping, h) aluminium mapping. Figure 4.34 EDS spectrum of the NCA-PVDF electrode surface. Figure 4.35 Pore parameters of the NCA-π-conjugated polymers electrodes from SEM. a) Estimation of the surface % area covered by macro pores and b) meso pores for the NCA-π-conjugated polymers electrodes. c) Mean Feret Diameter of the macro pores and d) meso pores NCA-π-conjugated polymers electrodes. Figure 4.36 Three-dimensional chart of the impedance response as function of state of charge at open circuit voltage. Nyquist plots as function of SOC during charge for a) Li-NMC-622-PVDF cell b) for the Li-MNC-622-PProDOT-Hx 2 cell under mass-efficient conditions with an approximate areal loading of 13 mg cm 2 . Figure 4.37. Charge transfer resistance as function of SOC obtained by ECC fitting from Figure 4.10c for NCA and NMC-622 based electrodes with PVDF and PProDOT-Hx 2. 350 Figure 4.38 a) Galvanostatic charge-discharge curves as function of C-rate for a Li-NMC-622-PProDOT-Hx 2 under non- limited conditions with an approximate mass loading of 5.5 mg cm -2 . Two formation cycles where carried at C/20. C- rate was calculated based on the reversible capacity of NMC-622 being 1C = 180 mA g -1 . b), corresponding specific capacity at the cell-level as function of cycle number at the different rates from the discharge curves plotted in a). Figure 4.39 Differential capacity vs voltage analysis (dQ dV -1 vs V) for Li-NCA-π-conjugated polymer cells. dQ dV -1 vs V curves measured as a function of number of cycles with a C/2 rate for discharge and C/5 for charging including the first formation cycle at C/10 (denoted with a red curve) under mass-efficient and non-limited conditions between 2.7 and 4.2 V vs Li + /Li for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20) cells, c) Li-NCA-P3HT cells, and d) Li-NCA- PVDF cells. Figure 4.40 Rate Capability for Li-NCA-π-conjugated polymer cells under mass-efficient and non-limited conditions. a) Specific energy at the cell-level as a function of cycling with increments in the discharge rate with a constant charging rate (C/5) and two formation cycles charged/discharged at C/10 for the Li-NCA-π-conjugated polymer cells under mass-efficient conditions. Mass-efficient cells had an approximate N/P ratio of 3 and an E/C ratio of 3. b) Specific capacity as a function of cycling at various discharge rates with a constant charging rate and two formation cycles for the Li-NCA-π-conjugated polymer cells under non-limited conditions. Corresponding galvanostatic discharge-charge curves in a) for c) Li-NCA-PProDOT-Hx 2, d) Li-NCA-(Hex:OE)(80:20), e) Li-NCA-P3HT, and f) Li-NCA-PVDF cells Figure 4.41 Rate Capability for Li-NCA-π-conjugated polymer cells under non-limited conditions. Corresponding galvanostatic discharge-charge curves from Figure 6b for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(Hex:OE)(80:20), c) Li- NCA-P3HT, and d) Li-NCA-PVDF cells in non-limited conditions. Scheme 5.1. Synthesis of (Hex:OE) PProDOT Random Copolymers using DArP Figure 5.1. Electrochemical performance of (85:15) PProDOT. (a) Cycling Voltammograms at 10 mV s-1 in various potential windows up to 4.5 V, (b) CV scans at different scan rates, and (c) long-term cycling at 10 mV s-1. Figure 5.2. CV data at 10 mV s -1 for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT- Hx 2 at various potential intervals. Figure 5.3 CV scan at varying scan rates between 100 and 20 mV s -1 for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT-Hx 2. Figure 5.4 Long-term CV curves for (a) (65:35) PProDOT, (b) (75:25) PProDOT, (c) (95:05) PProDOT, and (d) PProDOT- Hx 2 thin films at 10 mV s -1 . Figure 5.5 (a) Capacity retention of (Hex:OE) PProDOT copolymer family after 100 cycles at 10 mV s -1 as a function of oligo ether content. Figure 5.6 Representative cyclic voltammograms for each of the PProDOT variants, acquired at 1 mV/s in 1 M LiPF 6 in EC:DMC (1:1 v/v). (b) The number of radicals (or equivalently polarons) detected via EPR, normalized to the number of dopants, as measured by CV, as a function of electrochemical potential. The PProDOT-Hx 2 data (black curves) is adopted from our previous work. 45 Figure 5.7 (a) CV scans of (Hex:OE) PProDOTs at 5 mV s -1 in 1M LiTFSI dissolved in EC/DMC. (b) Electronic and (c) Ionic conductivities of the copolymer series. Figure 5.8 The integrated GIWAXS diffractograms of all the polymers in their neutral state. The neutral polymers are all relatively similar: there are two dominant peaks: the lamellar peak and the pi-stacking peak, indicating that they are not very crystalline. Figure 5.9 a) Full integration of GIWAXS diffractograms for (Hex:OE) PProDOTs, normalized to the lamellar peak under neutral conditions b) at the first oxidation peak, c) at the second oxidation peak, normalized to the π-stacking peak, d- f) select polymers with varying amounts of oligoether under neutral conditions, and doped at the first and second oxidation peak 351 Figure 5.10 2D diffractograms of the various polymers with differing amounts of oligoether content, at neutral (1 st column), at first oxidation peak (2 nd column), and at 2 nd oxidation peak (3 rd column). Figure 5.11 Full integration of GIWAXS diffractograms for (75:25) PProDOT under neutral conditions, doped at the first and second oxidation peaks. Figure 5.12 Swelling study of the (Hex:OE) PProDOTs using propylene carbonate (PC) electrolyte, showing an almost linear trend in % mass increase with increase in oligoether incorporation. Figure 5.13 Specific capacity as a function of cycle number at constant charge/discharge rate of 1C with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOTs cells with a cathode mass composition of 90% NCA, 4% binder, 6% carbon and an areal loading of 3.1 ± 0.4 mg cm -2 . Figure 5.14 Corresponding galvanostatic charge-discharge curves from Figure 5.13; a) Li-NCA-(85:15) PProDOT and Li- NCA-(95:05) PProDOT, b) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, c) Li-NCA-PProDOT-Hx 2 and Li-NCA- PVDF. Figure 5.15 Three-dimensional chart of the impedance response after 200 cycles as function of SOC. Nyquist plots as function of SOC for a) NCA-PProDOT-Hx 2, b) NCA-(95:05) PProDOT, c) NCA-(85:15) PProDOT, d) NCA-(75:25) PProDOT, e) NCA-(65:35) PProDOT, and f) NCA-PVDF cells. Figure 5.16 Nyquist plots of total impedances for (a) semi-infinite linear, (b) transmissive, and (c) reflective boundary. 85 Figure 5.17 Calculated Warburg coefficient σ as function of state of charge (SOC) for the NCA-(Hex:OE) PProDOTs, NCA-PProDOT-Hx 2 and NCA-PVDF cells after 200 cycles. Figure 5.18 Differential capacity vs voltage analysis (dQ dV ‐1 vs V) for the Li‐NCA‐(Hex:OE) PProDOTs cells from Figure 5.13-14 as a function of cycle number charged‐discharged at 1C with calculated between 2.7 and 4.2 V vs Li + /Li for a) Li‐NCA‐PProDOT‐Hx 2, b) Li‐NCA‐(65:35) PProDOT, c) Li‐NCA‐(75:25) PProDOT, d) Li‐NCA‐(85:15) PProDOT, e) Li‐NCA‐ (95:05) PProDOT, and f) Li‐NCA‐PVDF cells. Figure 5.19 a) Galvanostatic charge-discharge profiles for Li-NCA-PProDOT-Hx 2 and Li-NCA-(75:25) PProDOT cells at C/20 with an areal mas loading of ∼ 5 mg cm -2 and a charging cut-off of 4.2 V vs Li + /Li. The impedance response was measured at OCV as function of SOC (denoted with yellow circles). b) and c) Corresponding three-dimensional Nyquist plots of the experimental impedance response (denoted with color spheres) as function of SOC during charge for the same Li-NCA-(75:25) PProDOT (b) and Li-NCA-PProDOT-Hx 2 (c)cells used to generate the data in part (a). Real part of the impedance response against frequency for the Li-NCA-(65:35) PProDOT, Li-NCA-(75:25) PProDOT, Li-NCA-(85:15) PProDOT, Li-NCA-(95:05) PProDOT, Li-NCA-PProDOT-Hx 2, and Li-NCA-PVDF cells at d) 3.7 V e) 3.9 V, and f) 4.08 V vs Li + /Li. Figure 5.20 a) Specific capacity as a function of cycle number at various discharge rates with a constant charging rate at C/5 for Li-NCA-(Hex:OE) PProDOTs cells with an areal loading of 5.2 ± 0.8 mg cm -2 . b) Specific capacity as a function of cycle number with symmetric discharge/charge rate for Li-NCA-(Hex:OE) PProDOTs cells with an areal loading of 5.0 ± 0.8 mg cm -2 . In both cases, two formation cycles at C/10 were performed prior to rate testing. Figure 5.21 Corresponding galvanostatic charge-discharge curves from Figure 5.20a; (a) Li-NCA-(85:15) PProDOT and Li-NCA-(95:05) PProDOT, (b) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, (c) Li-NCA-PProDOT-Hx 2 and Li- NCA-PVDF. Corresponding galvanostatic charge-discharge curves from Figure 5.20b; (d) Li-NCA-(85:15) PProDOT and Li-NCA-(95:05) PProDOT, (e) Li-NCA-(65:35) PProDOT and Li-NCA-(75:25) PProDOT, (f) Li-NCA-PProDOT-Hx 2 and Li- NCA-PVDF. Figure 5.22 Differential capacity vs voltage analysis (dQ dV -1 vs V) for the Li-NCA-(Hex:OE) PProDOT family cells from Figure 5.20a with increments in the discharge rate and a constant charging rate of C/5 with an areal mass loading of 5.2 ± 0.8 mg cm -2 calculated between 2.7 and 4.2 V vs Li + /Li for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(65:35) PProDOT, c) Li-NCA-(75:25) PProDOT, d) Li-NCA-(85:15) PProDOT, e) Li-NCA-(95:05) PProDOT, and f) Li-NCA-PVDF cells. 352 Figure 5.23 Differential capacity vs voltage analysis (dQ dV -1 vs V) for the Li-NCA-(Hex:OE) PProDOT family cells from Figure 5.20b with increments in the discharge/charge rate with an areal mass loading of 5.0 ± 0.8 mg cm -2 calculated between 2.7 and 4.2 V vs Li + /Li for a) Li-NCA-PProDOT-Hx 2, b) Li-NCA-(65:35) PProDOT, c) Li-NCA-(75:25) PProDOT, d) Li-NCA-(85:15) PProDOT, e) Li-NCA-(95:05) PProDOT, and f) Li-NCA-PVDF cells. Figure 5.24 Real part of the impedance response Zre against frequency for the NCA-(75:25) PProDOT, NCA-PProDOT- Hx 2, and NCA-PVDF cells with cathode composition with a) 2% binder, 3% carbon, 95% NCA, b) 1% binder, 3% carbon, 96% NCA and c) 1.5% carbon, 4% binder, 94.5% NCA at 50% SOC Figure 5.25 Three-dimensional chart of the impedance response as function of SOC. Nyquist plots as function of SOC for a) 2% binder NCA-PVDF, b) 2% binder NCA-PProDOT-Hx 2, c) 2% binder NCA-(75:25) PProDOT, d) 1% binder NCA- PVDF, e) 1% binder NCA-PProDOT-Hx 2, f) 1% binder NCA-(75:25) PProDOT, g) 1.5% Carbon NCA-PVDF, h) 1.5% Carbon NCA-PProDOT-Hx 2, and i) 1.5% Carbon NCA-(75:25) PProDOT cell. Figure 5.26 Specific capacity as a function of cycle number with increments in the discharge/charge rate (symmetric charge/discharge rate capability) with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOT cells with a cathode mass composition of a) 95% NCA, 2% binder, 3% carbon with an areal loading of 5.5 ± 0.3 mg cm -2 and b) 96% NCA, 1% binder, 3% carbon with a loading of 5.3 ± 0.6 mg cm -2 . Figure 5.27 Corresponding galvanostatic charge-discharge curves from Figure 5.26a; (a) Li-NCA-(75:25) PProDOT, (b) Li-NCA-PProDOt-Hx 2, and (c) Li-NCA-PVDF. Corresponding galvanostatic charge-discharge curves from Figure 5.26b; d) Li-NCA-(75:25) PProDOT, e) Li-NCA-PProDOT-Hx 2, and f) Li-NCA-PVDF. Figure 5.28 a) Specific capacity as a function of cycle number with increments in the discharge/charge rate with two formation cycles at C/10 for the Li-NCA-(Hex:OE) PProDOT cells with a cathode mass composition of 94.5% NCA, 4% binder, 1.5% carbon and an areal loading of 4.9 ± 0.5 mg cm -2 . Corresponding galvanostatic charge-discharge curves in a) for; b) Li-NCA-(75:25) PProDOT, c) Li-NCA-PVDF, and d) Li-NCA-PProDOT-Hx 2. Figure 6.1 Li xNaNb 13O 33||Li cell galvanostatic charge/discharge (GCD) profile and differential capacity analysis. a) Galvanostatic profile during lithiation and b) delithiation between 1 and 3 V vs Li + /Li at C/15. c) Corresponding dQ dV -1 vs V plot for the GCD curves in a) and b). regions are labeled with Roman numerals and the phase transitions (peaks) with Arabic numerals. Figure 6.2 Rate Capability and incremental capacity analysis of a Li xNaNb 13O 33||Li cell. a) Specific capacity as function of cycle number with increments in C-rate. b) Three-dimensional ICA plot from the GCD curves from Figure 6.3. Figure 6.3 Representative galvanostatic charge-discharge curves at various C-rates for Li xNaNb 13O 33||Li cells with an areal active material mass loading of a) 2 mg cm -2 and b) 5.5 mg cm -2 . Figure 6.4 Incremental capacity analysis (ICA) as a function C-rate (two-dimensional plot) from the GCD curves in Figure 6.3 Identified peaks are denote with numeric labels. Figure 6.5. Li xNaNb 13O 33||Li cell peak analysis as function of C-rate during lithiation a) Peak position, b) Peak height, c) peak area, d) peak width. Figure 6.6 Li xNaNb 13O 33||Li cell specific capacity as a function of cycle number and differential capacity analysis. a) long-term cycling between 1 and 3 V vs Li + /Li at 2C. b) Three-dimensional ICA plot from the GCD curves in Figure 6.7 Figure 6.7 Representative galvanostatic charge-discharge curves as function of cycle number at 2C for a Li xNaNb 13O 33||Li cell with an areal active material mass loading of 2.2 mg cm -2 . Figure 6.8 Incremental capacity analysis (ICA) as a function of cycle number (two-dimensional plot) from the GCD curves in Figure 6.7 Identified peaks are denote with numeric labels. Figure 6.9 Li xNaNb 13O 33||Li cell peak analysis as function of C-rate during lithiation a) Peak position, b) Peak height, c) peak area, d) peak width. 353 Figure 6.10 Galvanostatic lithiation curve for the Li xNaNb 13O 33||Li cell at C/20. EIS measurements at open circuit voltage are denoted with yellow circles. Figure 6.11 Electrochemical impedance characterization as a function of potential vs Li + /Li during lithiation plotted in 3-D Nyquist plots for a) the 1 st cycle after formation, and b) after 600 th cycles Figure 6.12 a) Equivalent electrical circuit used for fitting the impedance data in Figure #. Change of the b) ohmic electrolyte resistance R1, c) the constant phase element Q1, d) a value, e) charge transfer resistance R2, and f) Warburg coefficient S1 as function of potential vs Li+/Li for the 1st (denoted with green dots) and 600th cycle (denoted with purple dots). Relative fitting errors are indicated with error bars. Figure 6.13 Real part of the impedance response Zre in the Warburg region as a function of the inverse square of the angular frequency ω -0.5 and their linear fitting for each potential vs Li + /Li during lithiation. Figure 6.14 Apparent diffusion coefficient by EIS during lithiation for the 1 st and 600 th cycle assuming semi-infinite conditions. Figure 6.15 Galvanostatic profiles for NCA||Li (yellow), NaNb 13O 33||Li (red) and Full-cell paring of both materials NCA|| NaNb 13O 33 (purple). Figure 6.16 Rate Capability testing and their correspond galvanostatic profiles for (a-b) NCA-PProDOT-Hx 2||Li and (c- d) NaNb 13O 33||Li cells. Figure 6.17 a) Rate capability and b) corresponding GCD curves for the NCA-PProDOT-Hx 2||NaNb 13O 33 cell with a N/P ratio of 1.3 an areal loading of 4 mg cm -2 . Figure 6.18 a) Cycle-life comparison and characterization of the NCA||NaNb 13O 33 full-cells at 1C with an NCA areal loading of 4 mg cm -2 and a N/P ratio of 1.3. Corresponding GCD curves from 6.18a for the b) NCA-PProDOT- Hx 2||NaNb 13O 33 and c) NCA-PVDF||NaNb 13O 33 cells. Figure 6.19 a) Schematics of the NCA and NaNb 13O 33 pouch cell electrodes b) representation of the electrode assembling were a polypropylene seperator is placed between the electrodes to prevent a short-circuit. Figure 6.20 Photograph of the NCA-PProDOT-Hx 2 and the NaNb 13O 33 pouch-cell electrodes coated onto Al foil with an Al tab and onto Cu foil with a Ni tab respectively. Figure 6.21 Photograph of the assembled SCALAR-EFRC first NCA-PProDOT-Hx 2||NaNb 13O 33 pouch-cell. Figure 6.22 First Cycle GCD for the SCALAR-EFRC pouch-cell NCA-PProDOT-Hx 2||NaNb 13O 33 with an approximate mass loading of 4.7 mg cm -2 and an N/P ratio = 2 cycled at C/27. Figure. 7.1 The first charge and discharge curves (C/50 charge and discharge rate) for the sulfur electrodes prepared using KB-300 and KB-600. Figure. 7.2 (a) The first discharge curves with 1.2 V voltage cut-off at (i) C/50 rate and (ii) C/20 rate, and (b) the discharge curve at C/20 rate with 0 V voltage cut-off. Figure. 7.3 EIS analysis during first discharge: Here (a) Nyquist plots at different DODs, and representative Nyquist plots for the cells discharged (b) below and (c) above 35 % DOD. Inset shows corresponding equivalent circuits. Figure. 7.4 (a) Open circuit voltage variation with DOD, (b) series resistance (R0) changes with DOD, (c) resistance from HFS (R1) variation with DOD, (d) a plot of lnR1 vs. OCV, (e) effect of lithium polysulfides on impedance of Li/Li symmetric cells, and (f) variation of charge transfer resistance (R2) with DOD. Figure. 7.5 (a). First discharge curves at C/50 for the sulfur electrodes prepared by (i) mixing and (ii) melt infusion processes and (b) Nyquist plots for the lithium-sulfur cell with sulfur electrode prepared using melt infusion technique. The variation in the (c) R0, (d) R1 and (e) R2 with DOD. Figure. 7.6. (a) TEM image of KB particle infused with sulfur, and element mapping of (b) carbon and (c) sulfur. 354 Figure. 7.7 (a) First discharge curves. Inset shows the enlarged region near 30 % DOD. (b) Nyquist plots at different DOD for a cell with hybrid sulfur electrode. The variation in (c) R0, (d) R1 and (e) R2 with the depth of discharge. Figure 8.1. Schematic representation of the solid-state composite-electrode/bi-layer electrolyte lithium-sulfur battery. Figure 8.2. (a) MCM coated with PEO, an example of the bilayer solid electrolyte. (b) Cross-section scanning electron microscopy image of the bilayer. Figure 8.3. (a) Comparison of the discharge curve at C/60 between the CEBE Li-S cell and a liquid electrolyte Li-S cell during the first cycle. (b) A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of discharge for a liquid electrolyte Li-S cell and (c) for the CEBE Li-S cell. (d) Equivalent circuits proposed to describe the Liquid and CEBE Li-S cells. Figure 8.4. (a) Change of the internal resistances as function of the of the state of discharge for the Liquid Li-S cell. (b) Change of the internal resistances as function of the of the state of discharge for the CEBE Li-S cell. Figure 8.5. (a) Galvanostatic charge/discharge curves of the CEBE Li-S battery for the first and fourth cycle at rate of C/20 (1C=1675 mA g -1 ). (b) Galvanostatic charge/discharge curves of the CEBE Li-S battery for the first and second cycle at C/8. Figure 8.6. Impedance response comparison of the CEBE Li-S cell before cycling (depicted with red circles) and after eight cycles (depicted with black squares). Figure 8.7. (a) Specific Capacity against cycle number plot comparison between the CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI at C/20 (depicted with squares) and a CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI- Al 2O 3 at C/20 (depicted with circles) (b) Galvanostatic charge/discharge curves for the CEBE Li-S cell with the bilayer made of MCM/PEO-LiTFSI-Al 2O 3 for the first and eight cycle. Figure 8.8 Shuttle current measurement of the mass-efficient CEBE Li-S cell (TAG). a) Potentiostatic specific current as a function time at 2.7 V vs Li + /Li. b) Steady-state specific shuttle current as a function of potential in the mass-efficient CEBE Li-S cell. Figure 8.9 Galvanostatic charge-discharge profiles for the mass-efficient CEBE Li-S cell at C/20 with a charging cut-off of 2.7 V vs Li + /Li. EIS was measured as function of SOC at OCV (denoted with yellow circles). Figure 8.10 A three-dimensional representation of the impedance response captured in Nyquist plots as function of the state of discharge the mass-efficient CEBE Li-S cell during a) discharge and b) charge. Figure 8.11 Real part of the impedance response Zre against frequency for the mass-efficient CEBE Li-S cell during the at 18%, 50%, and 75% SOC during a) discharge and b) charge. Figure 8.12 (a) Specific capacity against cycle number plot for the mass-efficient CEBE Li-S cell (b) Corresponding galvanostatic charge/discharge curves for the mass-efficient CEBE Li-S cell in at C/9 Figure 8.13 Galvanostatic discharge curve (potential vs specific energy at the cell-level) comparison of the CEBE Li-S cell baseline from 2020 against the mass-efficient 2022 TAG cell. Figure 9.1. Polymers investigated for binders in Lithium-Sulfur (Li-S) batteries. Figure 9.2. Cyclic Voltammograms of N2200 and N2200-OE. Initial CV data (a and b), CV data as a function of scan rate (c and d), and long-term cycling (e and f) for N2200 and N2200-OE, respectively. Figure 9.3. (a) Cyclic voltammograms of N2200 and N2200-OE at 5 mV s −1 . (b) Electronic and (c) ionic conductivity of N2200 and N2200-OE as a function of potential. 355 Figure 9.4. 2D diffractograms of N2200 (left column) and N2200:OE (87:13) (right column) at the neutral, undoped state (top row), doped at the 2.4 V vs. Li/Li+, and 2.1 V vs. Li/Li + . Figure 9.5. Integrated GIWAXS diffractograms. Panel (a) shows the full integration for N2200 plotted on a linear scale, and the inset shows the in-plane integration plotted on a log scale. Panel (b) shows the full integration of N2200-OE plotted on a linear scale, and the inset shows the in-plane integration plotted on a log scale. Figure 9.6. UV-Vis spectra of LiPS exposed to (a) PVDF, (b) N2200, and (c) N2200-OE. (d) Absorbance at 470 nm of LiPS solution over 3 hours of polymer exposure. Figure 9.7. Photographs of LiPS (a) before exposure to polymers and after 3 hours of exposure to (b) PVDF, (c) N2200, and (d) N2200-OE. Figure 9.8 (a) Potentiostatic intermittent shuttle current measurement curve of a Li-Sulfur-N2200 cell. (b) Voltage profile prior to and during shuttle current measurements on a Li-Sulfur-N2200 cell. Figure 9.9. Polysulfide shuttle current measurement. (a) Observed current as a function of time at approximately 2.3 V vs Li + /Li after an OCV relaxation step and (b) steady-state current or shuttle current as a function of potential during discharge for sulfur-N2200 (denoted with green circles), sulfur-N2200-OE (denoted with pink triangles), and sulfur-PVDF electrodes (denoted with yellow squares). Figure 9.10. Steady-state current or shuttle current as a function of potential after 35 cycles during discharge for a Li- Sulfur-N2200 (denoted with red circles) and Li-Sulfur-PVDF cells (denoted with black circles). Figure 9.11. (a) GCD profiles for Li-sulfur-N2200 and Li-sulfur-N2200-OE cells at C/20. Potentiostatic impedance response was measured at OCV as a function of SOC during charge and discharge (denoted with yellow circles). (b) Real part of the impedance response against frequency for the Li-sulfur-N2200, Li-sulfur-N2200-OE, and Li-Sulfur-PVDF cells at approximately 2.28 V, 2.15 V, and 2.1 V vs Li + /Li. Figure 9.12. Three-dimensional chart of the impedance response as function of SOC. Nyquist plots as function of SOC for (a) sulfur-N2200, (b) sulfur-N2200-OE, and (c) sulfur-PVDF during charge. (d) sulfur-N2200, (e) sulfur-N2200-OE, and (f) sulfur-PVDF during discharge. Figure 9.13. Rate Capability for Li-sulfur-n-dopable polymer cells under moderate and high loadings. Specific capacity as a function of cycle number with increments in the discharge rate and a constant charging rate at C/10 for sulfur electrode with (a) moderate loadings and (b) high loadings. (c) Specific capacity as a function of cycle number at a discharge/charge rate of C/2 for sulfur-N2200, sulfur-N2200-OE and Sulfur-PVDF. Impedance response as function of cycle number at approximately 2.2 V vs Li + /Li for the same (d) Li-sulfur-N2200, (e) Li-sulfur-N2200-OE, and (f) Li-sulfur- PVDF cells for the 15 th , 200 th , and 495 th cycles. Figure 9.14 Corresponding galvanostatic charge-discharge curves from Figure 9.13a; (a) Li-Sulfur-N2200, (b) Li-Suflur- N2200-OE, and (c) Li-Sulfur-PVDF. Corresponding galvanostatic charge-discharge curves from Figure 9.13b; (d) Li-Sulfur- N2200, (e) Li-Suflur-N2200-OE, and (f) Li-Sulfur-PVDF. Figure 9.15. Corresponding galvanostatic charge-discharge curves from Figure 9.13c; (a) Li-Sulfur-N2200 and (b) Li- Suflur-N2200-OE

Abstract (if available)

Abstract In the last couple of decades lithium-ion batteries have fulfilled most of the energy storage demands for mobile applications. However, there is a continuous increase in energy and power demands that current lithium-ion batteries will not be able to meet due to their theoretical limits for energy storage. Therefore, there is a necessity for understanding the pathways and designs to increase the capabilities of practical high-energy and high-power lithium batteries. This dissertation presents a systematic, comprehensive, and rationalized electrochemical study and electrode design for building next generation lithium-ion and lithium-sulfur batteries. We introduce the use new conducting polymers as cathode binders/additives, new niobium-oxide based intercalating materials with fast lithium-ion diffusivity and structural stability, and a novel composite solid-state electrolyte for lithium-sulfur batteries. Electrochemical and physicochemical characterization techniques such as Electrochemical Impedance Spectroscopy (EIS), Incremental Capacity Analysis (ICA), Cycling Voltammetry (CV), Electron Paramagnetic Resonance (EPR), and spectroscopy techniques have been utilized to achieve an in-depth understanding. We have found that using various conducting polymers as cathode binders significantly enhance the energy/power capability (320 Wh kg-1 at C/2), and cycle-life of a lithium-ion battery due to their high mixed conductivity and protection against the rapid growth of the solid electrolyte interphase. We found that NaNb13O33 as an anode material has an impressive rate capability (80 mAh g-1 at 20C) and a remarkable capacity retention (80% after 600 cycles). Finally, we introduce a unique solid-state lithium sulfur cell based on a flexible bilayer made of a lithium-ion intercalation material and a polymer electrolyte that operates at room temperature to yield as much as 85% of the theoretical capacity.

Linked assets

University of Southern California Dissertations and Theses

Conceptually similar

PDF

Advancing lithium batteries and related electrochemical technologies for a sustainable future

PDF

Electrochemical pathways for sustainable energy storage and energy conversion

PDF

Understanding the structure-property relationship in electrode materials for electrochemical energy storage

PDF

Design and characterization of flow batteries for large-scale energy storage

PDF

Modification of electrode materials for lithium ion batteries

PDF

Studies on lithium-ion battery electrolytes and three component Strecker reaction

PDF

Nanostructured silicon anode and sulfur cathode for lithium rechargeable batteries

PDF

Nanostructure design of sulfur cathodes and lithium metal anodes for lithium-ion batteries

PDF

Design of dioxythiophene conducting polymers as electrode binders in lithium-ion batteries

PDF

Electrocatalysts for direct liquid-feed fuel cells and advanced electrolytes for lithium-ion batteries

PDF

Understanding intercalation-driven structural transformations in energy storage materials

PDF

Selective fluoroalkylation methods and synthesis of water-soluble organic molecules for organic redox flow batteries

PDF

Design and modification of electrocatalysts for use in fuel cells and CO₂ reduction

PDF

Understanding the mechanism of oxygen reduction and oxygen evolution on transition metal oxide electrocatalysts and applications in iron-air rechargeable battery

PDF

Electrochemical investigations and imaging tools for understanding extracellular electron transfer in phylogenetically diverse bacteria

Asset Metadata

Creator Elizalde-Segovia, Rodrigo (author)

Core Title Understanding the role of electrode design in determining the electrochemical performance of high-energy/high-power lithium-ion and lithium-sulfur batteries

School College of Letters, Arts and Sciences

Degree Doctor of Philosophy

Degree Program Chemistry

Degree Conferral Date 2022-12

Publication Date 02/16/2024

Defense Date 07/12/2022

Publisher University of Southern California (original), University of Southern California. Libraries (digital)

Tag batteries,cell design,conducting polymers,electrochemical impedance spectroscopy,electrochemistry,electrode design,energy storage,lithium,lithium-ion batteries,lithium-sulfur batteries,OAI-PMH Harvest

Format application/pdf (imt)

Language English

Contributor Electronically uploaded by the author (provenance)

Advisor Narayan, Sri R. (committee chair), Prakash, Surya G.K. (committee chair), Ravichandran, Jayakanth (committee member), Thompson, Barry C. (committee member)

Creator Email elizalde@usc.edu,rodriloy@hotmail.com

Permanent Link (DOI) https://doi.org/10.25549/usctheses-oUC111376958

Unique identifier UC111376958

Legacy Identifier etd-ElizaldeSe-11141

Document Type Dissertation

Format application/pdf (imt)

Rights Elizalde-Segovia, Rodrigo

Type texts

Source 20220819-usctheses-batch-974 (batch), University of Southern California (contributing entity), University of Southern California Dissertations and Theses (collection)

Access Conditions The author retains rights to his/her dissertation, thesis or other graduate work according to U.S. copyright law. Electronic access is being provided by the USC Libraries in agreement with the author, as the original true and official version of the work, but does not grant the reader permission to use the work if the desired use is covered by copyright. It is the author, as rights holder, who must provide use permission if such use is covered by copyright. The original signature page accompanying the original submission of the work to the USC Libraries is retained by the USC Libraries and a copy of it may be obtained by authorized requesters contacting the repository e-mail address given.

Repository Name University of Southern California Digital Library

Repository Location USC Digital Library, University of Southern California, University Park Campus MC 2810, 3434 South Grand Avenue, 2nd Floor, Los Angeles, California 90089-2810, USA

Repository Email cisadmin@lib.usc.edu