The Pennsylvania State University. The Graduate School. College of Engineering A MULTI-FACETED APPROACH TOWARDS IMPROVING THE PERFORMANCE OF

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1 The Pennsylvania State University The Graduate School College of Engineering A MULTI-FACETED APPROACH TOWARDS IMPROVING THE PERFORMANCE OF SILICON ELECTRODES FOR NEXT-GENERATION LITHIUM-ION BATTERIES A Thesis in Mechanical Engineering by Michael J. Melnyk 2015 Michael J. Melnyk Submitted in Partial Fulfillment of the Requirements for the Degree of Master of Science May 2015

2 The thesis of Michael J. Melnyk was reviewed and approved* by the following: Donghai Wang Associate Professor of Mechanical Engineering Thesis Advisor Bo Cheng Assistant Professor of Mechanical Engineering Dan Haworth Professor of Mechanical Engineering Professor-In-Charge of MNE Graduate Programs *Signatures are on file in the Graduate School

3 iii ABSTRACT Although lithium-ion battery technology has been the catalyst in enabling modern electric vehicle, mobile device, and large-scale energy storage technology, the increasing power demands by end-users has motivated research in developing the next-generation of lithium-ion batteries. This next generation of batteries will need to achieve higher energy and power densities, while remaining chemically stable. Silicon-based active material has been proposed as a solution in achieving superior battery performance, as it can offer a lithium storage capacity (4200 mah/g) tenfold higher than the carbonaceous electrodes employed in commercial Li-ion cells, while also offering superior safety characteristics. Unfortunately, the higher lithium storage capacity translates into an immense volume expansion ( %) upon lithiation, and thus the mechanical integrity and electrochemical performance of the electrodes are very unstable. Within the past decade, the performance of Si-based electrodes has been greatly improved as active material morphologies, polymer binders, electrolyte additives, and theoretical models have provided solutions in alleviating the stresses and strains generated during Si lithiation/delithiation. A multi-faceted solution pathway is enacted in this research to develop a Si-based electrode that can achieve cycling performance relevant to industrial application, while also offering insight on the influence of several aspects of the Si-based electrode design on cycling performance. From this investigation, a Si-based electrode has been developed with carbon-coated silicon monoxide active material and polyacrylic acid polymer binder, both of which offer several complimentary attributes that enable a moderately stable cycling performance at high active mass loading while offering a gravimetric and areal lithium capacity magnitude relevant to industrial applications. Although much work lies ahead in further improving the capacity retention of the Si-based electrodes reported in this thesis, this work presents an economical platform for future work on the topic.

4 iv TABLE OF CONTENTS List of Figures... vi List of Tables... xi Acknowledgements... xii Chapter 1 Introduction... 1 Chapter 2 Fundamentals of Lithium-Ion Battery Mechanisms Battery System Components and Functions Lithium-ion Transport and Storage Mechanism Background on Alloy-Type Electrodes Challenges with the Silicon Lithium-Alloy Electrode... 9 Chapter 3 Experimental Procedures & Equipment Preparation of Active Material Preparation of the Electrode Slurry Calendaring and Porosity Control Cell Fabrication Electrochemical Characterization Chapter 4 The Influence of the Fabrication Parameters Introduction Literature Review Results and Discussion Conclusions Chapter 5 The Influence of the Active Material Chemistry Introduction Literature Review Results and Discussion Conclusions Chapter 6 The Influence of the Non-Active Polymer Binder Introduction Literature Review Results and Discussion Conclusions Chapter 7 Conclusions & Future Work Chapter 8 Early Development of a Battery Cell for In Situ Characterization... 79

5 v 8.1 Introduction Design and Fabrication of the In Situ Battery Cell Materials and Compatibility Early Stages of Characterization Future Outlook Appendix A In Situ Cell Bill of Materials Appendix B Schematic Diagram of the In Situ Cell Bibliography... 95

6 vi LIST OF FIGURES Figure 2-1 (pg. 8): Schematic diagram of the electrochemical reaction occurring within a lithium-ion battery during charge or discharge. This schematic was originally published by Nishi [3] Figure 3-1 (pg.17): Image of the casting process Figure 3-2 (pg.18): Correlation between the deformed thickness of an electrode sheet and the roller gap width. Figure 3-3 (pg.19): The device used for calendering the electrodes. Figure 3-4 (pg.21): The internal components of the coin cell assembly and the device used to crimp and seal the coin cells. Figure 4-1 (pg. 25): An example of the cracking and delamination that commonly occurs when preparing a high mass loading electrode that has not be properly handled or has had its moisture evaporated too quickly. Figure 4-2 (pg. 28): Electrochemical performance of high active mass electrodes composed of a novel pomegranate-inspired silicon electrodes. The electrodes were first cycled at 0.03 ma/cm 2 for three cycles and 0.7 ma/cm2 thereafter. Data referenced from Liu et al. [38] Figure 4-3 (pg. 29): The influence of active mass loading on the cycling performance of a Sibased electrode composed with a novel self-healing polymer binder developed and reported by Chen et al. [39] Figure 4-4 (pg. 32): Correlation between the thickness of NCM electrodes and a) rate capability, b) weight specific impedance, and c) capacity retention developed and reported by Zheng et al. [40] Figure 4-5 (pg. 34): Influence of calendaring and the induced porosity changes on the performance of a lithium-alloy negative electrode with TiSnSb active material as developed and reported by Nguyen et al. [41]

7 vii Figure 4-6 (pg. 36): The correlation between active material content fraction in an electrode composition on the cycling performance of Si nanoparticle-based electrode developed and reported by Beattie et al. [43] Figure 4-7 (pg. 38): Comparison of the a) gravimetric capacity, b) 1st and 10th cycle gravimetric capacity, c) 1st cycle coulombic efficiency, and d) areal capacity of carbon-coated silicon (Si-C) anodes prepared with varying active material mass loadings. The lithium half cells are cycled at 400 ma/g in a voltage window of V. Figure 4-8 (pg. 41): Comparison of a) gravimetric capacity, b) voltage profiles, and c) areal capacity of pristine silicon monoxide anodes prepared with 1mg/cm2 (red) and 2.5 mg/cm2 (black) mass loading. The lithium half cells are cycled at 400 ma/g in a voltage window of V. Figure 4-9 (pg. 43): Comparison of the a) gravimetric capacity and b) areal capacity of carboncoated silicon monoxide anodes prepared with high mass loading of 4.6 and 6.0 mg/cm2. The lithium half cells are cycled at 100 ma/g in a voltage window of V. Figure 4-10 (pg. 45): Comparison of the a) gravimetric capacity and b) areal capacity of carboncoated silicon monoxide anodes prepared with varying electrode porosity and active mass loadings. The lithium half cells are cycled at 50mA/g in a voltage window of V. Figure 5-1 (pg. 50): Schematic diagram of the multi-domain structure and porous carbon surface coating of a SiO-based active material developed and reported by Zhao et al. [49] Figure 5-2 (pg. 50): The electrochemical performance of SiO-based electrodes incorporated with an electronically-conductive type PFM polymer binder developed and reported by Zhao et al. [49] Figure 5-3 (pg. 52): Comparison of the electrochemical performance of pure SiO and mixed Si- SiO electrodes developed and reported by Doh et al. [50]

8 viii Figure 5-4 (pg. 53): A comparison of the electrochemical performance of SiO-based electrodes disproportionated at different temperatures prior to electrode fabrication, developed and reported by Park et al. [44] Figure 5-5 (pg. 55): The electrochemical cycling of a carbon-coated SiO electrode developed and reported by Kim et al. [54] Figure 5-6 (pg. 56): A comparison of the electrochemical performance of a pristine nanocrystalline Si electrode (top) and a carbon-coated Si nanocomposite electrode (bottom) developed and reported by Ng et al. [56] Figure 5-7 (pg. 58): A comparison in the electrochemical rate performance (left) and electrical impedance (right) of Si-based electrodes with carbon-coatings deposited via a chemical vapor deposition method at different temperature, developed and reported by Yi et al. [51] Figure 5-8 (pg. 60): Comparison of the cycling performance of pristine silicon monoxide (SiO) and carbon-coated silicon monoxide (SiO-C) anodes prepared at similar mass loading of ca. 1 mg/cm2 and material composition ratio of conductive additives and PAA polymer binder. The lithium half cells are cycled at 400mA/g within a voltage window of V. After cycle 10, the temperature of the SiO-C environment is increased from room temperature to 60oC. Figure 6-1 (pg. 66): A comparison of the cycling performance of electrodes employing CMC and PVdF derivative binders as developed and reported by Hochgatterer et al. [72] Figure 6-2 (pg. 67): Electrochemical cycling regime and associated XPS spectra for graphite/cmc anode and LiCoO2 full cell system as developed and reported by El Outani et al. [74] Figure 6-3 (pg. 68): The influence of the molecular weight of the CMC polymer binder, at 250,000 (grey squares) or 700,000 g/mol (black diamonds) on the cycling performance of Sibased electrodes developed and reported by Bridel et al. [75]

9 ix Figure 6-4 (pg. 69): Comparison in the cycling performance of SiO-based electrodes composed of a) PVdF, b) PVA, c) NaCMC, or d) PAA polymer binder as developed and reported by Komaba et al. [77] Figure 6-5 (pg. 70): Comparison of the cycling performance of Si/graphite electrodes composed of PVdF, PAA, or crosslinked PAA:PCD polymer binders with different ratios of PAA and PCD, as developed and reported by Han et al. [81] Figure 6-6 (pg. 71): Comparison of cycling performance of Si-based electrode with PVdF, PAA, CMC, or crosslinked PAA-CMC polymer binders as developed and reported by Koo et al. [31] Figure 6-7 (pg. 74): Comparison of the cycling performance of SiO-C electrodes composed with PVdF, NaCMC, PAA, or PAA:NaCMC polymer binders. All electrodes are cycled within a voltage range of V and a current density of 100 ma/g. Figure 6-8 (pg. 74): Voltage diagrams at the 1st, 2nd, and 40th electrochemical cycles of the SiO-C electrodes composed with PAA and PAA:NaCMC polymer binders. Figure 8-1 (pg. 82): Rendering of the initial conceptual design of the in situ battery cell. Figure 8-2 (pg. 83): An image of the actual fabricated in situ cell without the electrode assembly or transparent glass windows installed. Figure 8-3 (pg. 84): An image of the actual fabricated in situ cell with the transparent glass window and top bracket installed, but without the electrode assembly installed. Figure 8-4 (pg. 86): Image of the completely assembled and electrolyte-filled in situ cell in an argon-filled glove box. Figure 8-5 (pg. 88): Leak test of the in-situ cell. Lithium-foil is encased in the in-situ cell, previously assembled in an argon-filled glove box, and left in an air-filled environment. The insitu cell remains suitable air-free for about 6 days before signs of lithium oxidation occur. Figure 8-6 (pg. 90): Electrochemical evaluation of the influence of O-ring materials on cycling performance. For this evaluation, a SiO/C active material electrode is cycled at 0.1 C-Rate.

10 x Figure 8-7 (pg. 91): Optical microscope images of the material within the in situ cell. Figure B-1 (pg. 94): Schematic of the Bottom-Face of In Situ Cell Figure B-2 (pg. 94): Schematic of the Top-Face of In Situ Cell Figure B-3 (pg. 94): Schematic of the Side-Face of the In Situ Cell

11 xi LIST OF TABLES Table 2-1 (pg. 9): Comparison of electrochemical characteristics of carbon, lithium, and lithiumalloys, adapted from Zhang et al. [21] Table A-1 (pg. 93): Bill of materials for the in situ cell

12 xii ACKNOWLEDGEMENTS Project Sponsorship U.S. Department of Energy: My utmost gratitude goes out to the U.S. Department of Energy for their funding and sponsorship on the Advanced Battery Research (ABR) project (Project Title: High Energy, Long Cycle Life Lithium-Ion Batteries for PHEV Application). Through their support, I have been given an invaluable opportunity to evolve as an engineer and apply my knowledge towards state-of-the-art battery research. The Research Team Dr. Donghai Wang: Dr. Wang, I deeply appreciate your engineering mentoring, financial support, and research guidance over the past several years. I have been fortunate to have an advisor whom advocated my involvement in all stages of engineering research on fascinating projects related lithium-ion batteries, while also encouraging personal scientific endeavors. Dr. Bo Cheng: Thank you for reviewing my thesis in the midst of your very busy schedule. I wish you the best of luck with your new career at Penn State University! Mikhail Gordin, Michael Regula, Adnan Mousharraf, Ran Yi, Shuru Chen, Jiangxuan Song, Shi Hu, Duihai Tang, Zhaoxin Yu, Yue Gao, Qingquan Huang: Aside from recognizing your remarkable commitment, patience, and passion for your science; I would also like to acknowledge your interesting and fun personalities. I am very proud to call you my friends and wish you all the best of luck in your lives; I believe you are all destined for great things. Family Janet M. Melnyk (Mom): A proper thank you and recognition would be a thesis, in and of itself. From cooking the best food, to doing my laundry, to always having my back and cheering

13 me on; you truly deserve top honors! I love you very much, mom. Thanks for everything that you are! xiii Joseph M. Melnyk (Grandpop): You are the reason that I became an engineer. My motivation and commitment have been sparked by your endless and fascinating stories: everything from your trials and tribulations at Burpee s Seed Company to your awe-inspiring engineering stories about developing rocket and satellite systems. Thanks for all your love and support! Michelina L. Melnyk (Grandmom): You are the most kind-hearted and caring grandmother that a guy could ask for. Thanks for always making me feel special and sending newspaper articles about science. Love you very much! Carole A. Petner: Over the past nine years, you have grown into my best friend and true love. Thanks for always listening to me banter about deadlines and burdens over the course of my education. I look forward to our journey together. Roxie: Perhaps the most peculiar creature in existence, it is a shame you will never read this. Nonetheless, you always put a smile on my face after a long day, and that deserves recognition. May your life be full of sunny windowsills, shiny baubles, and ½ cup of food per day. =^.^=

14 1 Chapter 1 Introduction Green energy, electric vehicle transportation, long-lasting portable technology, and smart-grid energy distribution: these are the daunting objectives inadvertently established by today's digitized and power-demanding global society. A common denominator presents itself as a catalyst for the successful application of these technologies: secondary energy storage. That is, high energy- and power-dense electrochemical devices (e.g., batteries, ultracapacitors, etc.) that have the potential to store and deliver massive amount of electrical energy on-demand, all while maintaining a long life span in a variety of natural environments. Indeed, state-of-the-art batteries, lithium-ion batteries (LIBs), are already being applied as the keystone to our most advanced commercial and military technologies [1]. In the transportation industry, many car manufacturers have at least one vehicle model that currently utilizes the Li-ion battery, offering improved driving performance and mileage compared to older battery chemistries such as NiMH. Regardless if one is a battery enthusiast, it is difficult to deny the excitement that has been generated by Tesla Motors. In spite of being a new market and large price-tag vehicles, Tesla Motors has made ground-breaking progress in moving electric vehicle closer to the performance offered by conventional gas-engine vehicles. The technology inside many people's pockets, the smartphone, is also a product successful primarily from the advancements made in battery technology. Intelligent engineering has enabled high-energy Li-ion batteries to fit within a device crammed with circuitry, processors, and accessories and yet is about as thick as a pencil, while still offering a suitable life-cycle without over-heating. It was only a few decades ago that our cell phones were referred to as "bricks" and the battery charger was more of a permanent fixture than an external device. Both the public and industry have a desire to establish a smart-grid that will accomplish a more efficient power generation and distribution, and in turn lower the public's cost

15 2 of power consumption. However, extensive infrastructure will need to be developed to enable this rewarding technology, with a key component being the energy storage used during excess power generation. Perhaps the broadest ambition, which encompasses all of the technology previously mentioned, is the world's desire for a cleaner environment enabled through green energy. Green energy sources are those such as the solar energy from the sun or the kinetic energy from flowing water and, when combine with some intelligent engineering and capital funding, can provide a significant amount of electricity to modest-sized communities and alleviate the demand on power distributors. Unfortunately, the burning of fossil fuels has proven to be a much more efficient and ample source of electrical energy, when compared to alternative energy sources of the same physical foot-print. One method of optimizing the energy available from these alternative sources is through storing excess energy in large-scale batteries or capacitors. In this sense, power distribution from a solar array or wind farm is not limited to optimal weather conditions, but instead energy can be stored in a secondary battery during peak power generation and distributed later on-demand. For the vast majority of commercial LIB applications, the battery chemistry is based on a graphitic anode (negative electrode), transition metal oxide cathode (positive electrode), and ethylene carbonate-based electrolyte [2]. In 1991, Sony revealed the stability of this chemistry, opposed to older iterations that suffered volatile thermal-runaway problems due to chemical instability within the cell [3]. Compared to more traditional battery chemistries, such as lead-acid or nickel metal hydride (NiMH), LIBs offer high specific capacities (372mAh/g), rate capabilities, no memory-effect, and long cycle and shelf life [4]. Perhaps most important, the LIBs were finally engineered to be safe in a variety of environments. Technology, however, is outpacing the energy and power density offered by classic LIBs. Quite simply, these LIBs struggle to supply or store sufficient energy to drive our all-electric vehicles sufficient distances or keep our smartphone powered for more than a day. Naive strategies such as increasing the

16 3 number of cells or cycling close to operational limits leads to extremely high cost and chemical instabilities in the battery system [5, 6]. In effect, many academic institutions and government agencies are funding research for the next-generation of LIBs, an ambitious endeavor that moves closer to its target goals ever year, yet is plagued by persistent obstacles. The development of next-generation LIBs will not be solely achieved by modifying one component of the battery, but instead will require a multi-disciplinary engineering approach to synthesize and optimize battery components - active material, electrolyte, separator, polymer binder - that can complement the functionality of the others, while remaining stable within a certain voltage window and current density. In response to this ambitious pursuit, several nextgeneration Li-ion-based battery chemistries have been the center of rigorous research within the past decade, including silicon [7], lithium-sulfur [8], and lithium-air [9, 10]. In this thesis, the focus is on LIBs that incorporate the alloy-type silicon-based active materials, more thoroughly discussed in Chapter 2. Based on the progress that has been made within the last decade on the silicon electrodes and its constituents, this novel battery system is arguably the closest to commercialization of the three aforementioned chemistries. Compared to the carbon anodes in commercial LIBs, silicon anodes have an order-of-magnitude higher theoretical specific capacity (4200 mah/g w.r.t. Li 4.4 Si) and have a low discharge potential of about 0.2V (w.r.t. Li/Li+) [11]. In terms of economic and manufacturing feasibility, silicon is optimal as it is environmentally friendly and can found as an abundant compound in the Earth's crust. However, the much higher specific capacity of Si results in its immense volume change ( %) during lithiation and delithiation [12]. This is the primary issue that has prevented silicon electrodes from becoming practical and is typically the focus point on publications involving this material. This enormous volume expansion further leads to a plethora of other electrode- and particle-scale issues: particle pulverization, electrode cracking, unstable solid-electrolyte interphase (SEI) layer, irreversible surface reactions, delamination from current collector, and loss of conductive network [10, 13].

17 4 Ultimately, these issues manifest into extremely poor cycling stability and coulombic efficiency, in spite of a high initial capacity. The solution to this problem is multi-faceted (i.e., will involve the optimization of many battery components, not only the silicon active material) and will eventually require assembling an optimized version of each component into the battery cell. Accordingly, the objective of this thesis is to investigate the role of several battery components and parameters on the electrochemical performance of Si-based electrodes, and thus achieve a fundamental understanding of their influence and future insight on enhancing their application.

18 5 Chapter 2 Fundamentals of Lithium-Ion Battery Mechanisms 2.1 Battery System Components and Functions In its broadest definition, the lithium-ion battery is an electrochemical engine, the same as its predecessor battery chemistries (e.g., Pb-acid, NiMH, etc.). The primary function of this electrochemical system is to convert between electrical and chemical energy, the direction of which is dependent on a positive or negative electrical current. In the discharge mode, the Li-ion battery will convert energy stored in chemical bonds of the electrode material into electrical energy passed through an external circuit. In the same manner, when the battery is in charge mode, incoming electrons will force the formation of chemical bonds and energy will be stored. This mechanism is more thoroughly described later in this chapter. Primary or non-rechargeable Li+ batteries (i.e., lithium metal as one electrode) are energy-depleted after a single discharge, although still have many applications where a high-specific energy and long storage times are required. On the other hand, secondary, or rechargeable, batteries are those that may undergo many charge/discharge cycles and are the focus in this thesis. As with all battery systems, there are several primary components that are essential for operation: anode, cathode, electrolyte, binder, and separator. Ideally, the chemistry of each component will complement and be compatible with the other components (e.g., the binder should, in fact, be capable of binding to the selected active material), while being electrochemically-stable within a certain voltage window. The positive electrode (i.e., cathode) and negative electrode (i.e., anode) are systems of active materials (i.e., the material in which lithium ions will intercalate or insert, respectively), conductive additives, and polymer binders. In a conventional fabrication process, a certain

19 6 quantity of active material and conductive additive will be mixed with a polymer binder to form a slurry, followed by coating the slurry onto a current collector (e.g., Cu or Al foil). Once the slurry has been sufficiently dried on the current collector and the majority of moisture evaporated, the material is collectively known as a battery electrode. In conventional Li-ion batteries, the anode material is a lithiated carbon (Li x C), while the cathode is a lithium metal oxide (LiMO2) [2]. For practical applications, electrode materials are designed to have a stable electrode potential for Li + intercalation/insertion, rapid Li + diffusivities, stable SEI layer, and low manufacturing costs and toxicity. In this sense, the optimal design of an electrode would be a mechanical and chemical robust system with a certain porosity to enable access to the electrolyte. The electrode porosity is a crucial design parameter that partially governs the degree of electrolyte uptake, an in turn the diffusivity of ions and electrons. The polymer binder is a non-active component (i.e., lithium-ions are not stored within the material) and is responsible for constraining the active material to the current collector, and therefore maintaining an electronic and ionic conductive network within the electrode. As will be discussed later, the responsibility of the binder has been complicated by high-capacity active materials, which undergo large volume expansion upon lithiation/delithiation. The lithium-ion diffusion and migration between the anode and cathode is enabled by the electrolyte, typically lithium hexafluorophosphate (LiPF 6 ) lithium salt dissolved in an ethylene carbonate-based electrolyte liquid or gel-polymer. Aside from transporting ions between the electrodes, the electrolyte will also react with the surfaces of active material in the electrode, forming what is commonly referred to as the passivation or solid-electrolyte interphase (SEI) layer. Although the SEI is a necessary components for stable charge-discharge cycling, the electrolyte must be designed with high chemical stability within a particular voltage window to prevent unwanted/unstable growth of the SEI layer.

20 7 The non-active porous separator material has an important role, as well, in preventing electrons from migrating between electrodes within the cell, while allowing negative or positive ions (i.e., Li-ions in this case) to shuttle between the electrodes upon electrochemical cycling. Typical separator materials are polyolefin materials with a thickness range of several tens of micrometers. The separator is also a critical safety component, as it can prevent short-circuiting within the cell. A typical challenge for the separator is the formation of lithium dendrites (i.e., Li plating) at the electrode surface caused by cold temperatures or slow reaction kinetics. These dendrites may pierce the separator and contact the opposing electrode, which in turn may lead to thermal runaway reactions [14, 15, 16]. 2.2 Lithium-ion Transport and Storage Mechanism Figure 2-1 is a schematic of the electrochemical reaction that occurs in the battery system comprised of the components described above. The electrodes and separator have been immersed in electrolyte to enable the migration of Li-ions between electrodes, while the separator prevents the migration of electrons. When the battery is discharged, Li-ions will flow from the oxidized anode, through the electrolyte and separator, to the reduced cathode where the ions will be intercalated and stored in the spaces between the lattice structures of the active material. Concurrently, the discharge reaction will cause electrons to flow from the anode, through the current collector and external circuit, to the cathode. The flow of the electrons in the external circuit from anode to cathode thus causes current to flow in the opposite direction. In secondary batteries, this process is reversible and applying an external power supply will cause a charge process, which will force Li-ions to migrate from the cathode to the anode (i.e., the reverse of the process described above).

21 8 Figure 2-1. Schematic diagram of the electrochemical reaction occurring in a LIB. This illustration was originally published by Nishi [3]. 2.3 Background on Alloy-Type Electrodes A particular set of metal and metalloids (e.g., Mg, Ca, Al, Si, Ge, Sn, Pb, As, Sb, Bi, Pt, Ag, Au, Zn, Cd, Hg, etc.) are able to form intermetallic phases with lithium at room temperature; these are called the lithium-alloys, and they represent a set of high specific energy active material being investigated for next-generation LIBs [17, 18, 19, 20, 21]. The charge-discharge process proceeds according to the following reversible chemical reaction:

22 9 Compared to conventional graphitic anodes that have a low packing density (LiC 6 ), many of these lithium-alloy materials can alloy with approximately 4 Li atoms per metal atom (e.g., Li 22 Si 5 ). In some cases, such as the Li-Si alloy, the packing density of lithium (PD Li ) may even exceed that offered by pure metallic Li (e.g., Li 21 Si 5 : PD Li = mol/ml) [20]. In effect, the theoretical specific capacity of the lithium-alloy materials are typically an order-of-magnitude higher than commercial carbonaceous electrodes. Additionally, the lithium-alloy class of active material typically is known to have improved safety characteristics compared to commercial carbonaceous electrodes; this is due to a lithiation onset voltage well-above the potential of metallic lithium, compared to that of graphite anodes (~0.05 V vs. Li/Li+), and thus lithiumplating may be more easily avoided. The electrochemical properties for several of these lithiumalloy materials, as well as carbon and lithium metal, are summarized in Table 1. Table 2-1. Comparison of electrochemical characteristics of carbon, lithium, and lithium-alloys, adapted from Zhang et al. [21] 2.4 Challenges with the Silicon Lithium-Alloy Electrode Although the high lithium packing density and specific capacity of the silicon lithium alloy electrode is its greatest advantage over alternative electrode types, it also the source of its

23 10 most problematic challenge. Silicon active material, along with other lithium alloy counterparts, suffers from an immense volume expansion during the process of lithiation and delithiation. If the silicon active material were to be fully-lithiated (i.e., Li 4.4 Si), the volume expansion could theoretically be as high as 400%. Considering the dense packing environment of a typical battery cell, where strips of electrodes may be tightly compressed within a coin cell or a jelly-roll format, it is easy to imagine how such a dramatic periodic volume change could quickly degrade the entire system. Since the conventional carbon-based active material can only be host to a small amount of Li (LiC 6 ), these types of electrodes typically only experience a small volume change, ca. 10%, which is a type of elastic deformation and, therefore, well-suited for long-term electrochemical cycling. However, since silicon hosts such a large amount of lithium per silicon atom (i.e., 4.4 Li atoms per 1 Si atom), it will suffers from an inelastic deformation. This inelastic deformation and large volume changes are the source of a variety of issues that ultimately lead to poor cycling performance. One of the primary failure mechanisms of the silicon electrodes is from the internal mechanical stresses generated from the intense volume expansion. The stresses will cause the pulverization of active material particles, as well as lead to electrode-scale fracturing [22, 23]. With the fracturing occurring throughout the entire electrode, active materials can quickly become isolated from the electronic and ionic conductive network, and therefore can no longer contribute to the storage of Li-ions. Furthermore, this unstable expansion and contraction of the active material can easily destroy the adhesion of the electrode film from the surface of the current collector, and thereby increase the resistance of electrical current through the external circuit or completely eliminating the primary function of the battery. Other than the mechanical failure mechanisms that overwhelm the integrity of the electrode, there is also a variety of unstable chemical reactions which occur due to the constant exposure of fresh active materials surfaces to the electrolyte. In a stable electrode environment, a passivation or solidelectrolyte interphase (SEI) layer will form on the surface of the active material during the initial

24 11 charge process. Although the SEI layer can lead to internal electronic resistance of the electrode, its role is primarily beneficial, in that a stable SEI layer will prevent further chemical reaction between the electrolyte and active material, enable an ionically conductive surface layer for active material, and enable a stable and consistent cycling performance. In Si-based electrodes, however, the formation of a stable SEI layer becomes an extreme challenge. On one hand, particle pulverization is constantly causing fresh active material surfaces to be exposed to electrolyte, and thus constantly causing the consumption of active lithium ions and electrolyte solvent to form a new SEI layer, as well as increasing electronic resistance during the course of cycling. On the other hand, the active material continues to undergo dramatic volume changes, even after pulverizing into smaller particles, and thus the SEI layer is constantly being deformed and disconnected from the active material surface. In this way, fresh active material surface is exposed and an increasingly thick SEI layer is formed, which in turn causes a longer diffusion distance for Li-ions. In extreme circumstances, such as rapid charge-discharge cycling or highpower loading, the chemical instability caused by the constant availability of fresh active material surface may cause exothermic reactions and compromise the safety of the battery system. The dramatic volume changes and particle pulverization also create a challenge for the non-active polymer binder, which is responsible for maintaining the mechanical integrity of the electrode and adhesion to the current collector. The polymer binders that have been successfully employed in carbon-based anodes have been found to offer insufficient chemical bond strength for the Si electrode system, where binder-silicon and binder-current collector bonds will be severed during the first several cycles. The culmination of all these issues originating from the intense volume expansion of the Si particles results in a poor cycling performance, where the lithium capacity of the electrode diminishes greatly within the first several cycles. Although the active material is the primary source of the cell's failure, enabling the stable cycling of a Si-based electrode will undoubtedly require the other battery parameters and

25 12 components to accommodate the high-capacity active material. For instance, it has been shown that several parameters related to the electrode fabrication process (e.g., drying temperature, coating thickness, materials composition) are critical for ensuring a more favorable cycling performance. These parameters will ultimately dictate the porosity of the electrode, aggregation of solid materials, and the vacant space surrounding active material. An electrode that is made with a high loading of active material (i.e., mg/cm 2 ) will have a high gravimetric capacity, however will most likely suffer from rapid capacity fading due to the close vicinity of active materials and the amount of electrolyte and lithium ions being consumed on the many fractured active material surfaces. On the other hand, if the electrode is made with a very thin coating, a more stable cycling performance may be achieved at the cost of an impractically low gravimetric capacity. The morphology and surface modifications of the active material also play a pivotal role in the success of Si-based electrodes. Indeed, a critical diameter for spherical Si particles has been established, below which stresses generated during volume expansion are unlikely to cause particle pulverization, and thus the mechanical integrity of the electrode is greatly improved [13, 24, 25, 26, 27]. Other studies have shown Si nanowires, which are grown directly on the current collector, are ideal for remaining stable during lithiation and maintaining a constant electronic conductive network due to their direct contact to current collector [28, 29, 30]. As with most synthetic nanomaterials, however, there is often a problem in justifying the cost and yield of the synthesis route. Although these materials may at first seem to be the ideal solution, the restrictions of their uneconomical synthesis routes limit them to small-scale laboratory experiments, unsuitable for the large-scale manufacturing necessary for the commercialization of battery electrodes. As mentioned, the non-active polymer binder will also need to be selected to accommodate the large volume expansion of the Si. In this sense, polymer binders with

26 13 functionalities attuned to form stronger bonds with the surface of Si active materials, while remaining chemically-stable towards the electrolyte within the operation voltage window, need to be incorporated in the electrode system. With all of these factors considered, this thesis aims to develop a more stable Si electrode system by evaluating the influence of the fabrication parameters, active material type, and polymer binder selection on the cycling performance of the electrode. The following chapters will present a multi-faceted approach to improving the performance of silicon electrodes, opposed to optimizing a singular aspect of the electrode. Chapter 4 will investigate the influence of active mass loading and electrode pressing on the cycling performance of the electrode. With optimal fabrication parameters chosen, Chapter 5 will identify the beneficial influence of a easily scalable carbon-coating technique on the cycling performance of a silicon monoxide type electrode. Lastly, several improved polymer binders are investigated for their superior chemical bonding and stabilizing attributes towards the silicon electrode and related to an improved cycling performance.

27 14 Chapter 3 Experimental Procedures & Equipment 3.1 Preparation of Active Material In this thesis, the primary silicon-type active material being investigated is silicon monoxide (SiO). The active material is either of the baseline commercial silicon monoxide microparticles (Sigma Aldrich #262951, ~325mesh) or the active material has been modified with a thin surface layer of carbon. In some cases, commercial graphite flakes (Alfa Aesar #43209, 325 mesh) are also employed as a control sample. In regards to the process of carbon-coating the SiO microparticles, approximately 1 2 grams of silicon monoxide powder is measured and deposited into a ceramic holder for insertion into a horizontal quartz tube of a high-temperature tube furnace (MTI Corp., GSL-1100X-UL). Prior to the heat treatment of the SiO, the quartz tube is first purged with acetylene gas at a flow rate of 1500 sccm for 15 minutes. Following this purge stage, the flow rate is reduced to 100 sccm and the carbon-coating of SiO occurs by the thermal decomposition of the acetylene gas at 700 o C; this temperature is reached with a ramping rate of 10 o C/min followed by holding this temperature for 20 minutes. Once the heat of the tube furnace has been reduced to below 40 o C, the carbon-coated silicon monoxide (SiO-C) is removed and is ready for incorporation into an electrode slurry.

28 Preparation of the Electrode Slurry Unless otherwise stated, the electrode slurries are prepared with a silicon : carbon : polymer binder ratio of 60 : 20 : 20 (by mass %). The silicon active material incorporated into the electrode slurry is detailed in the previous section. The carbon material is introduced into the mixture as a conductive additive, which is essential for maintaining the electrical conductive pathways within the electrode. The carbon component is a mixture of several different varieties of carbon: ~10% Super P Carbon Black, ~70% graphite flakes (Alfa Aesar #43209, 325 mesh), and ~20% contribution from the carbon-coating of the SiO active material. This variety of carbon additives is chosen for its nano- to micro-scale size distribution, which is expected to maximize its dispersion throughout the electrode system while forming a strong electrical conductive network due to the bridging between surface coating, carbon nanoparticles (Super P Carbon Black), and micro-sized graphite flakes. For a single batch of electrode slurry, 60mg of silicon monoxide active material (i.e., SiO or SiO-C) and 20mg of carbon additives are first carefully weighed and deposited into an agate mortar, where the powders are ground and mixed over the course of 10 minutes. With the mixture of SiO and carbon additives appearing suitably uniform in particle size (i.e., no noticeable clustering) and color, the powder mixture is transferred to an empty borosilicate glass vial (20mL). With the weight of the glass vial previously set as a baseline, the total weight of the powder mixture is measured. It is important to make note of this weight because some material will unavoidably be loss during the transfer from mortar to vial, however the material ratio must be maintained for the final product. It is assumed that the powder mixture is still ca. 75% SiO and 25% carbon additives. With the weight of the powder mixture recorded, the powder is dry stirred for 1 hour via a magnetic stir bar apparatus. This is an important step in the process as it ensures that the dry particles are uniformly mixed prior to the addition of liquid polymer binder solution into the mixture, where clusters of materials are likely

29 16 to form permanent segregations in the electrode slurry. After the dry stirring process, the glass vial and powder mixture are weighed as baseline, prior to the addition of polymer binder solution. At this stage, the chosen polymer binder solution will be added to the glass vial. In this thesis, several polymer binders are investigated for their effects on the cycling performance of the SiO-based electrodes. These polymer binders include conventional polyvinylidene fluoride (PVdF, Sigma Aldrich #427179), sodium carboxymethyl cellulose (NaCMC, Sigma Aldrich #419303), poly(acrylic acid) (PAA, Sigma Aldrich #181285), and a crosslinked polymer binder of PAA and NaCMC. The crosslinking of PAA:NaCMC follows the process previously reported in the literature [31]. The PAA and PAA:NaCMC binder solutions are prepared by mixing the precursor polymer into water solution (8% polymer by weight), whereas the PVdF binder solution is prepared by mixing the precursor PVdF solid material into a N-Methyl-2-pyrrolidone (NMP, VWR # ) solution (8% polymer by weight). Through the use of a calibrated micropipette, 250mg of solution (i.e., ~20mg of polymer) is extracted and deposited into the glass vial. At this point, the mixture is now referred to as the electrode slurry. The slurry is mixed for hours via the magnetic stirring apparatus. At this point, the electrode slurry is ready for casting. For the casting process, it is imperative that the following procedures be performed in quick succession of each other, as exposing the slurry to air will immediately cause the slurry to begin drying and, in turn, affect its viscosity. The casting process is performed using a Doctor Blade Assembly (Tape Casting Warehouse), which is essentially a polished stainless steel blade that rides directly above the current collector film at a precisely controlled height, via a dual micrometer adjustment system. This height or blade gap is selected based on the desired thickness or mass loading of the cast electrode slurry. The Doctor Blade Assembly is driven by an Automatic Film Coater (MTI Corp., MSK-AFA-II), as shown in Figure 3-1. Prior to casting the electrode slurry, copper foil (MTI, EQ-bccf-9u, 11µm thickness) is secured flat to the casting

30 17 surface by taping its corners to the casting surface. Next, the copper foil is gently, as not to compromise its attachment to the casting surface, cleaned with ethyl alcohol to remove any fine dust particles. With the surface cleaned and any residual ethyl alcohol evaporated, the electrode slurry is deposited onto the copper foil in a small circular area (i.e., a small mound sufficiently thick to be affected by the blade gap). Quickly following the slurry deposition on the copper foil, the automatic casting is enabled at a speed of 30 mm/sec. At this point, no further contact should be made with the casted electrode slurry until it is visibly dry; contact should be avoided to ensure an uninterrupted bonding between the slurry and current collector. Once the electrode slurry is visibly air-dried, it can be gently removed from the casting surface and transferred to a vacuum oven. While transferring from the casting surface to the vacuum oven, it is imperative to keep the electrode film as flat as possible to avoid inducing fractures. With the electrode film in the vacuum oven, air is evacuated and the temperature set to 65 o C and held at this environment for 1 hour. The vacuum oven heat treatment is necessary for thoroughly removing any moisture and oxygen in the electrode system, which can be destructive to the electrode performance in an assembled battery cell. Figure 3-1. Image of the Automatic Film Coater and Doctor Blade casting device.

31 Calendaring and Porosity Control Unless otherwise mentioned, the electrodes evaluated in this thesis are pressed through a process known as calendering prior to their incorporation into a fabricated coin cell. The calendering process is a necessary stage of the electrode fabrication process in commercial applications, as it increases the packing density of electrode in a cylindrical cell. In this work, the calendering process is used as a means to control the porosity of the electrode and thus evaluate the influence of porosity on cycling performance. Prior to the calendering process, the nominal thickness of the electrode sheet prepared as described in the previous section is evaluated by measuring the thickness at 20 different location on the electrode sheet with a digital micrometer. This initial thickness of the electrode sheet allows for an educated guess on the necessary pressing gap needed for a given porosity, according to Figure 3-2. Figure 3-2. Correlation between the deformed thickness of an electrode sheet and the roller gap width.

32 19 A standard rolling press is employed for the calendering of electrodes and is shown in Figure 3-3. Before the actual calendaring process begins, the rollers are cleaned with ethanol to avoid any dust particles coming into contact with the electrode. Once cleaned, the rollers are driven at 10% speed. The roller gap is adjusted based on Figure 3-2 for specific deformed electrode thickness, which in turn can be used to determine the electrode s porosity. The electrode sheet is slowly guided in between the rollers until automatically drawn through, with an emphasis on keeping the electrode sheet flat throughout the process. Figure 3-3. The device used for calendering the electrodes.

33 20 After the calendaring process, individual electrode discs are punched from the electrode sheet using a Precision Disc Cutter (MTI Corp., MSK-T-06). The individual electrode discs have a circular area of 1.13 cm 2. Each electrode disc is measured for thickness is several locations; typically 5 locations within the circular area. With the thickness and mass of the electrode and substrate, as well as the density of the solid material incorporated into the electrode, the porosity of the electrode can be calculated as: P = (1 V s V tot ) 100 = (1 m Si ρ Si + m C ρ C + m PB t A ρ PB ) 100 where P is the porosity of the electrode, m is the mass of the material (Si = silicon, C = carbon, PB = polymer binder), ρ is the density of the material, t is the thickness of the electrode (minus the substrate thickness), and A is the area of the electrode (i.e., 1.13 cm 2 ). 3.4 Cell Fabrication For the assembly of the 2016-type coin cells, the SiO-based electrodes prepared as previously described are transferred to an argon-filled dry glove box (MBraun, Inc.). Immediately before being transferred to the glove box, the electrodes are thoroughly dried in a vacuum oven at 65 o C for 30 minutes. The internal components of the assembled coin cell are shown in Figure 3-4. The coin cell cases are 304 stainless steel with polypropylene O-ring (MTI Corp., EQ-CR2016). The electrode discs are mounted on 304 stainless steel spacers (MTI Corp., EQ-CR20- Spacer304), which ensure the proper positioning and electrical contact of the electrodes with the battery casing. The separator is a 25µm trilayer polypropylene-polyethyerlene-polypropylene

34 21 membrane (MTI Corp., Celgard, EQ-BSF C), which ensures the electrodes are electrically-isolated from each other, while allowing electrolyte to penetrate and thus allow Li-ion diffusion. A polycrystal Li metallic foil (MTI Corp., 0.06mm thick, Lib-LiF-35M) is employed as the counter electrode in the half cell assembly. The SiO and SiO-C electrodes are employed as the working electrodes in the electrode assembly. The electrolyte used for all of the coin cells evaluated in this thesis is composed of 1 mol/l LiPF 6 in a mixture of ethylene carbonate, diethyl carbonate, and dimethyl carbonate (EC : DEC : DMC, 2:1:2 by vol%) and 10 wt% fluoroethylene carbonate (FEC) (Novolyte Technologies, Independence, OH). Once the coin cell components have been assembled as described in Figure 3-4, the assembly is compressed by a Compact Hydraulic Crimping Machine (MTI Corp., MSK-110) at 1000psi. At this stage, the coin cell assembly is complete and the cells are transferred outside of the glove box for electrochemical characterization. Figure 3-4. The internal components of the coin cell assembly and the device used to crimp and seal the coin cells.

35 Electrochemical Characterization The electrochemical performance of the cells investigated in this thesis was evaluated by galvanostatic charge/discharge cycling on an Arbin BT-2000 or Landt Instruments CT2001 Battery Tester. Battery performance is evaluated at room temperature (ca. 25 o C) or under hightemperature conditions (60 o C), as specified. For the high temperature evaluation, coin cells are kept inside a temperature controlled oven (MTI Corp., EQ-DZF-6050-UL) for the duration of the study. The coin cells are evaluated within a voltage range of V versus Li + /Li and at various current densities, as specified based on the theoretical capacity of silicon of 4,200mAh/g- Si and the weight of silicon in the electrode. The specific capacities derived from this electrochemical testing are based on the mass of the active material.

36 23 Chapter 4 The Influence of the Fabrication Parameters 4.1 Introduction In this chapter, the influence of electrode fabrication parameters, such as the active mass loading and the pressing of the electrode, on the cycling performance of the electrode will be investigated. The active mass loading of the electrode is the mass of active material loaded onto a unit area of the electrode (i.e., mg/cm 2 ), and has been found to strongly correlate to variations in gravimetric capacity, capacity retention, and the energy density of the electrode. In this thesis and other literature related to Si electrodes, the omnipresent goal is to achieve the highest active mass loading possible before the onset of cycling instability. In this effort, this chapter will also investigate the influence of the calendering process on the electrode performance. The calendering process is essentially a high pressure pressing of the electrode, which can be used as a method for tuning the thickness and porosity of the electrode laminate following the slurry casting process. In commercial applications, the calendering process is invoked to improve the volumetric capacity of the electrodes; however its effects on Li-alloy electrodes has been only vaguely explored in the literature. In Section 4.2, a literature review on the topics of active mass loading and calendering effects on Si electrode performance is presented. Several high-impact publications on the most successful Si electrodes is first presented to highlight the low active mass loading that are typically reported and overshadowed by impressive cycling performance. Next, research particularly concerned with the impact of varying Si mass loading is presented and its implications discussed. Lastly, several reports are selected on the under-reported topic of calendering effects on Si electrode performance and their implications are discussed. In Section 4.3, the research conducted by this author on the topics of Si mass loading and calendering effects

37 24 on cycling performance is presented and the results discussed. Lastly, in Section 4.4, the implications of the results derived from this research and future work on the topic are discussed. 4.2 Literature Review The primary motivation for achieving higher mass loadings of active material in Li-ion battery electrodes is to maximize the gravimetric and areal lithium capacity of the electrode, and thus offer a more energy dense battery for commercial applications, which is especially important for the future of electric vehicles. In fact, the original motivation of this thesis originates from the Advanced Battery Research (ABR) project sponsored by the Department of Energy (DOE) with a demand for high energy and power density electrodes towards the improvement of Li-ion batteries in hybrid-electric vehicles. An increasing in active material loading translates into a decrease in weight fraction of inactive materials, such as the separator and current collector, and thus increase the overall energy density of the battery system. The challenge with accomplishing a high mass loading (i.e., >1mg/cm 2 ) and areal capacity (i.e., >2mAh/cm 2 ) in Li-ion battery electrodes, especially those of the Si variety, is the mechanical instability of the electrode during electrochemical reaction. During the fabrication of high mass loading electrodes, the casted electrode slurry will become increasingly brittle with higher thickness of coating and macroscale cracking is likely to occur throughout the electrode during the drying process. This was a consistent challenge when developing the high mass loading electrodes presented in this thesis; Figure 4-1 shows an example of the cracking that typically occurs with high mass loading electrodes if not properly handles or has been dried at a very fast rate. High mass loading electrode films will also suffer from a diminishing flexibility, which is problematic for electrode winding used in many commercial battery formats. In regards to the cycling performance of high mass loading electrodes, the electrode will obviously achieve a more desirable gravimetric and

38 25 areal capacity, however at the great expense of rapid capacity fading. This is a consequence of the system losing mechanical integrity as the thickness of electrode laminate increases. As the title of this thesis implies, enabling a high mass loading of an electrode is not achieved by improving a singular aspect of the battery design, but instead will require an optimal combination of system components, such as a mechanically-robust polymer binder and active material. Figure 4-1. An example of the cracking and delamination that commonly occurs when preparing a high mass loading electrode that has not be properly handled or has had its moisture evaporated too quickly. For the majority of the literature focused on Si electrodes, the importance of achieving a high mass loading is overshadowed by the influence of polymer binder or active material morphology. Typically, the active mass loading of Si electrodes in the literature is below 1 mg/cm 2, far from the mass loading necessary for commercial applications [32, 33, 34, 35, 36, 37, 29, 38]. Wu et al. recently reported on perhaps the longest cycling of a Si-based electrode material through the development of 3D continuous-conductive hydrogel network binder (i.e., essentially an electronically conductive binder that conforms to the surfaces of the active material) [34]. With the polymer binder and nanoscale Si active material as the anode, the group

39 26 is able to cycle the lithium half-cell for up to 1,000 cycles at a nominal gravimetric capacity of 1,600 mah/g in a voltage window of V at a relatively high 1 A/g current density. Although the electrode can maintain exceptional capacity and stability up to 1,000 cycles, the active mass loading of the electrode is a very low 0.2 mg/cm 2. At this low mass loading and with the use of nano-sized active material, the electrode is essentially a thin film that will be significantly less influenced by the volume change during Si lithiation/delithiation. However, both the use of nano-sized and low mass loading pose a roadblock for practical application. Chan et al. also report an impressive feat of one of the highest gravimetric capacity achieved by a Si electrode [29]. The group is among the first to develop a silicon nanowire electrode, where polymer binder is unnecessary since the nanowires are directly grown on the current collector substrate, and therefore do not need to be laminated onto the substrate via binder. Since the nanowires are attached directly to the current collector, this type of system benefits from high electrical conductivity that avoids the resistance that typically comes from electrons travelling through a system of semiconductive Si-C randomly-oriented particles. With this advanced active material system, the electrode is able to achieve a capacity of ca. 3,200 mah/g, which can remain stable for up to 10 cycles, cycled within a voltage of V at C/20 rate. Among the issues that present themselves when translating these results to practical application, the low mass loading of ca. 0.5 mg/cm 2 is apparent. The low mass loading definitely plays a large role in the magnitude of gravimetric capacity, as ionic and electronic conductive pathways are shorter in this thin film electrode, and thus a high percentage of active material is likely to be utilized. Liu et al. demonstrate how active mass loading, even with a highly sophisticated pomegranate-shaped active material, plays a critical role in the cycling performance of Si-based electrodes [38]. This study has a similar ambition to this thesis, in that the group addresses the challenges of Si anodes by optimizing several aspects of the battery components, in their case

40 27 primarily related to the active material morphology, as well as deliberate void spaces introduced during synthesis to accommodate Si expansion during Li-alloying. The active material is an impressive system inspired by the morphology of a pomegranate fruit, where nanoscale Si primary particles are encapsulated within a conductive carbon shell; a cluster of core-shell particles are attached to form the hierarchical micro-scale composite material. In spite of this accommodating active material morphology, the group shows that increases in active mass loading will dramatic reduce the cycle life of the electrode. The electrodes evaluated undergo three formation cycles at 0.03 ma/cm 2, followed by a current density of 0.7 ma/cm 2 within a voltage window of V, and an electrode composition of silicon : carbon : binder of 8:1:1 (mass%). At the highest mass loading of 3.12 mg/cm 2, the group can achieve an impressive nominal areal capacity of 3.67 mah/cm 2, which is similar to the areal capacity of commercial Liion cells. However, with the high mass loading comes a dramatic decrease in the number of cycles prior to unstable capacity fading. At the high mass loading of 3.12 mg/cm 2, the electrode has a stable cycling performance of ca. 100 cycles, whereas the 1.20 mg/cm 2 electrode can remain stable for ca. 200 cycles at the cost of a greatly reduced areal capacity of ca. 1.2 mah/cm 2, as shown in Figure 4-2. Unlike the aforementioned studies, several research groups have specifically reported on electrodes with high active mass loadings, and offered insight on how to achieve this goal [39, 30, 40]. Such studies have investigated electrodes which exceed an active mass loading of at least 2.0 mg/cm 2.

41 28 Figure 4-2. Electrochemical performance of high active mass electrodes composed of a novel pomegranateinspired silicon electrodes. The electrodes were first cycled at 0.03 ma/cm 2 for three cycles and 0.7 ma/cm 2 thereafter. Data referenced from Liu et al. [38] Chen and Wang et al. were motivated to improve the areal capacity of Si based electrodes based on the necessity of high gravimetric and volumetric energy density for practical applications, while incorporating a micro-scale active material to avoid the impractical cost that accompanies nanoscale particle synthesis [39]. Again, referring to the approach advocated in this thesis, the group successfully achieves stable cycling performance of a high areal capacity Si electrode through a multi-faceted approach of improving particle morphology and developing a polymer binder with a self-healing chemistry. The group synthesizes Si active material with a variety of sizes ranging from 250nm to 3.5µm through a solution precipitation-floatation selection process. From this size range, the group identifies that the electrode composed primarily of 800nm Si particles has the best combination of achievable capacity and cycling stability, and thus is a candidate for high mass loading tests. With the 800nm particles, this group was able to

42 29 fabricate electrodes with a relatively moderate mass loading of 1 1.6mg/cm 2. Complemented by the self-healing properties of the polymer binder and being cycled at a moderate C/30 or C/10 rate, the electrodes show excellent cycling stability for over 100 cycles. For example, electrodes prepared with a 1.6 mg/cm 2 mass loading could achieve a nominal areal capacity of ca. 4.2 mah/cm 2 and maintain stable cycling performance up to 140 cycles, as shown in Figure 4-3. At a lower mass loading of 0.72 mg/cm 2, the areal capacity significantly decreases to a nominal 1.97 mah/cm 2 but still retains the excellent cycling stability. This research is a clear demonstration of how optimizing several aspects of the battery system is necessary for greatly improving battery performance. Figure 4-3. The influence of active mass loading on the cycling performance of a Si-based electrode composed with a novel self-healing polymer binder developed and reported by Chen et al. [39]

43 30 With a similar ambition, Cui et al. seek to maximize the active mass loading of Si electrodes by developing C-Si core-shell active material through a chemical vapor deposition (CVD) synthesis route, where the carbon nanofibers are encased by an amorphous Si shell [30]. Since the carbon backbone will not undergo significant volume change during electrochemical reaction, it will act to reinforce the Si active material shell and alleviate the stresses generated during lithiation. With this logic, the group not only achieves a relatively stable cycling performance up to 50 cycles, but also can achieve a high areal capacity. When the electrode is cycled within a restricted voltage window of V (opposed to a lower limit of 0.01V) and at a C/15 rate, the high mass loading (~2.4 mg/cm 2 ) C-Si core-shell based electrode is able to achieve a nominal areal capacity of ~4 mah/cm 2. This work demonstrates the benefits of incorporating a stress-accommodating feature (i.e., the carbon-core) into the silicon particle system and how it can enable higher mass loading. However, the practical application of this material suffers from a CVD synthesis technique that is not readily scalable for industrial manufacturing. Additionally, stable cycling performance at a high areal capacity is mainly achieved through a limited lithiation where the full capacity of Si is not utilized. Although their investigation is focused on the cathode side of Li-ion full cells, Zheng et al. highlight the relevance of developing electrodes with high active material loading, among other optimized electrode-level design parameters, for increasing the energy density of battery systems in electric (EV) and plug-in hybrid electric vehicles (PHEV) [40]. The group is also partially motivated by a lack of correlation between active mass loading and cycling performance in the literature, which is more commonly focused on the active material morphology and polymer binder influences. This group correlates the mass loading of LiNi 1/3 Co 1/3 Mn 1/3 O2 (NCM) and LiFePO 4 (LFP) cathode material, proclaimed for their high energy density, to that of cycling performance. The laminate thickness of the electrodes ranges between µm and have a mass loading in the range of ca mg/cm 2. The impact of mass loading is obvious in both

44 31 electrode systems, as shown in Figure 4-4 for NCM. From Figure 4-4 (top), it is clear that higher mass loading will correspond to significant variations in the onset of capacity fading at practical discharge rates (i.e., C-rate > 1C). As shown in Figure 4-4 (middle), due to increase in distance of the electrical conduction pathway, electrical conductivity of the electrode will significantly increase by merely adjusting the laminate thickness by 20µm. Lastly, as shown in Figure 4-4 (bottom), capacity retention will suffer to greater extents as active mass loading increases. For the NCM material, capacity can be retained above 95% for nearly 500 cycles at a laminate thickness of 24µm; however, when the laminate thickness increases to 104µm, capacity retention drops to nearly 70% in ca. 20 cycles, and continues to decrease thereafter. This correlates strongly to the cycling behavior of Si electrodes and highlights to importance of this type of study. Based on the aforementioned studies, it is clear that Si electrodes can achieve a high active mass loading when a multi-faceted approach is used to design the electrode system such that both active and inactive material accommodate the volume expansion of the Si during electrochemical reaction. In this thesis, the influence of electrode pressing, or calendaring, during the fabrication process is investigated for its influence on cycling performance. By pressing the electrodes under immense pressure, several critical design factors are affected, such as particle-toparticle contact and the adhesion between the electrode and current collector. However, other aspects of the electrode design may be negatively impacted, such as reduction in specific surface area, increases in electrode tortuosity, and a hindered Li-ion migration throughout the electrode thickness due to a decrease in porosity.

45 32 Figure 4-4. Correlation between the thickness of NCM electrodes and a) rate capability, b) weight specific impedance, and c) capacity retention developed and reported by Zheng et al. [40] An investigation by Nyugen et al. is among the very few directly focused on studying the influence of electrode porosity and void fraction of Si-based electrodes through the process of

46 33 calendering [41]. The group investigates this correlation in an effort to judge the feasibility of transferring Si electrodes from the lab-scale to pilot-scale. The electrochemical performance of electrodes prepared with nano-sized Si, Super P carbon black, and SB-CMC polymer binder are evaluated within a voltage window of V at a C/5 rate and limited discharge capacity of 1,200 mah/g. From the cycling performance of these electrodes of different porosity ranging from ca %, it is clear that porosity plays an important role in extending the cycle life of the electrode. An increase of porosity from 60% to 70% showed an increase of cycles for the cell evaluated. Another interesting effect of the porosity variation is that the maximum mean discharge potential of the cell decreased with decreasing porosity in order to achieve the 1,200 mah/g limitation. As this indicates that the active material must undergo a deeper extent of lithiation to achieve the capacity limit, these results suggests that porosity has a notable influence on electronic/ionic resistance and the amount of active material that is electrochemically active. This makes intuitive sense, as the conductive bridges formed during electrode slurry casting may be severed or broken during the high pressure calendaring process, as well as imposing limitations on the void space for Si expansion during lithiation that can compromise the mechanical integrity of the electrode during cycling. Of the scarce literature that exists on the influence of electrode fabrication parameters on cycling performance of Li-ion batteries, Wilhelm et al. are among the very few that specifically investigate the influence of calendaring on lithium-alloy negative electrodes with TiSnSb active material [42]. The electrochemical performance of the cells is evaluated within a voltage window of V at a discharge rate of 1C. As shown in Figure 4-5, the volumetric capacity of the electrode is significantly improved from ca. 500 to 1100 mah/cm 3 as the porosity decreases from ca. 65% to 30%, respectively; the volumetric capacity, much like areal capacity, is another measure of improved energy density in the electrode. However, other factors that may not have an intuitive correlation to porosity are significantly affected. For example, the cycle life of the

47 34 electrode can be increased from 60 cycles at 30% porosity to 95 cycles at 65% porosity. The linear decrease in cycle life with porosity is attributed to the reversible breathing of the Si during the lithiation process. Essentially, as the porosity decreases so too does the amount of space available to the active material for volume change, and thus a more restricted environment (low porosity) will suffer from poor mechanical integrity. This trend agrees with Nyugen et al. [41]. The group offers that 45% is a suitable porosity for compromising between cycle life and volumetric capacity. Figure 4-5. Influence of calendaring and the induced porosity changes on the performance of a lithiumalloy negative electrode with TiSnSb active material as developed and reported by Nguyen et al. [41]

48 35 Beattie et al. indirectly provides insight on the role porosity can play in a Si electrode system [43]. The group reverse engineers a theoretical model where a close-packed structured of Si particles are uniformly-spaced on a substrate in multiple layers. Upon lithiation, the silicon particles have exactly sufficient room to expand in volume by 270% and barely touch the neighboring silicon particle; in this fashion, particles can expand without causing a significant displacement of the neighboring particles, which could lead to isolated active material in realistic electrode systems. With this scenario considered, the group backtracks to determine the theoretical optimal material composition of the electrode to enable sufficient spacing for Si expansion. This relates to the porosity of the electrode because, even though the group assumes the active material is completely surrounded by binder, void-space or pores may be substituted into this region without altering the results from the model. This group determines that for multilayered Si particle electrodes, the optimal material composition of Si : carbon additive : binder is 33 : 33 : 33 (i.e., about 1/3 for each component). As any realistic binder system is porous to some extent, a fraction of this electrode composition must be allowed for void space. Indeed, electrochemical testing of the nano-sized silicon active material electrode with the 33 : 33 : 33 composition and charge-discharged at C/7.5 rate, achieves exceptional cycling performance with stable cycling at ca. 2,000 mah/g up to 160 cycles. When the Si active material content is increased to 80%, and the pore space likely reduced in effect, the cycling performance dramatically decreases, as shown in Figure 4-6, even when cycled at a very slow C/26 rate.

49 36 Figure 4-6. The correlation between active material content fraction in an electrode composition on the cycling performance of Si nanoparticle-based electrode developed and reported by Beattie et al. [43] 4.3 Results and Discussion Figure 4-7 compares the cycling performance of electrodes composed of carbon-coated pristine silicon (Si-C) active material with varying degrees of Si mass loading. As described in Chapter 3, the mass loading of the electrode is controlled during the electrode slurry casting process, when the gap of the casting blade is varied to control the thickness of the electrode slurry on the Cu current collector foil. For the Si-C active material investigated here, mass loading ranges from a moderate 1.2 mg/cm 2 to a high 3.6 mg/cm 2, relative to the loadings typically reported in the literature [32, 33, 39, 34, 35, 36, 37, 29, 30, 40]. From the electrochemical performance of the Si-C electrodes shown in Figure 4-7, it is clear that a higher mass loading translates into lower gravimetric capacity in the long-term cycling. Indeed, all of the Si-C electrodes initially have a high specific capacity >2,400mAh/g, about half of the theoretical

50 37 capacity of silicon. However, a dramatic capacity fading is observed from the 1st to 2nd cycle; the Si-C electrodes with 1.2, 1.7, 2.4, and 3.6mg/cm 2 mass loading retain 84, 60, 55, and 48% of their specific capacity, respectively. Astoundingly, a difference of a mere 0.5mg of active material loading translates into a drastic decrease in specific capacity and capacity retention. This same trend is also observed for long-term cycling performance. Between the 10th and 100th cycles, the electrodes with mass loading 1.2, 1.7, 2.4, 3.6mg/cm 2 retain 73, 71, 47, 13% of their capacity. Undoubtedly, the contribution from the silicon active material to the lithium capacity of the electrode has greatly diminished by the 10th cycle. In fact, the 2.4 and 3.6 mg/cm 2 mass loading electrodes appear to have no capacity contribution from the silicon active material after approximately 20 cycles. Instead, the small nominal capacity ~300mAh/g is most likely contributed by the more stable graphite additives present in the electrode composition. An interesting trend is also observed from the perspective of areal capacity, as shown in Figure 4-7d. The highest mass loading electrode of 3.6 mg/cm 2 corresponds to the highest 1st cycle areal capacity of 9.41 mah/cm 2 compared to 6.63, 4.39, and 2.96 mah/cm 2 achieved by the 2.4, 1.7, and 1.2 mg/cm 2 electrodes, respectively. However, this trend is not consistent during the longterm cycling. Instead, the 3.6 mg/cm 2 electrode has the lowest nominal capacity (i.e., average between cycles 20 and 100) of 0.51 mah/cm 2 compared to 1.56, 0.97, 0.77 mah/cm 2 for 2.4, 1.7, and 1.2 mg/cm 2 electrodes, respectively. That is, the electrode with the lowest mass loading will eventually succeed the areal capacity of other electrodes. Overall, the cycling performance in terms of gravimetric and areal capacity is a testament to the deleterious effects of constricting a high concentration of silicon active material in a given electrode area. The significant influence of mass loading on cycling performance is further confirmed through the evaluation of alternative silicon active materials.

51 38 Figure 4-7. Comparison of the a) gravimetric capacity, b) 1 st and 10 th cycle gravimetric capacity, c) 1 st cycle coulombic efficiency, and d) areal capacity of carbon-coated silicon (Si-C) anodes prepared with varying active material mass loadings. The lithium half cells are cycled at 400 ma/g in a voltage window of V.

52 39 Figure 4-8 details the cycling performance of electrodes composed of commercial silicon monoxide (SiO) active material with mass loadings of 1.0 and 2.5 mg/cm 2. From the capacity and voltage profiles, it is clear that mass loading has a similar impact on the magnitude of gravimetric specific capacity and the capacity retention of the SiO electrodes as was observed for the Si-C electrodes. The lower 1.0 mg/cm 2 mass loading electrode achieves a high 1st cycle specific capacity of 2,254 mah/g, most likely because the lithium ions have a shorter distance to diffuse throughout the thickness of the electrode and therefore more active material will have the opportunity to be utilized compared to the higher 2.5mg/cm 2 mass loading electrode, which only reached an initial specific capacity of 1,392mAh/g. Both the high and low mass loading cells suffer a significant capacity fading between the 1st and 10th cycle with the capacity falling to 461 mah/g (33% retention) and 967 mah/g (43% retention), respectively. This is indicative of particle pulverization, electrode cracking, and loss of conductive pathways in the early cycles stemming from the dramatic Si volume expansion. However, capacity retention from cycle 10 to 100 is significantly different between the mass loadings, with the lower mass loading retaining 76% of its capacity and the higher mass loading retaining only 44% of its capacity. The superior short-term and long-term capacity retention in the 1.0 mg/cm 2 electrode suggests that the more spacious electrode environment enables a less destructive effect from particle expansion during cycling, and thus a higher reversible cycling performance can be realized. The cycling performance in terms of areal capacity, as shown in Figure 4-8c, also follows a similar trend as observed for the Si-C electrodes. That is, the highest mass loading of 2.5 mg/cm 2 corresponds to the highest 1st cycle areal capacity of 3.48 mah/cm 2, while the 1.0 mg/cm 2 electrodes only achieves 2.25 mah/cm 2. If these areal capacities could be retained for many cycles, they would be among the highest values reported in the literature, however there is significant capacity fading during the initial 10 cycles. In effect, the nominal areal capacity (i.e., average capacity between cycle 20 and 100) of the 1.0 and 2.5 mg/cm 2 electrode become 0.86 and 0.61 mah/cm 2,

53 40 respectively. That is, the areal capacity of the lower mass loading eventually exceeds that of the higher mass loading. Again, this is most likely attributed to the lower mass loading electrode offering a more spacious environment for SiO particle expansion, and thus less electrode-level cracking will occur. Although a more unstable cycling environment appears to be inevitable with increasing mass loadings, certain active particle surface modifications have been shown to enable moderately stable cycling performance with high mass loadings, the reasons for which will be explored in Chapter 4. Therefore, the effect of mass loading on carbon-coated SiO electrodes is next evaluated. Figure 4-9 details the cycling performance of a carbon-coated silicon monoxide (SiO-C) active material electrode, which can achieve moderately stable cycling performance at high mass loadings above 4.0mg/cm 2. The combination of the phase impurities intrinsically present in SiO systems along with the very high mass loading (i.e., >4 mg/cm 2 ), causing slow diffusion of lithium ions throughout the thickness of the electrode, results in a low initial lithium utilization of the active material, hence the 4.6 and 6.0mg/cm 2 electrodes investigated here have a relatively low initial capacity of 1,221 and 1,215mAh/g. However, at a sufficiently low mass loading and electrode thickness, the SiO-C can achieve a very high initial gravimetric capacity of 2,652 mah/g and by the 10 th cycle the capacity has only faded to 2,368 mah/g. This high capacity can attributed to superior lithium diffusion in the thin film, a less volatile environment from Si expansion, and improvements to the domain structure offered by the heat-treatment described in Chapter 3. At the 11 th cycles, the electrode is exposed to a high-temperature (65 o C) environment to tests its ability to maintain a stable cycling performance in extreme conditions. When the high temperature is induced at the 11 th cycles, the capacity suddenly increases to 2,710 mah/g. Over the course of 65 cycles, the electrode is able to retain ca. 86% of this capacity (i.e., 2,317 mah/g at cycle 65). Although capacity fading remains an issues, it has been significantly improved over

54 41 previously samples, and clearly demonstrates the influence of mass loading on cycling performance. Figure 4-8. Comparison of a) gravimetric capacity, b) voltage profiles, and c) areal capacity of pristine silicon monoxide anodes prepared with 1mg/cm 2 (red) and 2.5 mg/cm 2 (black) mass loading. The lithium half cells are cycled at 400 ma/g in a voltage window of V.

55 42 Nevertheless, by choosing a slow current density of 50mA/g, a relatively stable cycling performance can be achieved even for the 4.0 mg/cm 2 electrode. The initial coulombic efficiency of the high mass loading electrodes is ca. 73%. Unfortunately, the 1 st cycle of the 1.0 mg/cm 2 electrode fails to achieve a meaningful capacity, perhaps because the electrode was not given a sufficient opportunity to absorb electrolyte prior to cycling, and thus a true CE cannot be calculated. The effects of mass loading, once again, are obvious in the long-term capacity retention of the electrodes. While the 4.6 mg/cm 2 electrode can retain 57% (431 mah/g) of its capacity from cycle 10 to 100, the 6.0mg/cm 2 suffers greatly and can only retain 14% (110mAh/g-Si) within the same cycle range. Again, this capacity is lower than that of graphite (372 mah/g) and most likely suggests that the active silicon material is no longer contributing to the lithium capacity of the cell. The high mass loadings of both electrodes translates into silicon active material that is packed in a very constricted environment and, in spite of the low current density, the negative effects from particle volume expansion rapidly decays the amount of active material available for lithiation. As explained in Chapter 2, this high capacity fading has been shown to be from a loss of conductive networks, unstable SEI growth, and delamination from the current collectors, among a variety of other issues all sourced from silicon particle expansion. Once again, the areal capacity of the SiO-C electrodes follows the same trend observed for the Si-C and SiO electrodes. Although the 4.6 and 6.0 mg/cm 2 electrodes have a very high 1st cycle areal capacity of 5.62 and 7.29 mah/cm 2, respectively, the thickness of the electrodes leads to a rapid capacity fading over the course of 90 cycles. The 4.6 and 6.0 mg/cm 2 electrode areal capacity diminishes to 1.85 and 0.49 mah/cm 2, respectively, by the 90th cycle. The cycling performance of the 1.0 mg/cm 2 electrode is significantly more stable compared to the higher mass loadings. At the 2nd cycle, the 1.0 mg/cm 2 electrode achieves an areal capacity of ca mah/cm 2, and is further increased to ca mah/cm 2 during high-temperature testing beginning at the 11th cycle. Even at high temperature conditions, the 1.0 mg/cm 2 electrode can

56 43 maintain a nominal areal capacity of 1.50 mah/cm2 for up to 65 cycles. This value of areal capacity is at least competitive with the values reported in the literature, and with further optimization of the electrode system, can potentially reach the threshold for practical application. Figure 4-9. Comparison of the a) gravimetric capacity and b) areal capacity of carbon-coated silicon monoxide anodes prepared with high mass loading of 4.6 and 6.0 mg/cm 2. The lithium half cells are cycled at 100 ma/g in a voltage window of V.

57 44 During the fabrication of battery electrodes, one may also control the porosity of electrodes through a process known as calendaring, which is more thoroughly described in the Chapter 3. Through this calendaring process, which involves passing the electrode through a gap between two immense metal rollers, the solid materials of the electrode are condensed and the electrode porosity reduced. This is very relevant to industrial manufacturing of battery electrodes, as this process will increase the volumetric energy density of the cell. The SiO-C electrodes with a porosity of 30-37, 45-50, and 60% have SiO-C mass loadings of 8.5, 5.3, and 6.8 mg/cm 2. In spite of this wide range of active mass loadings, the calendering process seems to influence the cycling performance of the electrodes in a way that results in nearly identical gravimetric capacities throughout the course of 50 cycles. In fact, the cycling performance at all porosities appears very similar to that of the 4.0 mg/cm2 SiO-C electrode, shown in Figure The 1st cycle gravimetric discharge capacities of the electrodes are 1126, 1354, and 1144 mah/g for the 30-37%, 45-50%, and 60% porosity electrodes, respectively, with similar initial coulombic efficiencies of 71-77%. Additionally, the capacity retention of the electrodes are all very similar, with ca. 76% capacity retention between the 1st and 2nd cycle, followed by ca. 65% capacity retention between the 5th and 50th cycle. This is a remarkable finding, as the mass loading of the electrode was essentially able to increase from 5.3 to 8.5 mg/cm 2 without a significant change in the magnitude of gravimetric capacity or its retention. This clearly shows the importance of adopting the calendaring technique for industrial applications, as it is an effective means for tuning the areal capacity of the electrode without significantly altering its cycling performance, as shown in Figure 4-10b. From the perspective of areal capacity, the influence of the active mass loading is obvious, as a higher mass loading translates into a higher areal capacity. At the 1st cycle, electrodes with mass loading of 5.3, 6.8, and 8.5 mg/cm 2 can achieve an areal capacity of

58 , 6.81, 7.55 mah/cm 2. As mentioned before, the capacity fading of the electrodes is nearly identical. Figure Comparison of the a) gravimetric capacity and b) areal capacity of carbon-coated silicon monoxide anodes prepared with varying electrode porosity and active mass loadings. The lithium half cells are cycled at 50mA/g in a voltage window of V.

59 Conclusions From the investigation on the impact of active mass loading on cycling performance, it is clear that an increase in active mass loading will inevitably lead to a decrease in gravimetric capacity and capacity retention; however the onset of this deleterious behavior can be prolonged or improved depending on the type of active material chemistry and morphology employed. For instance, when the active mass loading of Si-C and SiO electrodes is increased to values slightly over 1 mg/cm 2, the gravimetric and areal capacity magnitude drop to values that indicate the majority of Si in the electrode is no longer contributing to the lithium storage. That is, gravimetric capacity falls to values below the theoretical gravimetric capacity of carbon (i.e., 372 mah/g). The SiO-C active material shows very promising performance, however, even at active mass loadings reaching 4.6 mg/cm 2. Although the gravimetric capacity of the electrode is reduced by ca. 1,600 mah/g when increasing the active mass loading from 1.0 mg/cm 2 to 4.6 mg/cm 2, the long-term stability of the electrode remains relatively similar. From the perspective of areal capacity, SiO-C electrodes at 4.6 and 6.0 mg/cm 2 mass loadings can retain an areal capacity above 3 mah/cm 2 for up to 50 cycles, which is a fair compromise in regards to the reduction in gravimetric capacity. In fact, the areal capacity achieved by the SiO-C investigated in this thesis represents a very competitive electrode performance (>3mAh/cm 2 ) in comparison to the sophisticated electrode systems presented in the literature review. Whereas some groups employed sophisticated and costly synthesis routes to develop active material morphologies and special polymer binder properties, the SiO-C presented here is the product of commercially available SiO particles and a simple heat-treatment process for carbon-coating. Additionally, the areal capacity of the SiO-C electrodes can be further improved through the calendering process. Electrodes with an active mass loading up to 8.5 mg/cm 2 are pressed through the calendering process and achieve similar gravimetric capacities as the non-pressed 4.6 mg/cm 2 mass loading

60 47 electrode. This is remarkable considering the areal capacity is increased by ca. 1.3 mah/cm 2 (e.g., an increase from 3.21 to 4.56 mah/cm 2 at cycle 20 between the 5.3 and 85 mg/cm 2 loading electrodes, respectively), without a significant sacrifice in capacity retention. Based on these promising results, it is clear that the SiO-C electrode system should serve as a platform for future work on Si electrodes. Although further work needs to done on improving the capacity retention of the electrodes, it is clear that these electrodes are capable of achieving very high areal capacity relative to what is commonly reported in the literature. Therefore, the SiO-C electrodes will remain the focus in the upcoming chapters, where the role of oxide content, carbon-coating, and polymer binder are investigated.

61 48 Chapter 5 The Influence of the Active Material Chemistry 5.1 Introduction Based on the superior performance of the SiO-C showcased in the Chapter 4, the objective of this chapter is to identify and discuss why a carbon-coated silicon monoxide active material is able to achieve superior electrochemical performance relative to its Si-C and SiO counterparts. As will be discussed, the nano- and micro-structure of the silicon monoxide active material offers a buffer region that can alleviate the stress and strain experienced by the overall SiO particle, and in this way allow for a more stable cycling performance compared to other pristine Si electrodes, at the cost of a capacity significantly lower than the theoretical capacity of Si. Similarly, the carbon-coating can offer this same stress-alleviating function, while also providing a strong electronic conduction pathway throughout the electrode. In Section 5.2, a literature review on the successful application and fundamental structure of SiO active material is first presented and discussed. Following this, a literature review on the role and benefits of a carbon-coating on the surface of the active material is presented and discussed. In Section 5.3, the electrochemical performance of carbon-coated silicon monoxide electrode developed and analyzed by the author will be discussed and compared to the non-coated counterpart. In Section 5.4, the conclusions from the experimental results and suggested future modifications to the SiO- C electrode will discussed.

62 Literature Review Silicon monoxide (SiO) has recently risen as a promising alternative active material for Si-based electrodes in Li-ion batteries. The microstructure of the SiO has been found to have two distinct nano-domains, Si (Si 0 ) and SiO 2 (Si 4+ ) [44]. The inactive SiO 2 layer has been found to encase the core Si domain, along with any intermediates that form during the lithiation process, such as Li 2 O and lithium silicate [45, 46], and can act a buffer layer to restrain particle expansion and alleviate the internal stresses generated during the lithiation of the active phase [47, 45, 46, 48]. Perhaps due to the inactive surface layer, most SiO electrodes that have been investigated in the literature have achieved capacities of only mah/g, which corresponds to ca % volume expansion on the particle level. Although the capacity is much lower than the theoretical capacity of pristine Si (4,200 mah/g-si), the smaller volume change can lead to a more stable cycling performance. Issues with the SiO electrode are with the electronic conductivity, which suffers due to the insulating SiO 2 phase on the surface of the particle. Nevertheless, here several studies are reviewed on the advancements of SiO electrodes and their derivatives. Zhao et al. highlight the successful application of a SiO electrode, as well as provide fundamental insight on the chemical composition of the constituent SiO particles [49], as shown in Figure 5-1. Here the group details how nanoscale silicon domain are surrounded by an inert silicon dioxide matrix. Impressively, the group develops an electrode with a high silicon-tobinder ratio of 98:2 (mass %), one of the highest reported in the literature. To overcome issues of electronic conductivity intrinsic to SiO electrodes, the group employs a carbon coating on the surface of the SiO particles. Additionally, the group is able to replace the common conductive additives, such as carbon black, with a conductive polymer binder (PFM), which enables a higher loading of active material.

63 50 Figure 5-1. Schematic diagram of the multi-domain structure and porous carbon surface coating of a SiObased active material developed and reported by Zhao et al. [49] The SiO-based electrode exhibits exceptional cycling performance, achieving a stable and reversible capacity of ca. 1,000 mah/g for up to 400 cycles when cycled at a C/10 rate in a voltage range of V, as shown in Figure 5-2. Although the sophisticated conductive polymer binder undoubtedly plays a large role in its success, previous testing with this binder in a pristine Si electrode showed a significantly degrade in cycling performance. Figure 5-2. The electrochemical performance of SiO-based electrodes incorporated with an electronicallyconductive type PFM polymer binder developed and reported by Zhao et al. [49]

64 51 Motivated by the improved mechanical integrity offered by the SiO 2 buffer layer present in SiO active material, Doh et al. compare the performance of SiO electrodes to that of Si:SiO mixture electrodes with different electrode compositions: Sample A) Si:SiO:C = 3:1:4, Sample B) Si:SiO:C = 2:2:4, Sample C) SiO:C = 1:1 (weight %) [50]. The electrodes employ a graphite carbon additive (<37microns) and a PVdF polymer binder (15 mass%). The cells are cycled within a voltage range of V (vs. Li + /Li) at a 0.1 C-Rate (based on the theoretical capacity of graphite, 372 mah/g). As shown in Figure 5-3, the samples with pristine Si in their compositions (i.e., Sample A and Sample B) achieve the highest maximum gravimetric capacities during the earlier stages of cycling: Sample A achieves 1,400 mah/g (4 th cycle) and Sample B achieves 1,450 mah/g (8 th cycle). The electrode devoid of pristine silicon (i.e., Sample C) achieves a smaller gravimetric capacity of 870 mah/g (4 th cycle). Although the SiO does not achieve the highest capacity, the benefits of the multi-domain SiO chemistry is clearly evident from the perspective of long-term cycling performance. The capacity retention at the 50 th cycle compared to the cycle with maximum capacity is 34% and 52% for the samples containing pristine Si (Sample A and B); that is, the capacity retention is decreasing as the amount of pristine Si increases relative to other electrode materials. However, the SiO shows an excellent capacity retention of 70% from its maximum capacity to the 50 th cycle. Although the cycling performance concludes at 50 cycles, it appears that the capacities of the samples with pristine Si (i.e., Sample A and B) would soon drop below the capacity of SiO sample (i.e., Sample C). The group attributes this superior performance to the improved mechanical integrity that is expected from the multi-domain SiO chemistry, with a layer of SiO 2 and Li 2 O intermediate surrounding the active silicon domain, and thus internal stresses may be more easily accommodated.

65 52 Figure 5-3. Comparison of the electrochemical performance of pure SiO and mixed Si-SiO electrodes developed and reported by Doh et al. [50] Park et al. [8] investigate further into the role of the multi-domain structure of the SiO material, being thermodynamically unstable at all temperatures and thus consistently separated into two distinct domains, one of Si (Si 0 ) and another of SiO 2 (Si 4+ ) [44]. To study how the presence of these domains influences the cycling performance of SiO-based electrodes, the group manually disproportionate the SiO into the two domains to different extents through a high temperature heat-treatment at 800, 900, 1000, 1100, and 1200 o C. These disproportionated-sio (d- SiO) samples were mixed with graphite and PVdF polymer binder to compose the electrode under investigation. Through electrochemical cycling in a voltage window of V and 100 ma/g current density, the group evaluates the electrochemical performance of the electrodes with the various heat-treatment temperatures, as shown in Figure 5-4. Although the pure SiO sample (i.e., no heat treatment) initially achieves the highest discharge capacity of 2,216 mah/g, the group finds that the SiO treated at 1000 o C exhibits the best long-term capacity retention and attributes this superior performance to the nano-domains of pristine silicon being more evenly distributed and sized throughout the Si x O matrix. However, a higher temperature 1,200 mah/g

66 53 leads to the poorest cycling performance, attributed to the wide spread growth of the SiO 2 domain which effectively isolates the inner Si core from being fully lithiated. The impact of this coreshell structure and its evolution with heat-treatment, therefore, can have a substantial effect on cycling performance and needs to be further explored. Figure 5-4. A comparison of the electrochemical performance of SiO-based electrodes disproportionated at different temperatures prior to electrode fabrication, developed and reported by Park et al. [44]

67 54 Since the most favorable cycling performance of the electrode investigated in Chapter 4 of this thesis was achieved by a carbon-coated SiO active material, here the role of a carboncoating is explored and literature on the topic is reviewed. The application of a carbon-coating on the surface of Si active material has been found to have several significant benefits. The carbon layer on the surface of the active material has two primary functions: 1) to impart enhanced electronic conductivity to the electrode system and 2) alleviate the stresses generated during active material lithiation [51, 52]. However there are a variety of secondary functions that the carbon-coating may fulfill. For instance, it is possible that the carbon-coating will react more favorably with the electrolyte and enable the formation of a more stable SEI layer [53]. Furthermore, as the expansion of particles during lithiation may cause electrochemical sintering (i.e., a type of fusing of solid particles), the carbon-coating may also prevent this from occurring during cycling [7]. Kim et al. report on several distinct stages of the electrochemical reaction during the charge-discharge of a carbon-coated SiO-based electrode [54]. The carbon-coating of the SiO is accomplished through the common CVD method and the electrochemical performance of the electrodes is evaluated within a voltage range of V (deep discharge) and current density of 600 ma/g. From the cycling performance, shown in Figure 5-5, it is clear that the carbon-coating enables a higher maximum gravimetric capacity (ca. 1,500 mah/g) than was previously reported for non-coated SiO electrodes; however, the carbon-coating does not completely resolve the issue of capacity fading prevalent in Si electrode systems. After the 15 th cycle, the capacity fading become significant and capacity falls to ca. 550 mah/g by the 100 th cycle. Although the group reports having incorporated a polyimide binder, one of the more effective binders for Si electrode systems [55], there is no exact information on the mass loading

68 55 of the cells or the extent of pressing (i.e., calendaring), factors which can greatly influence cycling performance, as was shown in previous Chapter 4 of this thesis. Figure 5-5. The electrochemical cycling of a carbon-coated SiO electrode developed and reported by Kim et al. [54] Ng et al. report on a unique spray-pyrolysis technique, which is advocated as a versatile and industrially-oriented electrode fabrication method, for preparing carbon-coated Si nanocomposite electrodes [56]. The group promotes the necessity of buffer layers in Si particle systems, such as the carbon-coating, which can accommodate the stresses and strains induced in the active material during lithiation/delithiation. In this study, electrodes are prepared by combining the spray-pyrolyzed carbon-coated Si composite with carbon black additives and a PVdF polymer binder. The electrochemical performance of the cells are evaluated within a voltage window of V and a current density of 100 ma/g. From the voltage diagrams of both the pristine Si electrode and the carbon-coated Si nanocomposite electrode, shown in

69 56 Figure 5-6, it is clear that the carbon-coating plays a significant role in the capacity retention over the course of 50 cycles. Although the pristine Si electrode initially has a discharge capacity of 3,474 mah/g, the capacity rapidly fades to 47 mah/g (i.e., cell death) by 20 cycles. The carboncoated Si electrode, however, initially achieves a 2,600 mah/g discharge capacity and only fades to 1,489 mah/g after 20 cycles. Although this cycling performance is relatively poor with respect to practical performance and capacity retention, it demonstrates the impact of carbon-coating and suggests that it may be a necessary component for the realization of commercialized Si electrodes. Unfortunately, the group uses a PVdF polymer binder, which notoriously offers inadequate chemical bonding at the electrode-binder and electrode-substrate interface and will definitely result in a poor electrode performance. In a sense, this is a testament that a multifaceted approach that incorporates several complementary components must be used to realize good performance. Figure 5-6. A comparison of the electrochemical performance of a pristine nanocrystalline Si electrode (top) and a carbon-coated Si nanocomposite electrode (bottom) developed and reported by Ng et al. [56]

70 57 As mentioned earlier, the process of applying the carbon-coating to the silicon surface is typically accomplished through a high-temperature chemical vapor deposition (CVD) technique, where carbonaceous gases, such as acetylene or propene, are thermally decomposed followed by carbon deposition on the silicon surface. However, when a SiO-based active material is employed, this high-temperature coating method will affect the extent of oxidation of the silicon, as shown by Park et al. [44]; therefore the nano/micro-structure of the SiO may form some intermediate silicate (i.e., SiO x with 0 < x < 2). The extent of oxidation will undoubtedly have an impact on cycling performance, since 1) SiO 2 structure is insulating and inactive, 2) SiO has impurities but can offer a more stable cycling performance, and 3) Si has a high capacity but a rapid capacity fading. The impact of CVD temperature is investigated by Yi et al., who have developed a micro-sized Si-C composite material composed of nanoscale Si building blocks interconnected with carbon from the CVD method [51]. Indeed, from an X-ray photoelectron spectroscopy (XPS) and Raman spectrum analysis, the group shows that a competing mechanism occurs that affects that extent of oxidation: at high temperature the SiO x will decrease while the disproportionation of SiO will cause an increase in SiO 2. Therefore, a direct correlation between the extent of oxidation and carbon deposition temperature becomes complex. However, the impact on cycling performance and electrical impedance is unmistakable. From the cycling performance, shown in Figure 5-7, it is clear that the active material coated at the higher temperature, 800 o C, can achieve a higher gravimetric capacity and 1 st cycle coulombic efficiency, as well as a superior rate performance (i.e., the ability to maintain a higher capacity at higher C- rates) compared to the active material coated at 600 o C. Additionally, based on the smaller semicircle in the EIS plot shown in Figure 5-7, the active material coated at 800 o C has a lower charge transfer resistance. The superior electrochemical performance of the sample prepared at 800 o C was attributed to both the quality of silicon and carbon-coating. At higher temperatures,

71 58 the impurities present in the silicate active material will be lessened (i.e., the extent of oxidation will decrease) and thus less intermediate Li phases (e.g., Li 2 O and Li 4 SiO 4 ) will be formed and detract from the capacity. Concurrently, a higher temperature will lead to higher quality graphitic coating, which will benefit the formation of SEI and the electronic conductivity. Figure 5-7. A comparison in the electrochemical rate performance (left) and electrical impedance (right) of Si-based electrodes with carbon-coatings deposited via a chemical vapor deposition method at different temperature, developed and reported by Yi et al. [51] 5.3 Results and Discussion Figure 5-8 details the cycling performance of electrodes composed of non-coated (SiO) and coated SiO (SiO-C) active materials, along with a comparison to a conventional graphite anode. The electrodes prepared with these active materials all have identical material compositions conductive additives (Super P nanoscale carbon) and polyacrylic acid (PAA) polymer binder as well as having similar active mass loadings in the range mg/cm 2. In this sense, the true impact of active material type can be evaluated without composition or loading influencing the correlation. From Figure 5-8, both the SiO and SiO-C electrodes achieve

72 59 high initial gravimetric capacities of 2254 and 1786 mah/g for SiO and SiO-C electrodes, respectively. In regards to the initial coulombic efficiency, SiO-C achieves a relatively low CE of 65.8%, compared to the 76.8 reached by the SiO electrode. However, the SiO-C electrode clearly has a superior reversible capacity over the course of cycling compared to its non-coated counterpart. The impact of the carbon-coating is clearly demonstrated by the loss of capacity from the 1 st cycle capacity and the subsequent cycle. For the SiO-C electrode, capacity retention between the 1 st and 2 nd cycle is approximately 76%. On the other hand, the non-coated silicon monoxide, SiO, electrode only has a capacity retention of 54%. This dramatic difference in the short-term cycling behavior of the silicon electrodes suggests that the carbon-coating is essential for accommodating the intense stresses generated within the electrode upon Si lithiation, whereas the non-coated SiO system obviously suffers from volatile loss of storage capacity during the first lithiation/delithiation cycle, which cannot be recovered in later cycles. To further demonstrate the positive influence of the carbon-coating on cycling performance, the SiO-C electrode is exposed to a more demanding cycling regime during its long-term cycling in comparison to the pristine SiO. Beginning at the 11th cycle, the SiO-C is exposed to a relatively high temperature (60 o C) environment, which typically bestows a temporary boast in specific capacity at the cost of a more rapid capacity fading. Remarkably, the increase in temperature has a very minimal impact on the capacity fading compared to room temperature cycling (i.e., cycles 1-10). Furthermore, the capacity retention between the 11 th and 40 th cycles of the SiO-C is unexpectedly higher (90%) than that of the pristine SiO (87%), which is cycled at the less demanding room temperature environment. This is extraordinary considering that the SiO-C can maintain a stable cycling performance at high temperature with a nominal gravimetric capacity of 1,447mAh/g, while the pristine SiO suffers a dramatic loss in capacity in the early stages of cycling, eventually only having a nominal capacity of 890mAh/g. Clearly the

73 60 carbon-coating on the surface of the silicon active material plays a vital role in the full utilization of lithium storage capacity and its maintenance over the course of many charge-discharge cycles. As previously discussed, the carbon surface layers can alleviate the stresses generated during Si expansion/contraction, thus reducing the likelihood of particle pulverization and disconnection from the conductive network. Figure 5-8. Comparison of the cycling performance of pristine silicon monoxide (SiO) and carbon-coated silicon monoxide (SiO-C) anodes prepared at similar mass loading of ca. 1 mg/cm 2 and material composition ratio of conductive additives and PAA polymer binder. The lithium half cells are cycled at 400mA/g within a voltage window of V. After cycle 10, the temperature of the SiO-C environment is increased from room temperature to 60 o C.

74 Conclusions Compared to the non-coated SiO electrode, the SiO-C electrode is able to achieve a much higher gravimetric capacity without a sacrifice in capacity retention, even at higher current density and temperature conditions. For instance, between the 11th and 40th cycle, the SiO-C achieves a nominal gravimetric capacity of 2,558 mah/g, while the SiO electrode only reaches a nominal gravimetric capacity of 890 mah/g; however the SiO-C electrode has a capacity retention of 91% while the SiO electrode has a capacity retention of 87%. Keeping in mind that the SiO-C was being cycled in a higher temperature environment, which typically causes a faster capacity fading, these results demonstrate the robust mechanical and electrochemical properties bestowed by the silicate and carbon surface structure of the active material. Furthermore, even when the current density is increased up to 800 ma/g, the capacity retention of SiO-C is retained at ca. 91%. Compared to the literature on SiO-based electrodes, the electrochemical performance of the SiO-C electrode presented here is among the best reported, at least in regards to the early stages of electrochemical cycling [50, 44, 54]. The most direct comparison in the literature is to Zhao et al. [49], who report on a carbon-coated SiO-based electrode with a self-healing polymer binder; although the gravimetric capacity achieved by the SiO-C electrodes reported here are significantly higher than those reported by Zhao et al. when tested at similar cycling parameters, further testing needs to be conducted to determine the life span of the SiO-C electrodes reported here. From the literature review on this type of active material, this successful electrochemical performance can be attributed to several factors. First, the multi-domain structure of the SiO material (i.e., a division of Si and SiO 2 ) provides a unique silicate buffer layer surrounding the Si active domains [44, 46], and thus can assist in alleviating the stress and strain that typically causes particle pulverization and rapid capacity fading in the early stages of the cycling. Secondly, the carbon-coating further fulfills this stress-alleviating role, while enhancing the

75 62 electronic conductivity of the overall electrode system. Lastly, the heat treatment used for the carbon deposition process has previously been shown to improve the cycling performance of the SiO electrodes due to a reduction of oxidized domains and increase in active domains [44]. In the next chapter of this thesis, the influence of the polymer binder on the SiO-C electrode performance is evaluated to determine the binder which can offer the most mechanically and chemically robust system, which will then conclude this comprehensive and multi-faceted investigation.

76 63 Chapter 6 The Influence of the Non-Active Polymer Binder 6.1 Introduction The polymer binder of battery electrodes is conceptualized as the "glue" that holds the active materials to the current collector and maintains the electronic and ionic conductive network among constituent particles. These functions of the polymer binder are indispensable for accommodating the repeated large volume changes of Si active material, and thus the selection of binder directly corresponds to the electrochemical performance of the battery [57, 58]. In commercial LIBs composed with graphitic anodes, the polymer binder employed is typically polyvinylidene fluoride (PVdF) [59, 60, 61]. Based on the relatively low volume expansion of carbon particles (ca. 10%) during the lithiation/delithiation process and their inherent electrical conductivity, the relatively weak binding and electrically-insulating PVdF polymer binder continues to be sufficient in commercial applications. For Si electrodes, however, the application of PVdF as a binder has been consistently shown to result in rapid capacity fading, unstable SEI formation, and delamination from the current collector [62]. In short, the chemical bonding offered by PVdF is insufficient to maintain a uniform particle distribution and conductive network when exposed to repeated large volume changes during electrochemical cycling. With a focus on improving mechanical integrity and Si surface-layer chemical bonding, a variety of polymer binder competitors have been showcased in the literature: CMC [63, 64], polyacrylates [62, 65], polysaccharides [66, 67], conductive polymers [68, 69], and self-healing polymers [32]. Although the cycling performance is still far less than ideal, these newly-developed binders have significantly mitigated the severe capacity fading and fracturing within Si-based electrodes. In Section 6.2, a literature review on polymer binders that have proven to be successful in the Si

77 64 electrode system is first presented. Of the identified candidates for suitable polymer binders, these will be employed towards the SiO-C active material discussed in Chapter 5 and the result will presented and discussed in Section 6.3. Based on the results acquired from the electrochemical characterization of each type of polymer binder system, Section 6.4 will present several conclusions and future outlooks. 6.2 Literature Review In the majority of commercial Li-ion batteries - those composed of a graphitic anode and transition metal oxide cathode - the most common polymer binder employed is the non-reactive fluoropolymer polyvinylidene fluoride (PVdF). Although the PVdF polymer binder offers a flexible and compliant environment for the small expansion of carbon active material, these properties have proven to be insufficient for Si electrodes [70]. This is primarily due to the fact that PVdF lacks the necessary functional groups to offer a strong bond to the SiO 2 surface layer of Si active material; instead the PVdF binder is only capable of bonding through weak van der Waal forces. With the intense volume expansion/contractions of the Si particles, these weak bonds will be severed and, in effect, broken conductive pathways, delamination from the current collector, unstable SEI formation, and electrode-scale fracturing will occur. Based on this inadequate performance of the PVdF binder, much interest has been taken to develop polymer binders that offer more in the way of stronger chemical bonding (i.e., hydrogen or covalent bonding) and more compliant mechanical properties. Carboxymethyl cellulose (CMC) and its derivatives (e.g., NaCMC) were among the earliest contenders for a Si electrode polymer binder, based on improved mechanical adhesion and chemical bonding properties [71, 58, 43]. Contrasting the non-functional PVdF polymer binder, CMC offers hydrogen bonding to the surface silanol groups of Si active material through

78 65 the carboxyl functional groups present on its polymer chains. This type of bonding of the carboxyl functional group to an oxidized surface has also led to an improved adhesion to the current collector, where a copper oxide layer is available for bonding. Furthermore, this bonding can help uniformly disperse the Si active materials during the casting process (i.e., casting electrode slurry on current collector foil) and thus avoid aggregation of active material, which could in turn lead to isolated active material from conductive networks. The CMC polymer binder has also been popularized due to its ability to form a solution in environmentally-friendly processing solvents, such as water, which makes it more suitable for large-scale electrode manufacturing. Originally the success of the CMC binder was perplexing since it was identified to be a brittle material [70], and thus could not offer the flexibility thought necessary for accommodating the intense Si expansion during lithiation. However, Hochgatterer at el. were among the first to reveal that the CMC chemical structure was primarily responsible for its successful application in the Si electrodes, opposed to its mechanical properties [72]. The influence of the carboxyl functional groups on the electrochemical performance of Si electrodes is investigated by comparing Na-CMC to other cellulose-based binders with cyanoethyl or terminal alcoholic functionalities substituted for the carboxyl functional group; all of which were compared to a PVdF-based binder, as shown in Figure 6-1. From the Figure below, it is clear the Na-CMC can enable both higher gravimetric capacity and a more stable cycling performance, compared to all other binders. The group offers further insight on the chemical interaction between the Si active material and Na-CMC binder through attenuated total reflection FTIR (ATR-FTIR) characterization on a partially-hydrolyzed SiO 2 film, which has been spray-coated with a Na-CMC solution. From a comparison of the IR spectra for the pristine Si/SiO2 film and the Na-CMC-coated sample, an emerging peak at 1629 cm -1 demonstrates the formation of an ester-like R1-CO2-R2 bond. From the groups interpretation, this bond is formed from an interaction when the free hydroxide groups on the Si surface are exposed to the terminal

79 66 carboxylic acid sites of the NaCMC. From this interaction, a condensation reaction occurs and a strong covalent bond is said to be formed. In later investigations, it was found that this bond type may actual be a hydrogen bond or covalent bond type depending on the ph level of the NaCMC processing solvent [73]. Figure 6-1. A comparison of the cycling performance of electrodes employing CMC and PVdF derivative binders as developed and reported by Hochgatterer et al. [72] Ouatani et al. provide valuable insight on the influence of the CMC binder on the formation of a stable SEI layer [74]. Although the group investigates the CMC polymer binder applied to a pristine graphite anode, opposed to the more relevant Si-based anode, this insight can be extended to electrode systems that incorporate carbon-coated active material, such as the SiO- C material investigated in this thesis. The formation of the SEI layer throughout the chargedischarge process is characterized through XPS spectra of a graphite anode incorporating an styrene-butadiene rubber (SBR)/CMC (anode), as shown in Figure 6-2. From the evolution of the main peak at ev in the XPS shown in the Figure below, the group shows that at 3.8 V the

80 67 graphite became nearly undetectable, suggesting an SEI with a thickness equivalent to the analysis depth of XPS (5-10nm) had been formed on the surface of the graphite active material. Furthermore, since the C 1s spectra does not show appreciable differences between 4.2 V during charge and 2.7 V at discharge, this indicates that the SEI layer is, indeed, very stable. Other groups have shown a variety of methods to further modify and improve the application of the CMC polymer binder, such as varying the number of carboxyl functional groups available along its chain length (degree of substitution, DS) or by optimizing its molecular weight. Both of these attributes have been shown to significantly impact the cycling performance of silicon electrodes [72, 64]. Figure 6-2. Electrochemical cycling regime and associated XPS spectra for graphite/cmc anode and LiCoO2 full cell system as developed and reported by Outani et al. [74]

81 68 Much like the multi-faceted approach used to improve silicon electrodes in this study, Bridel et al. endeavor to optimize the CMC polymer binder by selecting the optimal molecular weight and degree of substitution derived from previous studies, and develop a Si/CMC electrode with an optimized cycling performance [75]. Through this strategy, the group accomplishes an impressive cycling performance of the Si electrode, as shown in Figure 6-3. Based on these previous investigations, CMC is among the polymer binders that may lead to a suitable cycling performance and is thus chosen for investigation in this thesis. Figure 6-3. The influence of the molecular weight of the CMC polymer binder, at 250,000 (grey squares) or 700,000 g/mol (black diamonds) on the cycling performance of Si-based electrodes developed and reported by Bridel et al. [75] In addition to the widely successful CMC polymer binder, the polyacrylates class of polymer binders has also been proven to accommodate the extreme volume changes of the Si electrode [65, 62, 76, 77, 78]. Similar to CMC polymers, the polyacrylate polymers have a high availability of carboxylic functional groups for bonding to the oxidized Si surface. Komaba et al.

82 69 were among the first to investigate polyacrylate binders, such as polyvinyl acetate (PVA) and polyacrylic acid (PAA), applied to a SiO electrode [77, 79, 80, 65]. This group characterizes a variety of properties of the PAA polymer binder - adhesion strength, electrolyte uptake, XPS spectra, impedance, and influence on cycling performance - all of which clearly demonstrate PAA is well suited for the mechanical stress generated by the SiO electrode. The adhesion strength to the current collector is measured at 2.3 N/cm, significantly higher than PVdF (0.3 N/cm), PVA (0.7 N/cm), and NaCMC (0.2 N/cm), a result that has been confirmed in numerous publications [78, 79]. Furthermore, a low degree of swelling upon exposure to electrolyte (~8%) compared to PVdF (~43%) suggests that PAA polymer will retain its mechanical properties throughout the cycling process. The cycling performance of Si-PAA electrode is also superior to other binder counterparts, exhibiting a stable cycling performance for 50 cycles with a nominal capacity ca. 750 mah/g, as shown in Figure 6-4. From an analysis of the XPS spectra of the Si- PAA electrodes, the groups suggests that this superior cycling stability can attributed to an "artificial SEI layer" enabled by the PAA. That is, the PAA partially conceals the active material surface, increasing its resiliency towards an unstable SEI formation and thus contributing to high capacity retention. Figure 6-4. Comparison in the cycling performance of SiO-based electrodes composed of a) PVdF, b) PVA, c) NaCMC, or d) PAA polymer binder as developed and reported by Komaba et al. [77]

83 70 Han et al. demonstrate the positive effects of a polymer binder formed through the copolymerization of polyacrylic acid (PAA) and polycarbodiimide (PCD) on the cycling performance of a Si-based electrode [81]. The group explains that when PAA is solely applied as a polymer binder, the free carboxylic acid groups available for binding with the Si surface may form intramolecular bonds within the polymer. That is, the functionalities offered by the binder may bond within themselves, and thus decreasing that availability of bonding to the Si surface. By copolymerization with the PCD, there is an obvious improvement to both the adhesion properties and cycling performance of the Si-C electrode. The PAA-PCD polymer binder exhibits an adhesion strength of 0.32 N/cm, while PAA alone only exhibits 0.04 N/cm. Additionally, the PAA-PCD Si electrode achieves a specific discharge capacity ca. 800 mah/g compared to the 600 mah/g achieved by PAA Si electrode, both of which have a fair capacity retention over the course of 30 cycles, as shown in Figure 6-5. Figure 6-5. Comparison of the cycling performance of Si/graphite electrodes composed of PVdF, PAA, or crosslinked PAA:PCD polymer binders with different ratios of PAA and PCD, as developed and reported by Han et al. [81]

84 71 This concept of improving the base PAA polymer through condensation reaction with other polymers is further demonstrated by Koo et al., who have developed a cross-linked PAA- CMC polymer binder through the condensation reaction of the carboxylic groups of PAA and the hydroxyl moieties of CMC [31]. Again, when compared to PVdF, CMC, and PAA, the crosslinked PAA-CMC binder was shown to offer higher reversible capacity and more stable capacity retention over the course of 100 cycles, as shown in Figure 6-6. Moreover, this type of binder system was shown to be suitable in both high temperature (60 o C) and high current rate (10C) cycling scenarios. Based on this exceptional performance, the PAA-CMC is one of the binders chosen for investigation in this thesis. Figure 6-6. Comparison of cycling performance of Si-based electrode with PVdF, PAA, CMC, or crosslinked PAA-CMC polymer binders as developed and reported by Koo et al. [31] 6.3 Results and Discussion Based on the polymer binders successfully applied to other Si electrode system as reviewed in the literature, this section of the thesis presents the electrochemical performance of the SiO-C prepared with PVdF, NaCMC, PAA, and PAA:NaCMC (i.e., crosslink between PAA

85 72 and Na:CMC). The PVdF is chosen a control sample, being the conventional polymer binder employed in Li-ion battery systems. From the literature review, it is clear that NaCMC and PAA has proven to be capable polymer binders for similar electrode system due to superior chemical functionalities that offer strong hydrogen bonding to the surface silanol groups of Si active material. Additionally, a PAA : NaCMC crosslinked binder is prepared (as described in Chapter 3) to study its effect of electrochemical performance. The cycling performance of the SiO-C electrodes prepared with these various polymer binders and cycled at is presented in Figure 6-1. Since higher active mass loading electrodes offer a more substantial and relevant challenge (i.e., in terms of mechanical integrity) for polymer binder selection, the SiO-C electrode are prepared with a relatively high active mass loading of mg/cm2. The electrodes are cycled in a voltage window of V at a current density of 100 ma/g. It is immediately clear that the PVdF control sample is unsuitable for tolerating the lithiation and delithiation of the SiO-C active material. Though the 1st cycle gravimetric capacity of the PVdF-based electrode is 881 mah/g, this is followed a by a low initial coulombic efficiency of 52% and a rapid capacity fading. By the 4th cycle, the electrode has lost all functionality. Although slightly improved over the PVdF electrode, the NaCMC-based electrode also suffers from a severe capacity fading throughout the course of 40 cycles. The NaCMC electrode achieves a 1st cycle gravimetric capacity and coulombic efficiency of 1151 mah/g and 74%, however can only retain 17% of its capacity between the 2nd and 40th cycle. The electrode prepared with the PAA and PAA : NaCMC binder provide a suitably environments for the stable cycling of SiO-C active material; in fact, the electrodes prepared with these two binders perform almost identically. As shown in Figure 6-1 and 6-2, the PAA and PAA : NaCMC achieve a 1st cycle capacity of 1199 and 1259 mah/g, respectively, and have an identical initial coulombic efficiency of 74%. Additionally, the capacity fading over the course 40 cycles is nearly identical: PAA and PAA : NaCMC have a capacity retention of 60 and 62%, respectively.

86 73 Reflecting on the literature review, these results are not entirely unexpected. For instance, the NaCMC polymer binder is known to have a lower adhesion strength and higher swelling when exposed to electrolyte compared to polyacrylate binder types [77, 82]. When this characteristic is paired with these deliberately high active mass loading electrodes, the very poor cycling performance can most likely be attributed to classic Si electrode issues, such as particle pulverization, electrode-level cracking, and detachment from the current collector. The PAA polymer binder, on the other hand, has been shown to have a high resiliency towards the electrolyte, in terms of both mechanical integrity and stable SEI formation [80]. Hence, this explains the significant improvement in both the magnitude of gravimetric capacity and capacity retention; the PAA is more capable of tolerating the Si volume expansion and maintaining the electronic and conductive pathways of the electrode. Interestingly, cross-linking the NaCMC and PAA polymers neither negatively nor positively benefits the cycling performance. Therefore, the mechanical and chemical resiliency of the PAA most likely improvises for these lacking qualities in the NaCMC polymer; meanwhile, both of these polymers can offer the carboxyl functional groups that form strong hydrogen and covalent bonds at the Si particle surface. Nevertheless, none of the selected polymer binder showcased in this chapter appear capable of truly enabling a stable cycling performance; that is, the capacity retention for all electrodes is far less than ideal. The implications of this conclusion are discussed in the following section.

87 74 Figure 6-7.Comparison of the cycling performance of SiO-C electrodes composed with PVdF, NaCMC, PAA, or PAA:NaCMC polymer binders. All electrodes are cycled within a voltage range of V and a current density of 100 ma/g. Figure 6-8. Voltage diagrams at the 1 st, 2 nd, and 40 th electrochemical cycles of the SiO-C electrodes composed with PAA and PAA:NaCMC polymer binders.

88 Conclusions Based on the cycling performance of SiO-C electrodes presented in Section 6.3, it is clear that PAA and PAA : NaCMC offer the best lithium utilization and cycling stability; in fact, the electrodes composed with either binder have nearly identical performance. When evaluated within a voltage window of V and current density of 100 ma/g, the PAA-based electrodes achieved a relatively moderate 1st cycle gravimetric capacity and coulombic efficiency of ca. 1,200 mah/g and 74%, respectively. Furthermore, the capacity retention of the electrodes from the 2nd to 40th cycle are both ca. 61%. Although capacity fading is still a significant issue in these electrodes, it is important to keep in mind that these are high active mass loading electrodes (ca mg/cm2). Comparatively, the pure NaCMC and PVdF polymer binders appear to offer insufficient attributes in the way of mechanical integrity and for maintaining conductive networks, as the electrodes composed of these binders exhibit dramatic capacity fading within the first several cycles, degrading the lithium capacity to impractically small magnitudes. More specifically, the PVdF-based electrode has lost all lithium capacity by the 4th cycle, and the NaCMC suffers from a very low 17% capacity retention from the 2nd to 40th cycle. These results can definitely be justified from the experimental studies presented in the literature review in Section 6.2. Both NaCMC and PAA are pronounced as potential candidates as polymer binders for Si electrodes, as both offer the presence of carboxyl functional groups along their primary polymer chains, and thus can allow for strong hydrogen or covalent bonding to the Si active material, whereas PVdF can offer no functionalities. The attributes of the PAA binder exceed that of NaCMC, however, in terms of adhesion strength and swelling behavior when exposed to carbonate-based electrolyte employed in Si electrodes [79, 78]. For instance, studies have shown that the PAA has an adhesion strength of 2.3 N/cm, whereas NaCMC has a significantly lower value of 0.2 N/cm [77]. Additionally, the PAA swelling when exposed to

89 76 electrolyte (ca. 8%) is much less than NaCMC (ca. 13%), and thus its mechanical properties are likely to remain unchanged when exposed to the electrolyte [77]. With these factors considered, PAA is more capable of alleviating the stresses being generated by the Si active material during lithiation and assist in maintaining their distribution throughout the electrode. Concurrently, maintaining a uniform Si distribution can help ensure that electronic and ionic conductive pathways remain available throughout the electrode, and thus lithium can continue to alloy with the majority of Si throughout the electrode. However, the capacity fading present in all of the electrodes presented in this thesis is an indication that with every cycle, active material is either being isolated from these conductive pathways or SEI formation has become unstable and thus electrolyte and lithium are forming irreversible products that continually lead to capacity fading. Fortunately, further improvement to the cycling performance of SiO-C based electrodes is not necessarily limited to uneconomical and sophisticated polymer synthesis routes. For instance, the degree of substitution of the PAA binder (i.e., the number of functional groups present per unit length of its polymer chain) can easily be modified, and thus this increase in functional groups can likely be optimized to improve the cycling performance of the electrodes. Additionally, other studies have found that the ph level of the processing solvent of the polymer binder can have a profound effect on cycling performance [73], which is another relatively simple modification that can be optimized in future work. In this study, the polymer binder remains relatively unchanged from its commercial form, aside from solution-phase mixing with a water solvent, and therefore plenty of approaches exist for further optimization of this electrode before costly alternative solutions need to be considered.

90 77 Chapter 7 Conclusions & Future Work In this chapter, the conclusions derived from Chapters 4, 5, and 6 will be reviewed and discussed. In this thesis, a multi-faceted approach was taken to reach these conclusions, where multiple parameters of the electrode fabrication and design were investigated for their influence on the electrode's cycling performance. In this way, it is demonstrated that no single component of the electrode is solely responsible for achieving good cycling performance, but instead all components must be optimized and complimentary to enable a suitable electrode integrity and utilization. In Chapter 4, the influence of active mass loading and the calendering process on the cycling performance of Si-based electrodes was evaluated. It was found that the unpressed SiO- C-based electrode offered the best electrochemical performance, even at active mass loadings up to 4.6 mg/cm 2. In fact, through the calendering process, this active mass loading can be extended up to 8.5 mg/cm 2 without a significant alteration to the cycling performance. The pressed SiO-C electrode with active mass loading of 8.5 mg/cm 2 was able to retain an areal capacity of >3.5 mah/cm 2 for up to 50 cycles, which is among the highest areal capacity reported in the literature. In Chapter 5, the influence of the carbon-coating on the cycling performance of the SiO and SiO- C electrodes is evaluated. It was found that the carbon-coating enabled a higher nominal gravimetric capacity of 2558 mah/g from the 11th to 40th cycle, compared to the nominal gravimetric capacity of 890 mah/g of the non-coated SiO. Furthermore, the SiO-C was able to retain this superior cycling performance even at higher temperature conditions, without a significant loss in capacity retention. In Chapter 6, the influence of the polymer binder on the cycling performance of the SiO-C electrodes composed of PVdF, NaCMC, PAA, and PAA:NaCMC polymer binders was evaluated. It was found that for high active mass loading electrodes, the PVdF and NaCMC polymer binders were insufficient for enabling a stable cycling

91 78 performance; however the PAA and PAA:NaCMC offered a significantly improved gravimetric capacity and capacity retention, albeit nearly identical between the two binders. Based on these results, it is clear that the carbon-coated silicon monoxide (SiO-C) electrode paired with a polyacrylic acid (PAA) polymer binder is the most suitable electrode design presented in this thesis, either for moderate cycling performance at high active mass loadings (i.e., >4 mg/cm2) or a very stable cycling performance at lower mass loadings (~1 mg/cm 2 ). Future work on this SiO- C electrode should be aimed at improving the capacity retention and cycle life of the electrode, which will most likely entail active material or polymer binder modifications that allow for better mechanical and chemical stability during the process of electrochemical cycling. Meanwhile, the other goal of incorporating an economical and industrial-friendly electrode fabrication process should be maintained. In this effort, several approaches exist to modify the existing system without significantly increasing the processing cost. For instance, in substitute of the manual hand-milling of solid active materials, a high energy ball milling process can be used for decomposing bulk materials into significantly smaller and more uniform sub-micron particle distributions; the smaller the active material that can be achieved without resorting to costly nanoparticle synthesis, the better. Furthermore, several economical modifications can be made to the polymer binder, such as increasing its degree of substitution and molecular weight or adding other functionalities through a solution phase mixing process that would allow for a more conformal and stronger bonding at the particle-binder interface.

92 79 Chapter 8 Early Development of a Battery Cell for In Situ Characterization 8.1 Introduction As discussed earlier, the lithium-alloys are typically unstable electrodes primarily because of their immense volume expansion upon lithiation. From the previous investigations discussed in this thesis, it has been shown that an optimal combination of active and non-active material, paired with appropriate fabrication conditions, will result in an improved cycling performance for the silicon-based electrode. The capacity and voltage diagrams, however, only provide an indirect knowledge of the actual behavior of the silicon-based electrodes during electrochemical reactions, and hence a fundamental understanding of the real-time electrode behavior is left to be desired. A variety of ex situ analytical techniques exist that approach an understanding of real-time electrode behavior. The mechanical strength and structural integrity of an electrode can be surmised by a standardized peel tests [77]; however this lacks insight on mechanical property changes when wetted with electrolyte or during electrochemical reaction. From the electrolyte uptake of a pristine (i.e., non-lithiated) electrode, assuming the mechanical properties of the electrode will not significantly change during lithiation/delithiation seems unreasonable considering the dynamic behavior of the electrode morphology. Other spectroscopy techniques exist, such as Fourier transform infrared spectroscopy (FTIR) and Ramanspectroscopy, that have been invaluable for directly evidencing chemical bonding at the siliconbinder or silicon-electrolyte interface [74]; but again, these must be performed prior-to or following electrochemical cycling in order to transfer the specimen from battery casing to an analytical specimen holder.

93 80 For the silicon-based electrode, or other lithium-alloy materials, there is wealth of knowledge to be gained from a system that could enable in situ, real-time characterization of the electrode. For instance, certain phenomena that are only observable during electrochemical cycling, such as silicon particle growth or lithium dendrite formation, could be directly examined in real-time and correlated to electrochemical performance. This functionality could be extended to observe the dynamics of SEI growth or electrode fracturing, all of which may lead to direct evidence of failure mechanisms and, in turn, the more intelligent design of electrode systems. In situ observation and measurement of lithium-ion battery systems has recently received much attention due to the dramatic structural changes of alloy-type electrodes. Timmons and Dahn were among the first to employ in situ measurement on an alloy-type electrode [83]. In order to capture the dynamic behavior of an amorphous silicon-tin active material (a- Si0.64Sn0.36), the team developed an open Li-metal cell to contain the electrode, which could be observed by a modified microscope that could be mounted inside an argon-filled glove box. The research represents an impressive step towards in situ measurements, however the artificial electrode assembly being observed was not representative of an actual battery system. That is, the open cell design ignored the impact of a high compression electrode assembly, where particle mobility and expansion would be impeded by rigid boundaries. Furthermore, the microscope employed lacked the necessary resolution to observe particle surface interactions and could only observe the front face (i.e., looking down on the electrode disc) of the electrode, and thus it was virtually impossible to study thickness variations during cycling. Harris et al. aim to study the degradation of electrodes through an in situ measurement of time-dependent Li spatial distributions and transport throughout a graphite electrode [84]. The in situ measurement is enabled by an optical half-cell similar to that previous described [83]. Since graphite changes color based on the extent of lithiation, the optical half-cell and microscope were sufficient for the group to gain direct measurements of Li-ion transport in a porous electrode,

94 81 which in turn could be linked to degradation mechanisms in the electrode. Although the optical system employed in this study was only relevant for macroscopic observation on the electrode surface, the concept of linking direct in situ measurement to fundamental knowledge of degradation mechanisms was validated. Yue and Harris later showed that the same optical halfcell could be used for in situ measurements of strain mapping throughout the graphite electrode during the lithiation process, and thus validate previously established degradation models [85]. Again, through the use of a low-resolution optical microscope, the group used a digital image correlation technology to observe and measure particle displacements, and in turn the strain throughout the electrode. From the strain mapping, the group was able to conclude that porosity had significantly changed during the course of lithiation, which was contrary to a constant porosity employed in many models. Beaulieu et al. explored the use of in situ AFM paired with optical microscopy to study a SiSn electrode fracturing during the lithiation process [12]. The team was able to successfully observe in real-time the fracturing of SiSn-based electrode; but perhaps more interesting was the observation that the fracturing could be reversed (i.e., restored) by delithiating the material. Although the system under investigation was not representative of a commercial battery system, where electrodes and separator are in intimate contact, the conclusion offered a fundamental understanding on alloy-type electrode behavior and challenged previously reported elasticity theories on the subject. 8.2 Design and Fabrication of the In Situ Battery Cell Based on the concept described above, an electrochemical cell that could function as a conventional battery cell and provide realistic internal battery cell pressures was developed for use with a high-resolution (i.e., 100 nm resolution) optical microscope. Moreover, this device was

95 82 designed with affordable materials as to reduce the cost of implementing this type of battery cell in a normal laboratory environment. With this design philosophy in mind, here the 3D modeling design, material selection, and fabrication of a unique and simple in situ battery cell, hereafter referred to as the in situ cell, is discussed. From the design of the in situ cell, shown as both a concept rendering in Figure 8-1 and actual final product in Figure 8-2, several of the fundamental components of a conventional battery cell can be readily recognized. For the encasement of the internal battery components, a relative inert and affordable plastic material, polypropylene (McMaster # 8782K39), is machined to have a rectangular profile; these flat faces of the in situ cell will allow it to sit level on the stage of the optical microscope. In Figure 8-1a or 8-2, a top view of the in situ cell is shown, a perspective that offers a view through its transparent glass window. Figure 8-1. Rendering of the initial conceptual design of the in situ battery cell.

96 83 The glass window is a 1/16 thick, ground and polished, quartz glass disc (Technical Glass Products) that was selected for its compatibility with the optical microscope s objective lens working distance, as well as being chemically inert towards the ethylene carbonate-based electrolyte employed in the cell. It should be noted that a glass window of greater thickness would present serious compatibility issues with the working distance of the microscope s objective lens; this is considering that the objective lens must be able to focus on a specimen through the thickness of the glass, as well as through the depth of the liquid electrolyte between the glass and the sample. Figure 8-2. An image of the actual fabricated in situ cell without the electrode assembly or transparent glass windows installed. The quartz glass window is mounted into the in situ cell in a very careful and specific manner to avoid leakage of external air from the atmosphere to the inside of the cell. First, an O- ring is seated into a precisely machined groove directly below the glass window, which is most easily seen in Figure 8-2. Given that the chemical compatibility of the battery electrolyte has not

97 84 been thoroughly studied for materials other than typical battery components (e.g., polymer separator, metal casing), the O-ring material was chosen out of four candidates that were distinguished for having a high degree of chemical inertness, as will be described in Section 8.3. A chemically inert Aflas O-ring (McMaster AS568A) was selected for this application and was able to remain stable for several days without noticeable degradation. The O-Ring is the primary safeguard against the leakage of air into the cell, or electrolyte leakage out of the cell, ensuring a well-sealed gap between the plastic cell shelf and the mounted glass disc. As a secondary precaution for air leakage, a thin and uniform layer of silicone sealant (3M Super Silicon Sealant, McMaster #74955A53), a material that is expected to also be chemically inert towards the electrolyte, was applied in the gap between the plastic shelf and glass disc. Finally, a custommachined stainless steel bracket is mounted above the glass surface and compresses the O-ring through six low-strength steel bolts (McMaster #91465A107). The complete assembly of the top viewport is shown in Figure 8-3. Figure 8-3. An image of the actual fabricated in situ cell with the transparent glass window and top bracket installed, but without the electrode assembly installed.

98 85 Now attention is turned to the custom-machined electrode assemblies (dimensions in Appendix B) of the in-situ cell. As shown in Figure 8-1c, the copper current collector (CCC, anode-side) assembly is shown to consist of the copper plug with a large and small diameter section. An Aflas O-ring is mounted around the smaller diameter section of the plug, which has the function of preventing air leakage into the cell. As shown in Figure 8-3, the CCC assembly is compressed into a side-face through-hole of the plastic cell; this compression in enabled through a stainless steel bracket a similar concept as used for the glass window of the top viewport that compresses the O-ring through the tightening of bolts. The stainless steel current collector (SSCC, cathode-side) assembly is similar to the CCC assembly; however, instead of a smaller diameter section that penetrates into the in situ cell, a stainless steel spring (McMaster 8969T5) is mounted onto the SSCC and penetrated into the in situ cell. The incorporation of this spring in the SSCC assembly will enable fundamental mechanical property measurements of the electrode system, as described later in Section 8.4. Lastly, a glass disc window identical to the top-side viewport will be mounted on the bottom-side of the in situ cell. Unlike the top-side viewport, where the stainless steel bracket is mounted directly on the surface of the in situ cell, the bottom viewport is assembled in a depressed region. This depressed region was incorporated to enable the in situ cell to properly mount the microscope stage. For a more detailed summary of the bill of materials and design schematics, please refer to Appendix A and B, respectively. As mentioned above, a stainless steel spring has been incorporated in the SSCC assembly to enable fundamental measurements, based on Hooke s Law, of the mechanical properties of an electrode. For instance, with the material properties and initial length of the spring, the following fundamental relationships can be invoked:

99 86 F = kx P = F A where F is the force applied by the compressed spring, P is the pressure induced by the force of the spring on the area of the electrode disc, k is the spring constant, and A is the area of the electrode disc on which the force acts. By complimenting observations from the optical microscope with the theoretical pressures determined with the above equations, the influence of pressure on the electrode system could be deduced. Other simple measurements may also be achievable with this system, such as quantifying the thickness changes or delamination of the silicon electrode during the electrochemical reaction. This objective aligns with the motivations of the research defined in the brief literature review on this topic. In addition to providing a mechanism to estimate mechanical properties, the spring acts to mimic the compact electrode assembly found in commercial battery cells. Figure 8-4. Image of the completely assembled and electrolyte-filled in situ cell in an argon-filled glove box.

100 Materials and Compatibility Although the majority of the in situ cell components are of materials (e.g., copper, stainless steel, polypropylene) proven to be compatible and inert towards the ethylene carbonatebased electrolyte, the O-ring material is unique to this design, and thus the chemical inertness of several O-ring candidates were evaluated prior to battery operation. In this effort, four candidate O-ring materials were selected, namely, polytetrafluoroethylene (PTFE), Kalrez, Aflas, and fluorosilicone. These O-rings represent a class of materials known as fluoroelastomers that are very non-reactive with other materials due to the strong carbon-fluorine bonds along their primary polymer chains. Nonetheless, each O-ring material was evaluated for its chemical reactivity towards the ethylene carbonate-based electrolyte used in all the lithium-ion cells discussed in this thesis. To test the chemical reactivity between the O-rings and electrolyte, small pieces of each O-ring were submerged in fresh electrolyte and observed over the course of two weeks. All samples were prepared and evaluated in an argon-filled glove box. The samples were evaluated on the extent of their physical decomposition or the change of solution color. Initially, all O-ring samples are submerged in a clear liquid electrolyte solution. For the Aflas and PTFE samples, the solution can maintain this clear transparency for approximately 2 days before a very slight discoloration, similar to a beige color, is observable in the solution. Within a week, the electrolyte solution for the Aflas and PTFE has obviously transitioned from a clear liquid to a beige color liquid; although the Aflas and PTFE O-ring material still largely appeared intact, even after 3 weeks. Among the four O-rings evaluated in this manner, the Aflas O-ring material was capable of maintaining a clear electrolyte solution for the longest duration of time, approximately 2 days, and therefore was chosen as the O-ring material for the in situ cell.

101 88 The in situ cell was also evaluated on its design to prevent air leakage into the cell. In this evaluation, a piece of lithium-foil was deposited inside the assembled in situ cell (no electrolyte added, silicone sealant applied) within an argon-filled glove box. Since lithium is highly reactive with oxygen and will turn a dark color upon exposure, the material was used as an indicator for air leakage into the cell. As shown in Figure 8-5, once the in situ cell is removed from the glove box environment, the lithium foil remains pristine for approximately 4 days. After this period of time, the lithium foil shows clear indications of oxidation, noted by the darkened areas in Figure 8-5. Therefore, it is concluded that the in situ cell could maintain a suitable environment for electrochemical reaction for approximately five days, after which the integrity of the bracket compression or silicone sealant is uncertain. Figure 8-5. Leak test of the in-situ cell. Lithium-foil is encased in the in-situ cell, previously assembled in an argon-filled glove box, and left in an air-filled environment. The in-situ cell remains suitable air-free for about 6 days before signs of lithium oxidation occur.

102 Early Stages of Characterization Electrochemical Cycling As a proof-of-concept for the in situ cell to operate and behave as a normal battery cell, here the electrochemical cycling of the SiO-C electrodes in a lithium half-cell is conducted using the fully-assembled in situ cell. As described in Chapter 3, the SiO-C electrode, separator, lithium foil, and spacer platforms are tightly compressed in the center of the in situ cell, where they can be observed through the top viewport glass window. Once the in situ cell is removed from the argon-filled glove box, it is immediately and gently wiped down with ethanol to remove any chemical residue from the glove box. The in situ cell is prepared for the electrochemical test by attaching alligator clips to its copper and stainless steel current collectors. In this manner, a voltage may be applied to the cell and a current can pass through the external load. Once this electrical connection was established, a charge-discharge program is enacted through which the battery is cycled a 100 ma/g current density within a voltage window V over the course of 16 cycles. During the course of the cycling procedure, it was noted that both the internal environment of the in situ cell remained stable; that is, no volatile or exothermic reactions were detected. From Figure 8-6, the electrical operation of the in situ cell is shown to be successful as it has enabled the stable cycling of the SiO-C electrode over the course of 16 cycles. This is a significant achievement considering the novel and untested design of the in situ cell prior to this investigation.

103 90 Figure 8-6. Electrochemical evaluation of the influence of O-ring materials on cycling performance. For this evaluation, a SiO/C active material electrode is cycled at 0.1 C-Rate In situ microscopy The objective of the optical microscopy characterization of the in situ cell is to evaluate the cell design in enabling a high-resolution observation of electrode materials, such as active particles. This objective is complicated by requiring the working distance of the objective lens to observe a sample submerged beneath a thin layer of glass and liquid electrolyte, while maintaining a resolution of at least 1 micron. By designing the in situ cell based around the dimensions of a particular optical microscope, a 1micron resolution has been ensured as a minimum resolution for the material observed through the top viewport of the in situ cell. The optical microscope employed in this investigation is an Olympus BX51 Optical Microscope with a 100x Magnification Objective Lens, well-suited for the observation of specimens on the scale of 1 micron. The microscope is outfitted with an Olympus Nomarski Differential Interference Contrast Prism paired with a DP50 digital camera that enables high resolution images through a variety of surrounding medium. The in situ cell is mounted level on

104 91 the observation stage of the microscope and a fiber-optic illuminator light source is incident on the top viewport of the in situ cell to illuminate the internal cell chamber. In the ideal scenario, observations through the optical microscope would show a sideview of the electrode perimeter, in which the thickness change of the electrode and, possibly, the constituent active and non-active materials could be observed and measured. In this early-stage characterization, however, the observations show that fragments of the electrode have become detached from the assembly and are free-floating in the electrolyte, as shown in Figure 8-7. Furthermore, the fiber-optic light source proved to be insufficient for illuminating the tightlypacked electrode assembly. Nevertheless, the images from the optical microscope are encouraging in that the minimum goal of 1 micron resolution has been surpassed. Indeed, the free-floating electrode fragments observed are on the length scale of tens of nanometers. If the in situ cell can be modified to enable more light to penetrate the electrode chamber, this type of resolution should provide tremendous insight on the real-time thickness variation and morphology changes occurring during electrochemical cycling. From these early stage observations with the optical microscope, several design modifications for the in situ cell can be suggested to enable more suitable imaging of the intact electrode assemble. Figure 8-7. Optical microscope images of the material within the in situ cell.

105 Future Outlook From the optical microscope images acquired in this investigation, several conclusions and future directions can be made for improving the in situ cell. First, it is apparent that the internal electrode chamber of the in situ cell is unable to receive adequate illumination, and thus the electrode assemble appears as a dense dark area under the microscope. Therefore, the in situ cell needs to be modified in a way that light is not only incident on the top viewport of the in situ cell, but from different directions as well. For instance, machining through-holes on the sides of the in situ cell that are not occupied by the electrodes could enable a complete illumination of the perimeter of the electrode assemble, and thus allow for accurate thickness change observations. Also, to reduce the amount of electrode fragments in the solution, a new approach should be developed that better allows for the insertion of electrodes during cell assembly. At present, the electrode assemble is hindered by the effort taken to keep the spring compressed while tightening the outside electrode brackets. During the process, the spring undoubtedly induces a variety of shear and non-normal stresses upon the face of the electrode, and thus induces some fracturing. A possible solution to this problem would be to design a new clamping mechanism that relieves the user from stabilizing the spring while the brackets are compressed. In summary, the design and early-stage application of this affordable benchtop in situ battery cell represents an accessible approach towards in situ real-time battery cell observations.

106 93 Appendix A In Situ Cell Bill of Materials Table A-1. Bill of Materials for the In Situ Cell

107 94 Appendix B Schematic Diagram of the In Situ Cell #44 Drill Bit Figure B-1. Schematic of the Bottom-Face of In Situ Cell #44 Drill Bit #29 Drill Bit Figure B-2. Schematic of the Top-Face of In Situ Cell #29 Drill Bit Figure B-3. Schematic of the Side-Face of the In Situ Cell UNLESS OTHERWISE SPECIFIED: NAME DATE