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Engineering high performance electrodes for energy storage devices from low-cost, sustainable and naturally… Watson, Timothy 2016

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Engineering High Performance Electrodes for EnergyStorage Devices from Low-cost, Sustainable and NaturallyAbundant BiomaterialsbyTimothy WatsonB.Sc, The University of British Columbia, 2014A THESIS SUBMITTED IN PARTIAL FULFILLMENTOF THE REQUIREMENTS FOR THE DEGREE OFMaster of Applied ScienceinTHE FACULTY OF GRADUATE AND POSTDOCTORALSTUDIES(Biomedical Engineering)The University of British Columbia(Vancouver)December 2016c© Timothy Watson, 2016AbstractWith the increased global push towards sustainable energy utilization, the needfor advanced energy storage technologies has become increasingly important ascountries seek to integrate rapidly advancing renewable energy technologies likewind and solar. At the same time, the burgeoning electric vehicle and wearableelectronics industries are fuelling demand for lower-cost energy storage deviceswith high energy capacities and longer cycle lives. Currently, despite huge leapsin performance of Li-Ion batteries in recent years, the technology is approachingits predicted limits and new solutions will be needed to keep up with the demandof current and future electrical devices. At a time where scientific applications ofnanomaterials and nanofabrication is on the rise, there exists an opportunity to takeadvantage of our increased understanding of nanotechnology to significantly im-prove existing energy storage devices and to unlock the potential of next-generationenergy storage technologies.In this work, binder-free and porous graphitic nanofibre electrodes producedfrom low-cost and sustainable softwood kraft lignin are devised and proposed asa platform for the development of high-performance energy storage devices. Mo-tivated by difficulties facing some key energy storage technologies, scalable elec-trospinning of lignin and polyethylene oxide (PEO) precursor materials, followedby a hydrothermal treatment and carbonization in an inert atmosphere yields free-standing interconnected nanofibre electrodes with tunable porosity, high conductiv-ity and superior electrochemical performance. Electrical impedance spectroscopymeasurements of the optimized porous nanofibre electrodes demonstrate a con-ductivity reaching 18.39 S cm−1, while Brunauer-Emmett-Teller specific surfacearea measurements yield a specific surface area as high as 1258.41 m2 g−1. Su-percapacitor devices revealed highly symmetric cyclic voltammetry results whichdemonstrated a gravimetric capacitance approaching 112 F g−1 at a voltage scanrate of 5 mV s−1. Galvanostatic charge/discharge experiments show reversible su-percapacitor behaviour, a high capacity even at elevated voltage scan rates up to200 mV s−1 and exhibit excellent cyclic stability, retaining 91% of their initial ca-iipacity after 6000 cycles. This work demonstrates the use of sustainable and abun-dant softwood Kraft lignin as source for porosity-tunable electrodes with high ca-pacitance and stability as a demonstration of nature sourced and high performanceelectronic devices.iiiPrefaceThe author was responsible for all major areas of concept formation and exper-iment design with the assistance and guidance of Dr. Saeid Soltanian from theFlexible Electronics and Energy Lab at UBC.Results in Chapter 4 and 5 are based on work conducted by the author in UBC’sAdvanced Fibrous Materials Lab under the supervision of Dr. Frank Ko and inUBC’s flexible Energy and Electronics Lab under the supervision of Dr. PeymanServati. X-Ray Diffraction measurements were performed by Anita Lam fromUBC’s department of Chemistry. Raman Spectroscopy measurements were per-formed by Jiaxin Ke from the Department of Materials Engineering at UBC. All ofthe data analysis along with all other data collection was performed by the author.Portions of Chapter 4 and Chapter 5 are contained in a submission for publicationwritten by the author and co-authored by Dr. Frank Ko, Dr. Peyman Servati andDr. Saeid Soltanian. The submission is currently under review.ivTable of ContentsAbstract . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . iiPreface . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . ivTable of Contents . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . vList of Tables . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . viiiList of Figures . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xAcknowledgments . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . xiv1 Introduction . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 12 Objectives . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 63 Literature Review and Background . . . . . . . . . . . . . . . . . . 93.1 Batteries . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 93.2 Supercapacitors . . . . . . . . . . . . . . . . . . . . . . . . . . . 223.3 Carbon Nanofibre Composites . . . . . . . . . . . . . . . . . . . 283.4 Electrospinning . . . . . . . . . . . . . . . . . . . . . . . . . . . 323.5 Lignin . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 404 Lignin-based Porous Carbon Nanofibre Electrodes . . . . . . . . . . 454.1 Materials & Methods . . . . . . . . . . . . . . . . . . . . . . . . 464.1.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . 464.1.2 Electrospinning . . . . . . . . . . . . . . . . . . . . . . . 464.1.3 Removal of Sacrificial Polymer . . . . . . . . . . . . . . 474.1.4 Stabilization and Carbonization . . . . . . . . . . . . . . 484.2 Characterization and evaluation . . . . . . . . . . . . . . . . . . . 504.3 Results and Analysis . . . . . . . . . . . . . . . . . . . . . . . . 51v4.3.1 Carbonization Parameter Optimization . . . . . . . . . . . 514.3.2 Microstructural Observations . . . . . . . . . . . . . . . . 534.3.3 Surface Area and Porosity . . . . . . . . . . . . . . . . . 574.3.4 X-Ray Diffraction Analysis . . . . . . . . . . . . . . . . 614.3.5 Lignin Fibre Conductivity . . . . . . . . . . . . . . . . . 644.3.6 Raman Spectroscopy . . . . . . . . . . . . . . . . . . . . 654.3.7 Thermogravimetric Analysis . . . . . . . . . . . . . . . . 685 Supercapacitors From Porous Lignin-based ECNF Electrodes . . . 725.1 Materials & Methods . . . . . . . . . . . . . . . . . . . . . . . . 725.1.1 Materials . . . . . . . . . . . . . . . . . . . . . . . . . . 725.1.2 Electrode Preparation . . . . . . . . . . . . . . . . . . . . 735.1.3 Test Cell Fabrication . . . . . . . . . . . . . . . . . . . . 735.2 Characterization and Evaluation . . . . . . . . . . . . . . . . . . 745.3 Results and Analysis . . . . . . . . . . . . . . . . . . . . . . . . 745.3.1 Cyclic Voltammetry . . . . . . . . . . . . . . . . . . . . 745.3.2 Galvanostatic Charge/Discharge . . . . . . . . . . . . . . 805.3.3 Cycle Life Testing . . . . . . . . . . . . . . . . . . . . . 835.3.4 Electrical Impedance Spectroscopy . . . . . . . . . . . . 856 Discussion . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 886.1 Structural Suitability of Porous Nanofibre Mats . . . . . . . . . . 886.2 Electrochemical Suitability of ECNF Mats for Supercapacitor Ap-plications . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 936.3 Commercial Considerations . . . . . . . . . . . . . . . . . . . . . 967 Conclusions and Recommendations . . . . . . . . . . . . . . . . . . 987.1 Porous Lignin-based Nanofibrous Electrodes for Energy StorageDevices . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 987.2 Recommendations for Future Work . . . . . . . . . . . . . . . . . 103Bibliography . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 106viA Additional Cyclic Voltammograms . . . . . . . . . . . . . . . . . . . 116A.1 Cyclic Voltammograms of ECNFs(99/1) at Various Scan Rates. . . 116A.2 Cyclic Voltammograms of P-ECNFs(99/1) at Various Scan Rates. 118A.3 Cyclic Voltammograms of P-ECNFs(98/2) at Various Scan Rates. 120A.4 Cyclic Voltammograms of P-ECNFs(97/3) at Various Scan Rates. 122B Additional Galvanostatic Charge/Discharge Traces . . . . . . . . . . 127B.1 Galvanostatic Charge/Discharge Traces for ECNFs(99/1) at Vari-ous Current Densities. . . . . . . . . . . . . . . . . . . . . . . . . 127B.2 Galvanostatic Charge/Discharge Traces for P-ECNFs(99/1) at Var-ious Current Densities. . . . . . . . . . . . . . . . . . . . . . . . 129B.3 Galvanostatic Charge/Discharge Traces for P-ECNFs(98/2) at Var-ious Current Densities. . . . . . . . . . . . . . . . . . . . . . . . 130C Additional SEM Images . . . . . . . . . . . . . . . . . . . . . . . . . 132C.1 SEM Images of P-ECNFs(99/1). . . . . . . . . . . . . . . . . . . 133C.2 SEM Images of P-ECNFs(98/2). . . . . . . . . . . . . . . . . . . 134C.3 SEM Images of P-ECNFs(97/3). . . . . . . . . . . . . . . . . . . 135viiList of TablesTable 3.1 Voltage and theoretical specific energy of lithium based energy-storage devices. . . . . . . . . . . . . . . . . . . . . . . . . . 18Table 4.1 Conductivities of lignin/PEO ECNFs(99/1) after stabilization at230◦C and carbonization at various temperatures. . . . . . . . 52Table 4.2 Conductivities of lignin/PEO ECNFs(99/1) after stabilization atvarious temperatures and carbonization at 900◦C. . . . . . . . 53Table 4.3 Effect of stabilization time and rate on the conductivities oflignin/PEO ECNFs(99/1). . . . . . . . . . . . . . . . . . . . . 53Table 4.4 Average diameters of different ECNFs and P-ECNFs. . . . . . 57Table 4.5 BET specific surface area and pore volume of different ECNFsand P-ECNFs acquired from N2 adsorption isotherms. . . . . . 59Table 4.6 Conductivity of prepared ECNFs and P-ECNFs. . . . . . . . . 64Table 5.1 Gravimetric capacitance in F g−1 of ECNFs calculated from CVcurves. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78Table 5.2 Volumetric capacitance in F cm−3 of ECNFs calculated fromCV curves. . . . . . . . . . . . . . . . . . . . . . . . . . . . . 78Table 5.3 Capacitance of ECNF electrodes after cycling as a percentageof the capacitance measured on the first cycle. . . . . . . . . . 83Table 6.1 Conductivity of different carbon nanofibres as reported in theliterature. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 89Table 6.2 Specific surface area of different carbon nanofibres as reportedin the literature. . . . . . . . . . . . . . . . . . . . . . . . . . 91Table 6.3 Specific surface area of different carbon materials as reported inthe literature. . . . . . . . . . . . . . . . . . . . . . . . . . . . 92Table 6.4 Pore Volume of different mesoporous carbon nanofibres as re-ported in the literature. . . . . . . . . . . . . . . . . . . . . . . 93viiiTable 6.5 Specific capacitance of various carbon materials measured inaqueous electrolytes as reported in the literature. . . . . . . . . 94Table 6.6 Volumetric capacitance of various carbon materials measured inaqueous electrolytes as reported in the literature. . . . . . . . . 95Table 7.1 Performance of porous and lignin-based nanofibrous electrodescompared to non porous PAN based electrodes. . . . . . . . . . 101ixList of FiguresFigure 2.1 Schematic showing the current and expected energy densitiesfor various battery technologies. . . . . . . . . . . . . . . . . 7Figure 3.1 Schematic of a simple battery. . . . . . . . . . . . . . . . . . 10Figure 3.2 Specific and volumetric energy density of some of the estab-lished secondary battery types. . . . . . . . . . . . . . . . . . 15Figure 3.3 Schematic of the polysulfide shuttling effect in Li-S batteries. 20Figure 3.4 Battery technology roadmap and characteristics. . . . . . . . 22Figure 3.5 Schematic of a supercapacitor during charge and during dis-charge. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 25Figure 3.6 Different strategies for encapsulating sulfur in 3D frameworks. 29Figure 3.7 Typical electrospinning setup. . . . . . . . . . . . . . . . . . 35Figure 3.8 Close-up of the structure of the electrospinning jet. . . . . . . 36Figure 3.9 3 Step process for lignin-based carbon nanofibre preparation. . 43Figure 4.1 Graphical representation of the stabilization and carbonizationprocess. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 48Figure 4.2 Fabrication process for porous lignin-based nanofibre mat elec-trodes. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 49Figure 4.3 SEM images of ECNF(99/1) at magnifications of 10k and 40k. 54Figure 4.4 SEM images of P-ECNFs(99/1) at magnifications of 10k and40k. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 54Figure 4.5 SEM images of P-ECNFs(98/2) at magnifications of 10k and40k. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 55Figure 4.6 SEM images of P-ECNFs(97/3) at magnifications of 10k and40k. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 56Figure 4.7 N2 adsorption isotherms of the prepared P-ECNFs and ECNFs. 58Figure 4.8 Pore size distribution histograms of the prepared P-ECNFs andECNFs. . . . . . . . . . . . . . . . . . . . . . . . . . . . . . 60xFigure 4.9 Specific surface area of ECNFs and P-ECNFs as a function oftheir precursor PEO content. . . . . . . . . . . . . . . . . . . 61Figure 4.10 Total pore volume of ECNFs and P-ECNFs as a function oftheir precursor PEO content. . . . . . . . . . . . . . . . . . . 62Figure 4.11 XRD patterns for different ECNF and P-ECNFs. . . . . . . . 63Figure 4.12 Conductivity of ECNFs and P-ECNFs as a function of theirprecursor PEO content. . . . . . . . . . . . . . . . . . . . . . 65Figure 4.13 Raman spectra for different ECNF and P-ECNF. . . . . . . . 66Figure 4.14 ID/IG values of ECNFs and P-ECNFs as a function of theirprecursor PEO content. . . . . . . . . . . . . . . . . . . . . . 68Figure 4.15 TGA profiles of P-ECNFs and ECNFs. . . . . . . . . . . . . 69Figure 5.1 Schematic representation of the electrochemical testing cell. . 73Figure 5.2 Cyclic voltammograms of different ECNFs at a scan rate of30 mV s−1. . . . . . . . . . . . . . . . . . . . . . . . . . . . 76Figure 5.3 Cyclic voltammograms of different ECNFs at a scan rate of50 mV s−1. . . . . . . . . . . . . . . . . . . . . . . . . . . . 77Figure 5.4 Cyclic voltammograms of P-ECNFs(97/3) at scan rates of 5,10, 30, 50, 70, 100, 150, and 200 mV s−1. . . . . . . . . . . . 79Figure 5.5 Specific capacitance of P-ECNFs calculated from CV curvesat different sweep rates. . . . . . . . . . . . . . . . . . . . . . 80Figure 5.6 Galvanostatic charge/discharge curves of ECNFs and P-ECNFsat a current density of 400 mA g−1. . . . . . . . . . . . . . . . 81Figure 5.7 Galvanostatic charge/discharge curves of P-ECNFs(97/3) at cur-rent densities of 6000, 5000, 4000, 3000, 2000, 1000, and400 mA g−1. . . . . . . . . . . . . . . . . . . . . . . . . . . 82Figure 5.8 Cyclic stability of P-ECNFs and ECNFs(99/1) after 8000 cy-cles in 6M KOH. . . . . . . . . . . . . . . . . . . . . . . . . 84Figure 5.9 Nyquist plots from EIS analyses of ECNFs(99/1) and P-ECNFs. 86Figure A.1 CV trace of ECNFs(99/1) at a scan rate of 5 mV s−1. . . . . . 116Figure A.2 CV trace of ECNFs(99/1) at a scan rate of 10 mV s−1. . . . . . 117Figure A.3 CV trace of ECNFs(99/1) at a scan rate of 30 mV s−1. . . . . . 117xiFigure A.4 CV trace of ECNFs(99/1) at a scan rate of 50 mV s−1. . . . . . 118Figure A.5 CV trace of P-ECNFs(99/1) at a scan rate of 5 mV s−1. . . . . 118Figure A.6 CV trace of P-ECNFs(99/1) at a scan rate of 10 mV s−1. . . . 119Figure A.7 CV trace of P-ECNFs(99/1) at a scan rate of 30 mV s−1. . . . 119Figure A.8 CV trace of P-ECNFs(99/1) at a scan rate of 50 mV s−1. . . . 120Figure A.9 CV trace of P-ECNFs(98/2) at a scan rate of 5 mV s−1. . . . . 120Figure A.10 CV trace of P-ECNFs(98/2) at a scan rate of 10 mV s−1. . . . 121Figure A.11 CV trace of P-ECNFs(98/2) at a scan rate of 30 mV s−1. . . . 121Figure A.12 CV trace of P-ECNFs(98/2) at a scan rate of 50 mV s−1. . . . 122Figure A.13 CV trace of P-ECNFs(97/3) at a scan rate of 5 mV s−1. . . . . 122Figure A.14 CV trace of P-ECNFs(97/3) at a scan rate of 10 mV s−1. . . . 123Figure A.15 CV trace of P-ECNFs(97/3) at a scan rate of 30 mV s−1. . . . 123Figure A.16 CV trace of P-ECNFs(97/3) at a scan rate of 50 mV s−1. . . . 124Figure A.17 CV trace of P-ECNFs(97/3) at a scan rate of 70 mV s−1. . . . 124Figure A.18 CV trace of P-ECNFs(97/3) at a scan rate of 100 mV s−1. . . . 125Figure A.19 CV trace of P-ECNFs(97/3) at a scan rate of 150 mV s−1. . . . 125Figure A.20 CV trace of P-ECNFs(97/3) at a scan rate of 200 mV s−1. . . . 126Figure B.1 Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 400 mA g−1. . . . . . . . . . . . . . . . . . . 127Figure B.2 Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 1000 mA g−1. . . . . . . . . . . . . . . . . . . 128Figure B.3 Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 2000 mA g−1. . . . . . . . . . . . . . . . . . . 128Figure B.4 Galvanostatic charge/discharge trace for P-ECNFs(99/1) at acurrent density of 400 mA g−1. . . . . . . . . . . . . . . . . . 129Figure B.5 Galvanostatic charge/discharge trace for P-ECNFs(99/1) at acurrent density of 1000 mA g−1. . . . . . . . . . . . . . . . . 129Figure B.6 Galvanostatic charge/discharge trace for P-ECNFs(99/1) at acurrent density of 2000 mA g−1. . . . . . . . . . . . . . . . . 130Figure B.7 Galvanostatic charge/discharge trace for P-ECNFs(98/2) at acurrent density of 400 mA g−1. . . . . . . . . . . . . . . . . . 130xiiFigure B.8 Galvanostatic charge/discharge trace for P-ECNFs(98/2) at acurrent density of 1000 mA g−1. . . . . . . . . . . . . . . . . 131Figure B.9 Galvanostatic charge/discharge trace for P-ECNFs(98/2) at acurrent density of 2000 mA g−1. . . . . . . . . . . . . . . . . 131Figure C.1 SEM images of P-ECNFs(99/1) at various magnifications. . . 133Figure C.2 SEM images of P-ECNFs(98/2) at various magnifications. . . 134Figure C.3 SEM images of P-ECNFs(97/3) at various magnifications. . . 135xiiiAcknowledgmentsI want to express my gratitude to Dr. Frank Ko and to Dr. Peyman Servati fortheir continued support and mentorship throughout the completion of this work. Ialso would like to thank Dr. Saeid Soltanian for all of the advice and guidance thathe provided me during this project, for his patience and willingness to answer myquestions and for the work he did to ensure that all of the necessary materials andequipment needed for this work were available. This project would not have beenpossible without his support and the support of all of the students and staff in boththe Advanced Fibrous Materials Lab and the Flexible Electronics Engineering Lab.I would also like to thank Lynn Wan for her expertise surrounding lignin and elec-trospinning, Anita Lam for her help with XRD testing, Jiaxin Ke for his help per-forming Raman spectroscopy, Sally Finora for her expertise regarding BET testing,Mijung Cho for assisting me with TGA and Yuta Dobashi for his advice regardingsupercapacitor testing.Portions of this work were made possible through funding from a Strategic Re-search Grant from the Natural Sciences and Engineering Research Council ofCanada along with funding from the Canada Foundation for Innovation.xivChapter 1IntroductionOver the past decades, the importance of efficient electrical energy storage hasgrown exponentially following an explosion of devices that rely on electrical powerand a significant push towards sustainable energy practices. As of 2012, the globalproduction of electricity reached about 22,200 TWh annually, 70% of which wasgenerated through the burning of fossil fuels including coal, oil and natural gas[1]. Today, international targets for reducing fossil fuel related emissions havemotivated a push towards sustainable energy usage in both the public and privatesectors. Problematically, renewable energy sources like wind and solar can be un-reliable, a problem that has proven to be a great challenge for sustainable energygeneration, reliability and power network efficiency and stability. Efficient, cheapand reliable energy storage has long been proposed as a solution to the problemsof renewable power reliability and great efforts are currently underway to developbetter energy storage devices to allow for continued progress towards the adoptionof sustainable energy practices. At the same time, adoption of personal communi-cation devices like cell phones, tablets and computers on a massive scale as well asthe advent of wearable electronics continues to fuel demand for higher performingenergy storage devices.Recently, the conversation around energy storage has grown in volume follow-1ing the success of companies like Tesla and Solarcity that are seeking to changethe way that society interacts with and makes choices about their energy consump-tion. For one, Tesla’s electric vehicles have convinced many that electric vehicles(EVs), once thought to be far inferior to hydrocarbon fueled vehicles as a mode oftransport, can indeed perform at least as well as, and in some cases better than tradi-tional vehicles without many of the negative side-effects. This realization has seenalmost all of the big car companies invest heavily in EV research and developmentand has contributed greatly to advances in electrical energy storage technologies.The transportation industry is expected to continue to contribute to the demand forelectrical energy storage in a massive way in the upcoming decades. Electric vehi-cles are predicted to make up 35% of all new vehicle sales and consume 2,700 TWhof electricity by 2040, a figure that is equivalent to 11% of todays global energyuse [2].As well as being driven by the global sustainability movement, the success ofEVs recently has been fueled by dramatic advances in energy storage technolo-gies. The Li-Ion battery for example, which is the energy storage technology thatunderpins the majority of EVs has been the subject of rapid technological advanceover the past decades and is currently experiencing a dramatic fall in the price perkWh of power, from over $1,200 per kWh in 2010 to under $400 per kWh in 2015[3]. This fall in cost is further contributing to the demand for energy storage aslower costs facilitate the development of energy storage reliant industries like theEV industry. Despite the apparent success in recent times of the EV industry, EVtechnology has immense room for improvement. Many EVs today still have woe-fully low range with even the best EVs requiring recharging after only 300 or somiles [4]. The current industry target of 500 miles of range on a single charge2requires further development of advanced and drastically improved energy storagetechnologies with higher energy storage capacities, increased safety characteristics,lower cost and weight and using more sustainable materials.Simultaneously and also on the energy front, the cost of production for renew-able energy like solar and wind is falling dramatically as technological advance-ments appear in these sectors as well. As a result, the demand for PV cells hasbeen increasing dramatically as the price per watt has fallen over recent years.Similar advancements are changing the economics and demand for wind powerand demand is expected to continue increasing at dramatic rates for the foreseeablefuture. In the renewable energy sector, it is estimated that wind and solar powerwill add 8.6 TW of new power generating capacity worldwide in the next 25 years[2]. The success of wind and solar depends on the ability to store the electricalpower produced so that it can distributed and used efficiently and the market forthese technologies is expected to rise from 31.88 million in 2015 to 2.9 billion in2024 [5]. Similarly to EV technology, the success of energy storage in the renew-able energy industry depends on increasing storage capacities and lowering the costand environmental impact of energy storage technologies.A third driver of energy storage technologies, and perhaps one of the biggestdrivers of demand for electrical energy storage today is the portable electronic de-vices industry. Without taking into account tablets, laptops and other portable elec-tronic devices, the number of smartphones alone in circulation is expected to riseto 6.1 billion by 2020, further increasing the demand and importance of efficientelectrical energy storage [6]. With 5 billion people expected to gain access to theinternet between now and 2020, the demand for portable devices has never beengreater [7]. Additionally, recent technological advances have opened the door for3an entirely new industry that will further contribute towards the demand for highperformance energy storage devices: wearable electronics. This industry, muchlike the portable electronics industry, is demanding smaller batteries with high en-ergy density and lower cost and weight and is expected to grow to a market size ofover $35 billion USD in the next two years [8].Outside of the renewable energy industry, the transportation industry, and theportable electronics industry, energy storage remains a crucial component of manyof the world’s other industries including manufacturing and telecommunications.In high tech manufacturing for example, some $80 billion USD is lost by Amer-ican industry each year due to electrical power outages and interruptions due tounreliable power storage and transmission [9]. The progression of these industries,in conjunction with the key industries described above, is fuelling the demand forenergy storage technologies with higher storage capacities, efficiencies and cyclelives like never before. The explosion of electrical devices has seen off grid electri-cal use expand from a niche market in the early 20th century to a booming presentday industry. While there exists today a plethora of methods for storing energy,from low cost hydro storage systems to specialized fuel cells, two electrical energystorage devices dominate the portable electronics, transportation, manufacturingand renewable energy industries and will be the focus of this work: batteries andcapacitors. Batteries and capacitors have significantly different attributes and ap-plications but are similar in the way that they store energy electrochemically andcan be scaled up or down for a wide range of applications. Both of these technolo-gies are the subject of rapid technological advance and are slated to change the waythat modern society uses electrical energy on a daily basis.One factor that is facilitating the advance of these key energy storage technolo-4gies is the massive progress that has been seen in the field of nanomaterials over thepast decade. When combined with advances in materials engineering and nanofab-rication techniques, well engineered nanomaterials have the potential to contributeto the advancement of energy storage devices in a number of key ways includ-ing 1) improving energy storage capacity of devices, 2) enhancing cycle life andstability, 3) increasing charge and discharge rates and 4) improving safety char-acteristics. By carefully selecting precursor materials, significant improvementsin sustainability and cost can also be achieved. The progression of nanomaterialsengineering has allowed for significant improvements in existing energy storagetechnologies, such as the ubiquitous Li-Ion battery and may be the key to unlockingnext-generation energy storage technologies that have, up to this point, remainedcommercially non-viable.With the goals of high capacity, low cost and weight, stability and sustainabil-ity in mind, this study explores advances in nanomaterial engineering as it relatesto energy storage device development and details the development of novel andsustainable electrodes for high-performance electrical energy storage devices. Theproposed electrodes are developed from naturally abundant and sustainable bio-materials and take advantage of advanced nanofabrication techniques to create ahighly-engineered 3D framework for electrodes in energy storage devices. Thework is motivated by existing difficulties facing some of todays most promisingnext generation energy storage devices including the Lithium-Sulfur battery. Thesolutions presented in this work, while applicable to Lithium-Sulfur batteries, aremore widely applicable to energy storage devices in general and are characterizedby fabricating and testing supercapacitor devices.5Chapter 2ObjectivesAs the market for energy storage devices continues to grow under the pressureof increasing demand from the renewable energy, electric vehicle, manufacturingand consumer electronics industries, technological breakthroughs will be requiredto keep up with industry targets. Today’s Li-Ion technologies have improved somuch in recent years that the performance of this technology is approaching itstheoretical maximum (Fig 2.1). With benchmarks for energy device performancealready set by governments and industry alike, research and development effortshave turned to the fields of nanomaterials and nanofabrication to look for waysof improving device performance and unlocking next generation storage technolo-gies. Following an increasing world wide focus on sustainability, the issues ofsustainable materials and fabrication techniques have become equally importantto performance when designing the next generation of devices. In this work, ad-vances in nanofabrication techniques are applied to development of high perfor-mance electrodes from sustainable materials for use in next generation energy stor-age technologies.While incremental improvements to energy storage technologies continue to bea regular occurrence in academia, scaling up breakthroughs from the lab setting tothe commercial setting requires more than just improvements in key performance6Figure 2.1: Schematic showing the current and expected energy densities forvarious battery technologies (reproduced with permission from [10]).metrics. The commercial success of energy storage technologies depends on anumber of factors. In addition to having exemplary electrochemical performancecharacteristics, next generation energy storage devices require excellent structuraland physical properties, including low weight and high mechanical resilience andfelexibility and should use primarily sustainable component materials. Also, costof materials should be low and fabrication techniques should be inexpensive, rela-tively simple and should be highly scalable to allow for mass production. Withthese considerations in mind, this work will attempt to look at energy storagedevices with their commercial feasibility in mind, which will propose carefullythought-out and holistic designs and materials for next generation storage devices.7Of the avenues through which the development of energy storage devices isprogressing, electrode research and development is perhaps the most ready for thearrival of modern nanomaterial and nanofabrication science. Due to the complexnanoscale processes involved in storage device operation, the application of nano-materials and nanofabrication to electrodes presents an opportunity to control andto optimize these processes like never before. With the push towards better storagetechnology arriving in tandem with increased understanding of nanoengineeringtechniques, there is an opportunity to use nanofabrication practices to push en-ergy storage into the future by designing and fabricating intelligently engineeredelectrodes. With this in mind, the primary objective of this work is to develop aplatform from which high performance electrodes for applications in commercialnext generation energy storage devices can be based.In order to achieve the objectives of this work, naturally abundant and sustain-able biomaterials will be used to engineer intelligently designed electrodes witha specific goal of alleviating some of the concerns facing next generation batterytechnologies. Highly scalable fabrication techniques including electrospinning willbe studied with an eye on the future commercial feasibility of the proposed devicesand low cost precursor materials will be used whenever possible. A thorough studyof performance considerations for energy storage devices electrodes will be con-ducted and the proposed electrodes will seek to meet and exceed the performanceparameters of commonly used carbon based materials through intelligent nanoscaledesign. The proposed electrodes will be presented as a platform on which to basethe development of future and even higher performing electrodes for next genera-tion devices.8Chapter 3Literature Review and Background3.1 BatteriesThe most ubiquitous and probably the more familiar of the two major electrochem-ical energy storage technologies is the battery. According to The Freedonia Group,the global market for batteries is expected to reach $120 USD billion by 2019, upfrom $83 billion in 2015 [11]. While today, many of us rely on batteries to poweressential technologies on a daily basis, the history of the battery began in earnest inthe 17th and 18th centuries. Despite evidence of early electrochemical cells as longago as 2,000 years ago, the first semblance of a battery as we recognize it todayappeared in 1800 with the invention of Alessandro Volta’s “Voltaic Column”. Ear-lier, Volta had discovered that specific liquids have the ability to initiate chemicalreactions between dissimilar metals and instigate the flow of electrical current. Byalternatingly stacking plates of copper and zinc and separating these plates with athin sheet of cloth saturated in acid, Volta was able to generate sizeable electric cur-rent in a wire connected to the ends of the stack [12]. The result was a rudimentarybattery, the first of its kind.The Voltaic Column functioned by converting chemical energy into electricalenergy through a series of electrochemical reactions. As with all other batteries9that would come after, during discharge, the electrolyte, in this case a strong acid,reacts with the metal anode and through oxidation, releases ions, which will thenmigrate across the cell. The free electrons, which are liberated from the anode canthen travel through a wire to the cathode, where the electrons re-combine with theions in the electrolyte in a reductive reaction. An electrical device can be poweredby this flow of electrons. A thin separator, pieces of cloth or cardboard in the caseof the Voltaic Column, is usually added to prevent contact of the two electrodesand electrical shorting of the cell. A simple schematic of a battery is provided inFig 3.1.Figure 3.1: Schematic of a simple battery (reproduced with permission from[13]).Volta’s battery spurred a huge amount of technological innovation at the time10and lead, among other things, to the invention of electrolysis and to the discovery ofthe ability to isolate chemical elements. In 1802, William Cruickshank developedthe first battery suitable for mass production which consisted of a voltaic columnlaid sideways and encased in concrete [12]. In 1836, a British chemist namedJohn Daniell invented a new variant of the battery, known as the “Daniell Cell”,which relied on a zinc electrode submersed in sulfuric acid and then placed in anearthenware pot which was then submerged in a copper pot filled with sulfuric acid.The resultant cell provided a longer and more reliable current while also being saferthan the Voltaic Column. It soon became the standard battery in industries such astelecommunication [12].Over the next decades, battery technology progressed rapidly and higher per-forming batteries where introduced to the market regularly. Up until 1859 however,all batteries where primary batteries, meaning that they could not be recharged andwere usually disposed of after running out of usable charge. In 1859, a frenchphysicist named Gaston Plante´ invented the first ever rechargeable battery thatcould be charged by passing current through the cell in the reverse direction. Hislead-acid battery was made using a lead anode and a lead dioxide cathode withsulfuric acid as the electrolyte. During discharge, the electrolyte reacts with thelead anode and releases electrons. At the same time, electrons are absorbed bythe lead dioxide cathode as it too reacts with the sulfuric acid. This produces acurrent, which, if reversed, causes the chemical reactions in the cell to proceed inthe opposite direction [14]. Despite its size and bulk, Plante´’s lead-acid battery be-came very popular in applications requiring large currents and is still in use todayin industries like the automobile industry.In 1866, Georges Leclanche´ developed a battery relying on a zinc anode and a11manganese dioxide cathode dipped into a jar of ammonium chloride. At the time,the 1.4 volts that this cell could provide made it useful in applications from teleg-raphy to signalling [14]. Despite its commercial success, the Leclanche´ cell andthose before it suffered from a major drawback: they relied on a liquid electrolytewhich made the cells heavy, bulky and non-portable. In 1886, Carl Gassner in-vented the first dry cell that did not require the use of a liquid electrolyte. Instead,Gassner’s cell used plaster of Paris, mixed with ammonium chloride as a solid elec-trolyte. This, combined with a manganese dioxide cathode and a zinc anode madea portable and maintenance free battery capable of a 1.5 volts output [14]. The con-venience of this design allowed the battery to flourish and the Gassner cell becamethe first truly ubiquitous battery, allowing for the proliferation of portable elec-tronic devices, such as the flashlight. In 1899, the first alkaline battery appearedafter a Swede named Weldemar Jungner combined nickel and cadmium electrodeswith a potassium hydroxide solution [14]. Nickel-cadmium batteries (Ni-Cd) hadsignificantly better energy density compared to lead-acid batteries.By 1903, Thomas Edison had patented the nickel-iron battery in hopes of com-mercializing a more light weight and durable alternative to lead-acid batteries forelectric vehicles. Despite never catching on in the personal transportation indus-try, the design became very popular in other industries like the rail and miningindustries. Until the 1950s, Gassner’s zinc-carbon battery remained popular, de-spite its low battery life. In 1959, a researcher named Lewis Urry developed alonger lasting alkaline battery from manganese dioxide and powdered zinc. Thenew battery’s longer life quickly made it popular with consumers. In 1989, nickel-metal hydride (Ni-MH) batteries appeared which promised even longer lifespansfor smaller applications [14]. These batteries used a nickel oxide hydroxide cath-12ode like in Ni-Cd batteries but employed a hydrogen absorbing alloy as the anode.Ni-MH cells can have 3 times the capacity of an equivalent Ni-Cd cell and by 2006had all but replaced Ni-Cd cells [15].Meanwhile, much research was being done on an even more promising batterytechnology based on lithium batteries which had hit the market in 1970. Lithiumhad been identified in 1912 as being suitable for batteries due to its low density,and high electrochemical potential and energy-to-weight ratio. In 1985, a Japaneseresearch team at Asahi Chemical built the first lithium-ion battery prototype, im-proving on the existing primary lithium battery technology by increasing stabilityand adding rechargeability [16]. In 1991, the first rechargeable lithium-ion batterywas commercialized by Sony and utilized a lithium cobalt oxide cathode [17].The introduction of the lithium-ion (Li-ion) battery revolutionized the recharge-able battery market with its portability, high energy density, low self-discharge rateand small memory effect. In a Li-ion cell, lithium ions travel travel back and forthbetween the positive and negative electrodes during charging and discharging. Typ-ically, these cells employ an intercalated lithium compound as the anode, usuallygraphite, and a lithium oxide cathode. During charging, positively charged lithiumions flow from the cathode and intercalate into the graphite anode. Unlike someother battery types, this intercalation mechanism does not involve a chemical reac-tion. During discharge, the opposite process occurs and lithium ions are releasedfrom the anode and can travel freely back to the cathode. Over the ensuing years,Li-Ion battery technology saw continuous improvements as researchers developedand discovered new electrode and electrolyte materials, some of which are detailedin the following paragraphs.13LCO Batteries: Since Sony’s introduction of their lithium cobalt oxide (LCO)lithium-ion battery in 1991, the popularity of this system has exploded and the LCObattery has quickly become the most popular battery for consumer and portableelectronics. LCO batteries utilize a Cobalt-Oxide cathode and a graphitic carbonanode. Compared to nickel-based alkaline batteries, LCO cells are smaller, have ahigher energy density, and can operate over a wider range of temperatures. Addi-tionally LCO batteries lose a fraction of the charge per month compared to nickel-based cells: 5-10% compared to more than 30% [18]. LCO batteries do come witha downside however. More expensive to make than nickel-batteries, LCO cells aresusceptible to thermal runaway if over charged and so require the addition of pro-tective circuitry. This circuitry can accelerate self-discharge and limit dischargerate. Additionally, LCO cells accumulate internal resistance as they cycle and sohave limited lifetimes.LFP and LMO Batteries: In 1996, the group of John Goodenough at the Uni-versity of Texas invented a new variation of the Li-ion battery relying on lithiumiron phosphate (LFP) as the cathode. This material is much cheaper than lithiumcobalt oxide as well as being inherently more stable. LFP batteries have a lowerenergy density than LCO cells but have longer cycle lives, a higher discharge rateand make for a battery that is 14% cheaper than an equivalent LCO cell [18]. Athird popular Li-ion cell type is the lithium manganese oxide (LMO) cell whichwas also developed in 1996. This variation has very high thermal stability and lowinternal resistance, making it ideal for applications requiring fast recharging andhigh currents - they can output 20-30 amps safely. Problematically however, thesecells have very low energy density and are only able to hold 1200 mAh of power in14a typical 18650 cell, about half of the capacity of LCO cells [18].Figure 3.2: Specific and volumetric energy density of some of the establishedsecondary battery types (reproduced with permission from [19]).Li-Ion batteries drastically improved performance of many of the portable de-vices that we are familiar with today and represented a huge step forwards in thedevelopment of battery technology. Fig 3.2 presents a summary of the perfor-mances of some of the major secondary battery types developed in the last century.Despite the radical change brought about by the Li-ion battery, the technology isstarting to approach its theoretical limits and the need for improvements in batterytechnology has become apparent (see Fig 2.1). The incredible progress of portableelectronics, the rise in the use of electric vehicles and the advance of renewable en-ergy technologies are all demanding increases in energy storage performance and15capabilities beyond what is commercially available today. In November 2013, theU.S. department of energy announced a $120M effort to develop next generationbattery technology with a goal of extending battery life by 5 times that of 2012batteries [20]. This effort, others like it, and pressure from key industries has leadto intensive research on next generation battery technologies, a few of which arehighlighted in the following paragraphs.Lithium-Silicon Cells: Lithium-Silicon cells employ a silicon anode which can,in theory, hold more than ten times the number of lithium ions than typical graphiteanodes, leading to an extraordinarily high theoretical specific capacity of 4200mAh/g compared to 372 mAh/g for graphite. This is because silicon atoms canbind up to 4.4 lithium ions compared with one ion per 6 carbon atoms in graphite.Problematically however, silicon anodes have been found to expand to over 3 timestheir original volume during anode lithiation. This expansion can lead to cracksand crumbling in the anode as well as detachment from the current collector. Re-search stage cells have been shown to lose almost all of their capacity in as fewas 10 cycles. Additionally, lithium-silicon cells are plagued by the formation ofa solid electrolyte interface layer (SEI), which, due to cracking caused by volumeexpansion, can continue to grow and eventually leads to very high internal cell re-sistance and low efficiency. Despite these drawbacks however, efforts to optimizelithium-silicon cell performance continue.Lithium-Air Cells: Lithium-Air batteries are another highly promising next gen-eration battery technology. In lithium air batteries, oxygen reacts with lithium ionsdissolved in the electrolyte at the cathode to form lithium peroxide. In the reverse16reaction, lithium peroxide disassociates to produce lithium ions and oxygen. Thissystem is a precipitation-dissolution system and is very unlike the intercalation re-action that typical Li-ion batteries rely on. The major advantage of the lithium-airbattery is that, because oxygen is found abundantly in air, it does not have to bestored on board the battery and instead, can be supplied to the battery by flowingair into the cell. This means that theoretically, cells can achieve an energy den-sity of 11,140 Wh kg−1 [21]. There are significant challenges associated with thisbattery design however. Moisture and contamination has to be filtered out of theair in order to prevent impurities from entering the cell and degrading battery per-formance. Even more problematically, the precipitation-dissolution can be slowand so lithium-air batteries are known to have low discharge and charge rates. Ad-ditionally, there is a mismatch in the charge vs. discharge voltages in these cellsmeaning that lithium-air batteries can be very inefficient.Lithium-sulfur Cells: Lithium-sulfur (Li-S) batteries have been the subject of in-creasing research recently and are thought to be the next generation battery technol-ogy that is closest to commercialization. This increase in activity has been fuelledby the progression of the electric vehicle industry, an industry that is demanding theprogress of battery technology to allow for electric vehicles with ranges of around500 miles+ per charge. This type of range requires a battery with a cell level spe-cific energy of 350-400 Wh kg−1 [22], a figure that is significantly higher than the200 Wh kg−1 that current Li-ion cells can provide. In order to meet the energy den-sity target of 400 Wh kg−1 by 2017, a target set officially by the US Departmentof Energy [10], new battery technologies must be explored and Li-S batteries area prime candidate for meeting such a target because of their high theoretical ca-17pacities and other highly desirable characteristics. As indicated in Table 3.1, Li-Sbatteries have a high theoretical specific energy of 2,567 Wh kg−1, almost 5 timesthat of Li-Ion cells. In practice, the predicted possible energy density is likely to becloser to 500 Wh kg−1. Additionally, the active ingredient in Li-S cells is sulfur, anaturally abundant, widely available and low cost material, making Li-S cells rel-atively cheap to produce and more environmentally sustainable than other batterychemistries. The combination of high energy density, low-cost and sustainabilityof component materials makes Li-S cells worth exploring further.Battery Cell Voltage (V) Theoretical Specific Energy(Wh kg−1)Li-ion 3.8 387Li-S 2.2 2,567Li-Air (non-aqueous) 3.0 3,505Li-Air (aqueous) 3.2 3,582Table 3.1: Voltage and theoretical specific energy of lithium based energy-storage devices (adapted from ref [10].)Li-S batteries are commonly constructed using a lithium metal anode and asulfur cathode. During discharge, lithium metal is oxidized, freeing up lithium ionsand electrons at the anode. The lithium ions then travel through the electrolyte andreacts chemically with sulfur and with the free electrons, which have by then passedthrough the external circuit. The reaction is a reduction reaction and produceseither lithium sulfide (Li2S) or lithium disulfide (Li2S2). The overall reaction is asfollows:Li−−→ Li+ + e–S8 +16Li+ +16e– −−→ 8Li2S16Li+S8 −−→ 8Li2S18Li-S cathodes usually contain sulfur in its most stable form, S8. As previouslymentioned, during charge, lithium ions interact with stable sulfur to form one oftwo solid products in the cathode: Li2S2 and Li2S [23]. This reaction is far fromsimple however and is still being fully understood. One of the challenges of Li-Sbatteries arises from this reaction and is known as the polysulfide shuttling effectwhich arises because, when S8 is reduced by lithium ions, it goes through a seriesof reduction reactions before it reaches its final solid form. In between, it passesthrough various polysulfide forms, beginning with long chain polysulfides, Li2S8and Li2S6, before passing through shorter chain forms like Li2S4, Li2S2 and Li2S[24]. Problematically, some of these polysulfides are highly soluble in typical or-ganic electrolytes, meaning that they can dissolve into the electrolyte and migratearound the cell. It has been found that higher order polysulfides tend to migrate tothe anode, where they are reduced to lower order polysulfides that migrate back tothe cathode before being oxidized to high order polysulfides again. This process isreferred to as the polysulfide shuttling effect.The shuttling effect, shown schematically in Fig 3.3, can repeat itself endlesslyand can lead to loss of active material in the cell, dramatically decreasing efficiencyand cycle life [24]. Additionally, dissolved polysulfides have a tendency to platethemselves onto the surface of the anode, forming insoluble products in the process.This too can significantly hinder the free flow of ions to and from the anode anddrastically and negatively impact the performance of the battery [25]. Finally, theinsulating nature of sulfur poses an additional problem for Li-S cells. Becauseelectrons can not flow freely from reaction sites within the cathode, Li-S cathodesneed to incorporate conductive additives so that the flow of electrical current canbe established. this can have a negative impact on the overall specific capacity and19volumetric capacity of Li-S cells.Figure 3.3: Schematic of the polysulfide shuttling effect in Li-S batteries (re-produced with permission from [26]).In order to combat these limitations, groups around the world have adopteda wide variety of approaches. One popular avenue of research is to develop newelectrolyte chemistries with the goal of decreasing polysulfide solubility within theelectrolyte. This approach has lead to experimentation with non-aqueous, ionic,solid polymer and glass-ceramic electrolytes among others [27]. Other groups are20focusing on novel separator designs to help limit the flow of polysulfides and haveeven begun introducing what is known as an “interlayer” [28]. This additionalconductive layer between the sulfur cathode and the separator is designed to helpblock the diffusion of polysulfides and increase the electrical contact to the sulfurcathode. While this strategy can provide benefits in terms of cycle life, it can dras-tically decrease the specific and volumetric capacities of the cell. Still other groupsare looking at ways of altering the chemistries at the electrodes to improve perfor-mance and safety. For example, there is much research into the idea of using a Li2Sbased cathode instead of one relying on elemental sulfur [29]. This strategy allowsfor the use of non-lithium anodes which may be safer than conventional lithium-based anodes. Finally, and perhaps one of the more popular research directions isthe search for new and novel cathode materials and structural designs that mighthelp Li-S cells circumnavigate some of the problems that they currently face.Within the battery market, we are starting to see the appearance of commer-cial Li-S batteries with Oxis Energy appearing to lead the way. The UK companyrecently announced a Li-S cell that they claim achieves over 400 Wh kg−1. Evenwithin the context of this apparent success, cyclability problems caused by theshuttle effect continue to plague even the best experimental Li-S cells. There re-mains much improvement to be done in the field before Li-S batteries can achievecommercial success on a wider scale and further investigation and experimenta-tion on electrode materials and geometries is necessary. Despite these difficulties,the sentiment within the research community is that, with the progression of ad-vanced nanofabrication techniques and an ever increasing understanding of mate-rials used in battery research coupled with drastically increased interest in the fieldof rechargeable batteries, we may be getting closer to solving the problems asso-21ciated with the technology. For these reasons, the development of new materialsand intelligently engineered cathode structures is hugely exciting. Fig 3.4 presentsa summary of the characteristics and development trajectory of recent battery tech-nologies.Figure 3.4: Battery technology roadmap and characteristics (reproduced withpermission from [30]).3.2 SupercapacitorsLike batteries, research on supercapacitors has seen huge growth over the lastdecade as teams continue to push for better ways of storing and using electricalenergy. Indeed, there are many parallels between high performance supercapaci-tors and batteries, and in particular, in the design of the electrodes for both sys-tems. Electrochemical capacitors (EC), otherwise known as supercapacitors or ul-22tracapacitors, have been the subject of rapid technological advancement since theirappearance in 1957 in the labs of General Electric (GE). In 1962 and improvingon GE’s experimental designs, the Standard Oil Company of Ohio (SOHIO) devel-oped the first supercapacitor in the format that we are familiar with today. SOHIOhad high hopes for the newly developed technology and sought to develop it for ap-plications in power rectification. By 1971, these efforts had failed and the companyhad accumulated too much debt from building the Alaska Pipeline to continue withdevelopment of their supercapacitors. In 1975, SOHIO licensed their EC technol-ogy to the Nippon Electric Company (NEC) of Japan, who, in 1978, began thefirst commercial production of their NEC SuperCapacitor, which was marketed asa back-up device for clock chips and computer memories [31]. Since then, super-capacitors have progressed rapidly through many generations of designs and havefound applications in a huge variety of industries, from wireless communication totransportation.The best of the early commercial devices had voltage ratings of a few voltsand low capacitance values ranging from fractions of farads up to a few farads. Incontrast, today’s supercapacitors can achieve performance ratings of hundreds ofthousands of farads and operate at voltages of thousands of volts, demonstratingthe level of progress that the field has seen over the past 40 years [31]. Shortly af-ter the release of the first commercial supercapacitor, Panasonic entered the marketwith its “Goldcaps” brand, followed shortly by ELNA’s “Dynacap” supercapacitors[32]. Because of the high internal resistance in these early supercapacitors, theywere limited to low current applications like powering computer chips and backingup data. In 1982, the Pinnacle Research Institute (PRI) introduced its ultracapac-itor, the first such double-layer capacitor capable of being used in high power ap-23plications. Designed for military applications, the PRI ultracapacitor incorporatedelectrodes made from metal-oxide and its release triggered renewed interest in thetechnology by the US Department of Energy. This lead to the formation of theUltracapacitor Development Program at Maxwell Laboratories in 1992 [33].Supercapacitors are double-layer capacitors with low voltage limits and withhigh capacitance values. In a supercapacitor, two electrodes are immersed in anelectrolyte containing both positive and negative ions. Upon application of a volt-age across the two electrodes, the positive and negative ions migrate across theelectrolyte in different directions and gather at the surface of the electrodes, creat-ing an electric double layer at each of the electrodes as shown in Fig 3.5. Super-capacitors can be classified as either electrical double-layer capacitors (EDLCs),which rely purely of the accumulation of electrostatic charge at the electrode/elec-trolyte interface, or as pseudocapacitors, which produce electricity via fast andreversible redox reactions at active sites on the electrode [34]. In practice, manysupercapacitors display characteristics of both types.Supercapacitors have numerous applications and have a variety of advantagescompared to other energy storage devices. For one, supercapacitors have very longcycle lives, often operating for 10 to 20 years without losing significant amounts ofcapacity. Additionally, supercapacitors can function well in extreme environmentsand are able to provide energy at temperatures reaching -40◦C. The primary advan-tage to using supercapacitors however is their ability to provide high load currentsmeaning they can charge in mere seconds. Supercapacitors have very high powerdensity when compared to batteries. On the other hand, supercapacitors have rel-atively low energy density and so can not provide power continuously for longperiods of time. A typical supercapacitor has a cell voltage of around 2.7 V and24Figure 3.5: Schematic of a supercapacitor during charge and during dis-charge (reproduced with permission from [35]).can easily be connected in series to boost voltage. Supercapacitors are in use todayin a huge variety of industries, from consumer electronics to medical devices andand the market for supercapacitors is expected to grow quickly in the next decade.Since the conception of the Ultracapacitor Development Program, there hasbeen vast improvements in supercapacitor performance in all performance metrics.While there has been some work in the research community focused on improvingthe break-down characteristics of electrolytes in order to raise the operating volt-age of supercapacitors, the lion’s share of efforts have been dedicated to improvingmaterials and material properties of supercapacitor electrodes. The energy storagecapability of a supercapacitor is referred to as capacitance (C) and is often mea-sured in Farads per gram. Capacitance is a function of the surface area (A), thepermittivity (ε) and the distance between plates of the cell (d) (Eqn3.1).25C = ε ∗ (A/d) (3.1)Following along with this relationship, high performance supercapacitors usu-ally include electrodes with very high specific surfaces areas which helps to in-crease the amount of charge that a cell can hold. High surface area is often achievedthrough developing highly porous electrode materials. Additionally, supercapaci-tor performance can be linked to the electrical conductivity of the electrodes, mean-ing that highly conductive materials are preferable for high performance superca-pacitors. For these reason, modern supercapacitors are generally constructed usingactivated carbons and other carbon derivatives like carbon-fibre cloth and carbonaerogels. By far the most widespread source of carbon for commercial supercapac-itors today is activated carbon, which is usually derived from coconut husk.The materials currently being used for supercapacitors have high surface areabut continue to suffer from low mesoporosity, meaning that electrolyte accessibil-ity is still weak within the electrode. This can lead to low energy density in thedevice, which, when considered in parallel with the poor electrical conductivity ofmany carbon materials and therefore high internal resistance, means today’s super-capacitors continue to deliver much lower power density then theoretically possi-ble. Additionally, typical supercapacitor electrodes require the addition of bindingagents and conductive agents as well. A typical commercial supercapacitor is madeup of around 85% activated carbon, 10% conductive additive (carbon black for ex-ample) and 5% binding agent (PTFE for example) [36]. The addition of theseagents further decreased volumetric power and energy density while contributingto increased electrode weight and cost. Currently available supercapacitors based26on activated carbons can provided power densities in the range of 1-2 kW kg−1and energy densities of between 4 and 5 kWh kg−1 with specific capacitances of25-30 F per gram of activated carbon [37][38].With a view on developing commercially viable solutions to improving su-percapacitor performance, it is important to understand the costs typically involvedwith such devices. A typical mid size commercial supercapacitor can provide about350 Farads [38]. Assuming a performance of 30 F g−1, this means that a mid sizecommercial device contains about 11.66 g of activated carbon, 0.69 g of bindingagent and 1.37 g of conductive additive. At a rough price of $15 USD/kg for ac-tivated carbon from coconut husk [39], $15 USD/kg for PTFE binding agent [40]and $0.08 USD/kg for carbon black [41], this equates to an electrode material costof approximately $0.19 USD per device with additives making up about 5.5% ofthe cost.The limitations in supercapacitor performance indicate that special attentionneeds to be paid to increasing porosity and conductivity in supercapacitor elec-trodes. It is noteworthy that the characteristics that make for a high performance su-percapacitor, namely high surface area, high porosity, high conductivity and goodstructural resilience and integrity, are also characteristics that research teams arelooking for in high performance battery electrodes and in particular, high perfor-mance Li-S batteries. With these comparisons drawn, it should be clear then thatthe search for intelligently engineered, sustainable and high performance materialsand geometries for electrodes applies equally to the development of supercapaci-tors as it does for next generation Li-S batteries. It follows that the performance ofsupercapacitor devices developed with these electrodes should give a good idea ofthe potential performance of these electrodes in applications in Li-S batteries.273.3 Carbon Nanofibre CompositesWithin the umbrella of electrode material research and structure design, there is awide variety of strategies aimed at developing high performance and intelligentlydesigned electrodes for batteries and for supercapacitors. In the next few pages,strategies for developing structures and materials for high performance electrodesfor both Li-S batteries and supercapacitor electrodes will be discussed. In the caseof Li-S batteries, proper electrode design should seek to contain polysulfide for-mation and migration while increasing electrode conductivity and while seekingto accommodate the most active material possible in order to optimize cell ca-pacity. This can be achieved by designing highly conductive materials with highporosity and high surface area. Because carbon is naturally abundant and widelyavailable, a wide variety of intelligently designed 3D and free-standing carbonframeworks have been proposed to achieve high sulfur loading and facilitate chargetransport. For example, microporous carbon spheres were used to trap elementalsulfur within the pores and achieved a loading of 42% and stable cycling up to500 cycles [42]. Mesoporous carbon with tuneable porosities have been used ina similar fashion and have achieved sulfur loading up to 50% and a capacity of1390 mA h per gram of sulfur [43]. Other carbon materials like multi-walled car-bon nanotubes have been used to create composite cathodes with similar capacitiesof 1352 mA h g−1[44], and graphene has been employed to wrap and encapsulatesulfur particles, leading to a stable specific capacities of 600 mA h g−1 over 100 cy-cles [45]. Other strategies include the use of conductive polymers, carbon-foams,carbon blacks, hollow carbon, carbon nanofibres, carbon cloth and hierarchicalporous frameworks. The different strategies used to encapsulate sulfur into high-28performance Li-S cathodes are summarized in Fig 3.6.Figure 3.6: Different strategies for encapsulating sulfur in 3D frameworks(adapted with permission from [46]).Because supercapacitor performance is tied to some of the same characteristicsas Li-S electrode performance, it is not surprising that many of the same strategiesfor developing high performance Li-S electrodes have been applied to supercapac-itors. In particular, 3D architectures with high specific surface areas are regularlyemployed for supercapacitor electrodes. Typically, a carbon material with a sur-face area of 1000 m2 g−1 could achieve a theoretical specific capacitance of be-tween 200-500 F g−1. For example, Zhu et al. developed supercapacitor electrodesfrom activated graphene and achieved a gravimetric capacity of 150 F g−1 [47]. Inanother work, electrodes were prepared from carbon nanotubes and achieved high29capacitances of 170 F g−1 [48]. Like for batteries, functioning research stage super-capacitors have been prepared from a variety of other materials including carbonblacks, carbon aerogels and carbon/polymer composites.Of the myriad of strategies for developing intelligently engineered 3D elec-trode architectures for Li-S batteries and for supercapacitors, the use of carbonnanofibres (CNFs) stands out as one of the more promising. CNFs are well suitedfor use in electrodes because of their structural integrity and high conductivity.Additionally, they can be prepared in interconnected mats with varying degreesof connectivity, which, when used in Li-S electrodes, have the ability to containlarge amounts of sulfur while maintaining interconnected and conductive pathwayswithin the cathode. On top of its contribution to the conductivity properties of theelectrodes, this interconnectivity is attractive because it contributes to the struc-tural integrity of the electrode and can mean that electrodes can be prepared binderfree and without a current collector. Because of this interconnectivity, structuralrobustness and high conductivity, carbon nanofibres are suitable for many differenttypes of energy storage electrodes including for supercapacitors and not just forLi-S specific applications. The use of carbon nanofibres as an electrode materialis a popular strategy and high performance electrodes have been demonstrated forlithium-air batteries [49]. To date, this approach has been used sparingly for ap-plications in Li-S batteries but there does exist examples of successful applicationsin the literature. For example, hollow carbon nanofibres were used to encapsu-late sulfur and deliver a discharge capacity of 1170 mA h g−1 and retained goodperformance even after 200 cycles [50]. In another example, a nano-carbon/sulfurcomposite used carbon nanofibres as the electrical conductor and achieved a sim-ilar initial discharge capacity of 1200 mA h g−1 and good stability after 50 cycles30[51]. In supercapacitor applications, the use of carbon nanofibres is already quitepopular with many devices being demonstrated to date [52].While generally structurally strong and highly conductive, ECNFs suffer froma specific surface area that is typically lower than those of carbon nanotubes orgraphene (240 m2 g−1 vs 1315 and 909 m2 g−1 respectively)[53–55]. In the case ofLi-S composite cathodes, high specific surface area of the supporting carbon matrixcan benefit performance by providing additional interfaces between sulfur and theconductive framework. Additionally, it has been demonstrated that high porosityin carbon composite cathodes can help to capture the various polysulfide speciespresent during cycling and can facilitate the passage of ions in the cathode [56].These benefits can lead to increased long-term cycling stability and to reducedinternal cell resistance. In supercapacitor applications, both electrical double layercapacitance and pseudocapacitance rely on surface effects at the electrode and sohigh surface area is necessary for supercapacitors requiring high charge capacity.The low specific surface area and low porosity of traditional carbon nanofibres hashampered carbon-nanofibre based electrodes in the past.To solve this problem of low surface area and porosity, porous CNFs havebeen developed using a variety of different methods. Recently, porosity in nanofi-bres has been achieved through chemical treatment after electrospinning. Wang etal. demonstrated porous fibres made using this method with specific surface ar-eas reaching 699 m2 g−1 and developed supercapacitor devices with capacitancesof 170 F g−1 [57]. Another popular method is to electrospin blends of immisciblepolymers followed by selective removal of one of the polymers during the heattreatment step [58, 59]. Abeykoon et al. prepared fibres in this way from PANand PMMA, and achieved a very high specific surface area of 2419 m2 g−1 with a31capacitance of 140 F g−1 for their supercapacitor devices [60]. Similarly, Tran andKalra used PAN and Nafion to produce fibres with surface areas of 1600 m2 g−1[61]. Finally, highly porous ECNFs have been created by electrospinning com-bined with a template method in which sacrificial polymers are dissolved from thefibres after electrospinning. Using this method, Wang et al. dissolved PVA fromPAN fibres and obtained nanofibres with specific surface areas of 1232 m2 g−1 anda capacitance of 202 F g−1 at low scan rates of 2 mV s−1 [62].Examples of porous carbon nanofibres for Li-S specific applications are fewand far between. In one example, Ji et al. synthesized porous carbon nanofibres foruse in Li-S cathodes using the PAN/PMMA template method and loaded the pre-pared fibres with sulfur via a solution-based chemical reaction-deposition method[63]. The resulting cathodes achieved a sulfur loading of 42% and a discharge ca-pacity of 1400 mA h g−1 with promising signs of cycling stability after 30 cycles.In another work, Zhou et al. doped microporous polypyrrole based nanofibres withnitrogen to achieve Li-S cathodes with a sulfur loading of 60% and a dischargecapacity of 1532 mA h g−1 [64]. While these works remain promising, the issuesof cost and sustainability remain to be considered in these works. There remainssignificant opportunity for developing high performance Li-S batteries from porouscarbon nanofibre based electrodes.3.4 ElectrospinningCarbon nanofibres can be made in the research setting in a number of highly ef-fective ways, including by phase separation methods, by synthesis by template andby polymer drawing [65, 66]. Despite the success of these methods, there remainsgreat challenges in adapting these research-scale methods to the industrial scale32[67]. Of the ways of fabricating nanofibres, electrospinning stands out as a fab-rication method that suits large-scale production well. In the emerging field ofnanofabrication, electrospinning has become a popular and widely used method toobtain large quantities of carbon nanofibres with a high degree of control duringfabrication. Electrospun carbon nanofibres (ECNFs) have a high surface area tovolume ratio and have been used in a variety of applications, from energy storageto clothing. Significantly, ECNFs can be spun from a diverse set of materials in-cluding polymers, semiconductors, ceramics and polymers and are made as verylong continuous fibres making them ideal for applications in electronics where con-ductivity is of high importance [68].The beginnings of modern day electrospinning technology first appeared inthe 19th century, but it wasn’t until the early 20th, with the advent of high volt-age power sources and soluble polymers, that patents relating to electrospinningfirst began to appear. In the early 90’s electrospinning saw a huge rise in popular-ity after Doshi and Reneker demonstrated its use in the preparation of nanofibres[69]. The method that they proposed and that is still in use today is extremely at-tractive for nanofibre production in experimental settings because of the simplicityin its design. Specifically, electrospinning in its most basic form requires only ahigh voltage power supply, a collector and a syringe with a flat tip. As the 90’sprogressed, the research community experimented with this new method of nanofi-bre fabrication and much work was published on the topics of characterization ofelectrospun materials and analysis of optimal electrospinning conditions. Today,electrospinning as a process is relatively well understood and groups are movingtowards the fabrication of devices using electrospun materials.The electrospinning process relies on charge repulsion in an electroactive poly-33mer induced by an external magnetic field. In a modern electrospinning apparatus,a small droplet of viscoelastic solution is held at the end of a spinneret, usuallya flat-tipped needle on a syringe. In its rest state, the droplet is affixed to the tipof the syringe by its own surface tension. When the device is switched on andan electric field is applied, charge is induced at the surface of the droplet. As themagnitude of the electric field is increased, additional charge is induced, creatingmutual charge repulsion which causes a force directly opposite to that of the sur-face tension. With further increase in the intensity of the external field, the dropelongates along a uniaxial axis creating what is referred to as the “Taylor Cone”[69]. At a certain critical field intensity, the force due to the surface tension is over-come by that of the repulsive electric force, causing the charged jet of solution to beejected from the tip of the needle towards a collector. As the jet of solution travelsthrough the space between the syringe and the collector, the solvents present in thesolution have the opportunity to evaporate providing they are volatile enough. Bythe time the jet reaches the collector, what remains is a long continuous chargedpolymer fibre which collects in a random pattern on the collector, thus producing anon-woven and highly interconnected mesh.The process, while conceptually relatively easy to understand, is effected by anumber of different electrospinning parameters. In particular, the properties of theresulting nanofibres are influenced by the viscosity, the conductivity, the surfacetension, the magnitude of the electric field, the atmospheric humidity, the hydro-static pressure in the needle, the temperature, and the air velocity in the chamberamong other things. Electrospinning can be tailored to produce a variety of fibresizes ranging from less than 0.05 microns to above 5 microns in diameter and over100 different types of electrospun nanofibres have been reported [71]. Addition-34Figure 3.7: Typical electrospinning setup (reproduced with permission from[70]).ally, modern applications of electrospinning have yielded fibres with tuneable crosssectional shapes and have even provided the ability to spin composite fibres withmore than one polymer. A typical electrospinning set-up is shown in Fig 3.7.The jet produced from the tip of the spinneret can be divided into three mainportions during electrospinning which are visualized in Fig 3.8. In the cone jet,or the Taylor cone, which is the cone jet portion of the stream directly connectedto the spinneret tip, the viscoelastic solution begins to pick up velocity as it exitsthe spinneret and is pulled by the electrostatic force. Once the critical voltage hasbeen reached, the solution can break free of the cone jet and is pulled out into astable jet region. In this jet, the electrostatic forces continue to stretch the materialand the mass density of the jet remains constant along the stable jet region. As35Figure 3.8: Close-up of the structure of the electrospinning jet (adapted withpermission from ref [72]).the electrostatic forces pull the charged particle in the jet towards the collector,the elongational velocity in the jet resists the motion, causing the jet to elongateas the solvent evaporates from the jet. The elongation and the evaporation of thesolvent dynamically alter the viscoelastic parameters in the jet until eventually theradial forces caused by the electric charges in the jet become larger than the forcescausing the jet to maintain its cohesion and uniaxial direction. Once this occurs,a spiral jet is formed. The forces present during this stage cause the spiral jet tomove in a very haphazard and whipping motion. It is during this whipping motionthat the fibres are stretched into their final nanometer scale diameter as they aredeposited randomly on the collector [73]. By the time the fibres have depositedonto the collector, the solvent has, for the most part, evaporated and what is left is36a continuous non-woven mat of polymer.One of the most important parameters to consider during the electrospinningprocess is the surface tension associated with the solution being spun. It is this ten-sion that causes a liquid to resist forces that try to pull against it and is a functionof the level of cohesion of molecules at the surface that face the external atmo-sphere. Within the bulk of the liquid, individual molecules are pulled equally in alldirections by the forces associated with the molecules that surround them. At thesurface however, the molecules there are subject to bulk forces in only one direc-tion, causing a them to be pulled towards the bulk liquid and form a surface. Thesurface molecules contract towards the bulk of the liquid which in turn, causes acertain amount of pressure to build up in the liquid. During electrospinning, inter-nal repulsive forces are created and when these forces exceed the surface tensionforces, the liquid is expelled in a jet. The surface tension of a liquid therefor playsa big role in determining the shape of the jet and, if not balanced correctly, canlead to formation of beads on the jet. If the surface tension is too low, as is oftenthe case if the concentration of polymer in solution is too low, then the solutioncan agglomerate and result in beads on the fibres. Higher concentrations of poly-mers in the spinning solution tend to decrease the chance of bead formation. Beadformation can sometimes be rectified by lowering the voltage or by increasing thedistance between the spinneret and collector [74].As well as surface tension, viscosity plays a key role in the eventual shape ofthe electrospun fibres. Since viscous forces tend to resist the pull from the electro-static forces, solutions with higher viscosities have been shown to result in fibreswith larger diameters while low viscosity solutions lead to thinner fibres. Addition-ally, if the viscosity is too low, fibre formation can be inhibited all-together leading37to the development of a spray rather than a jet [75]. Along with viscosity, solu-tion conductivity and the dielectric properties of the solvent used can greatly effectthe final diameter of the electrospun nanofibres. With a highly conductive solutioncomes a greater amount of charge accumulation in the jet, leading to higher repul-sive forces and therefore a lower critical voltage. This can lead to more uniformfibres which can be further made uniform by increasing the number of ions in so-lution [68]. Adding to this effect, high solution dielectric properties can furtherincrease the repulsive forces within the jet which in turn will increase the speedof the jet as it leaves from the spinneret. This has been shown to reduce beadformation [76].Finally, two more factors that can effect the final dimensions of the preparednanofibres are applied voltage and the feed rate of the electrospinning solution. Ahigher applied voltage leads to higher accumulation of electric charge inside thejet. As previously mentioned, it is when the repulsive forces inside the jet exceedthe forces due to the surface tension that the jet is expelled from the spinneret. Ahigher applied voltage means that the jet is released sooner, leading to a shorterTaylor cone. Additionally, the increased repulsive forces means there are higherelongational forces within the jet. This leads to faster solution movement and inturn, smaller fibre diameters [77]. On the other hand, the increased forces meansthat the jet reaches the collector sooner and it has been demonstrated that, for cer-tain high voltages, this can result in less time for elongation and so thicker fibres[78]. Beyond this, high voltages can lead to splitting of the jet, which again, maylead to thinner fibres [79]. Solution feed rate can also lead to changes in the finaldiameter of the prepares nanofibres with higher feed rates resulting in increasedfibre diameters and lower feed rates leading to thinner fibres.38Due to its ability to continuously output high amounts of high quality nanofi-bres, electrospinning has been identified as an ideal candidate for scaling up theproduction of research quality nanofibres for industrial and commercial applica-tions. With the progress in electrospinning research, a number of high output de-vices featuring high controllability of the parameters discussed above have beendeveloped. One example of such a device, is the Katotech Nanofibre Electrospin-ning Unit. These devices, while still designed for use in the research setting, aredemonstrating the scalability of the technique and have allowed for further under-standing of the effect of the electrospinning parameters on the final nanofibres.With the development of devices relying on electrospun nanofibres, we can expectfurther development of industrial and commercial scale electrospinning units. In-deed, companies like Donaldson Torit DCE are already operating commercial scaleelectrospinning units.In the research setting, electrospinning has been used successfully for appli-cations in sensors, biomedical devices, filtration, textiles, magnetic devices andenergy storage among many others [80]. For energy storage applications, electro-spinning is particularly appealing as an electrode fabrication technique due to itsability to create highly interconnected and conductive mats of carbon nanofibreswith highly tuneable properties and from many different precursor materials. Theuse of ECNFs in energy storage has extended from fuel cells, to battery anodesand cathodes, to supercapacitors, to electrolytes and even to separators. Addition-ally, these applications have taken advantage of the simplicity of the technique toengineer highly tuneable components from a wide variety of materials, from Poly-acrylontrile (PAN), to polyvinylpyrrolidone (PVP) to Nafion. The versatility ofelectrospinning, the advantages associated with the technique and the nanofibres39that it can produce make it highly suitable for nanofabrication in the field of en-ergy storage and the search for the right combination of parameters and materialscontinues to be important for developing revolutionary energy storage devices.3.5 LigninTo date, the large majority of research on ECNFs has focused on PAN derived fi-bres due to the ease of electrospinnability and favourable electrical and structuralproperties provided by the organic polymer. Despite its chemical and physicalsuitability, PAN is expensive and, because it is a derivative of petroleum, is nonsustainable. Recently, increased efforts have been dedicated to searching for inex-pensive, naturally abundant and environmentally sustainable sources of precursorcarbon for ECNFs. This is particularly important for applications in energy storagewhich, as a field is moving towards more environmentally friendly and sustain-able practices. While some groups have had success producing CNFs from highlysustainable precursors such as fungi [81], seaweed [82], eucalyptus and potatoes[83], one of the more promising materials is a material called lignin. As a majorcomponent in the cell walls of many plants, lignin is, after cellulose, the secondmost abundant renewable carbon source on earth and comprises around 30% of theweight of dry wood [84]. lignin is a byproduct of the pulp and paper industry andclose to 50 million tons of lignin are generated as waste worldwide each year [85].Despite traditionally being viewed as a waste product, developing environmen-tal waste regulations are advancing the appeal of the use of lignin in commercialproducts in industries ranging from energy production to agriculture. For example,Volkswagen is considering using lignin based carbon nanofibres in the productionof automotive components and can achieve a 50% reduction in the cost of manu-40facturing relative to the use of PAN [86]. There are a number of different varietiesof lignin, each a byproduct of different pulping processes. While each lignin vari-ety differs in terms of its chemical and structural properties, Lignosulfonates havelong been the dominant lignin variety in the commercial sector, often being used inthe construction and bitumen industries and to develop food and feed ingredients.A second popular type of lignin is called Kraft lignin which is used for a varietyof commercial applications. Kraft pulping is the most common form of chemicalpulping in use today in the pulp industry and relies on a strong sodium sulfide cat-alyst to remove the lignin from the surround cellulose fibres in which it is found.Once the pulping stage is complete, the cellulose is bleached to further remove anyremaining lignin. Often, the lignin that is removed from the cellulose is burned atthe mill to supply energy for the functioning of the mill but just as often, it is dis-carded as waste. Kraft lignin is distinct from lignosulfonates and does not containthe sulfonate groups present in lignosulfonates. Kraft lignin is used primarily inapplications such as dyes and pesticides and currently, is only used commerciallyto a very limited degree [87].In addition to environmental waste regulations, the use of lignin in commercialapplications is being pushed forwards by a number of drivers including high crudeoil prices, booming bioeconomies around the world, and rapidly increasing inter-est in sustainability across all industries. These drivers are being fuelled by rapidlyadvancing production/treatment processes, increased collaboration between indus-tries and the advent of modern biorefineries. Despite these seemingly favourableconditions, widespread adoption of lignin as a material for commercial products re-quires new business models, further technological advancement and better supplychain networks.41As a potential source of carbon for battery electrodes, lignin is an attractiveprecursor material for a global battery industry that is demanding cheap and sus-tainable component materials. If lignin can be incorporated into high performancenext generation batteries, the energy storage industry may be the catalyst requiredto help the lignin industry realize its long-awaited potential. Interestingly, due tothe high aromatics content in the macromolecules of lignin, the biopolymer is wellsuited as a precursor for the synthesis of carbon materials that are being consideredfor next generation battery applications. Recently, lignin has been used success-fully in energy storage applications in a number of studies. For example, ligninderivatives were used in polypyrrole composites to create renewable cathodes withhigh specific capacitance values [88]. Even more recently, groups have taken ad-vantage of the ability of lignin to form into nanofibres to explore the use of ligninderived carbon nanofibres for energy storage applications.Depending on the type of lignin, a lignin molecule is composed of various com-binations and ratios of three basic phenyl propane monomeric sub-units: CoumarylAlcohol (H), Coniferyl Alcohol (G) and Syringyl Alcohol (S). Typically, hard-wood lignin has a S/G ratio of 2:1 while softwood lignins have a ratio of around1:2 or 1:3 [89]. The production of lignin-based carbon nanofibres from lignin canbe divided into three basic steps: i) spinning of lignin solution; ii) thermal sta-bilization of as-spun lignin nanofibres at temperatures around 250◦C in air andiii) carbonization of stabilized fibres at temperatures upwards of 800◦C in an inertenvironment (Fig 3.9). After spinning is complete, a thermal stabilization pro-cess is essential to the production of carbon nanofibres because it prevents meltingof fibres during the carbonization step. During thermal stabilization, moleculesbecome crosslinked, weak intermolecular bonds become broken, ketones and car-42bonic acids form and autoxidation of aldehydes occurs [90]. These processes in-crease the glass transition temperature of the lignin, which allows the fibres toremain intact during carbonization[89]. During the carbonization at high temper-atures, functional groups are eliminated from the lignin molecules leaving behinda conductive carbon backbone and graphitic structures. Additionally, during car-bonization, macroscopic fusions between fibres are created, resulting in a highlyinterconnected structure that is ideal for energy storage electrodes because it elim-inates the need for binders or current collectors [89].Figure 3.9: 3 Step process for lignin-based carbon nanofibre preparation.Lignin can easily be formed into carbon nanofibres, a strategy that has beenpursued for applications in supercapacitors and in both anodes and cathodes forbatteries. Successful integration of lignin into energy storage technologies willrepresent a significant step in the search for sustainable storage devices. A varietyof different methods can be used to prepare carbon nanofibres from lignin pre-cursors. For example, Chatterjee et al. used a melt spinning processing to preparecarbon fibres with graphitic nanoscale domains [91]. In recent years and because ofits scalability, electrospinning has been proposed as a method for producing ligninbased carbon nanofibres. Non porous CNFs have been successfully fabricated us-ing this method [92, 93]. By adding porosity to these fibres to increase the surfacearea, electrospun lignin-based carbon nanofibres may constitute an ideal electrodematerial for Li-S batteries and energy storage devices in general. Porous lignin43nanofibres have been prepared successfully ([94, 95]), but despite these works,there have been few reported studies of porous lignin based ECNF electrodes forenergy storage applications and fewer examples of porous lignin based ECNF elec-trodes with tuneable fibre porosity and high surface area.44Chapter 4Lignin-based Porous Carbon Nanofi-bre ElectrodesWith the vast amounts of waste lignin being produced by the pulp and paper indus-try, and because of the suitability for this material for carbon nanofibre preparation,there exists a significant opportunity for development of high performance energystorage device electrodes from this material. In reviewing the literature, it is clearthat optimal energy storage electrodes should be highly conductive, should havegood mechanical resilience and flexibility, ideally should be free-standing and lightweight, should have high porosity and surface area, should be low-cost and easy toproduce from highly scalable fabrication techniques and should be sustainable. Inthis work, the electrospinning process was selected as a highly scalable and simplemethod of fabricating highly conductive and interconnected carbon nanofibres tobe used as a conductive matrix for free-standing electrodes. Because high surfacearea and porosity has been demonstrated as a characteristic of high performanceelectrodes in both supercapacitor and battery applications, ECNFs require furthermodification.In order to develop porosity in ECNFs, a template method was selected inwhich lignin is co-electrospun with small amounts of polytheylene oxide (PEO).45The PEO is then selectively removed from the fibres prior to carbonization viaa facile hydrothermal treatment step. This step will be shown to be effective atremoving PEO from the surface of the fibres to create high levels of porosity andto increase the specific surface area of the fibres significantly. The resultant matrixis highly conductive and consists of a web of highly porous ECNFs with fusedjunctions.4.1 Materials & Methods4.1.1 MaterialsPowder of softwood kraft lignin was obtained from FP-Innovations and was driedin a vacuum oven for 24hrs at 50◦C before use. PEO powder with a molecularweight of 900,000 was obtained from Sigma-Aldrich and was used as-is with noadditional treatment. Spectroanalyzed N,N-Dimethylformamide (DMF) was ob-tained from Fisher Scientific.4.1.2 ElectrospinningSpin dopes with mass ratios of lignin to PEO of 99/1, 98/2 and 97/3 were preparedby first dissolving PEO in DMF and stirring at 80◦C for 1 hour. PEO served asa plasticizer to aid in electrospinning. It was found that spinning solutions withhigher or lower proportions of PEO resulted in fibres containing high amounts ofbeads or in a lack of fibre formation all together. Lignin was simultaneously andseparately dissolved in DMF and stirred by magnetic stirrer at 80◦C for 30 minuteswith intermittent periods of heavy agitation using a vortex mixer, followed by ul-trasonication with a Sonics Vibracell tip sonicator for 30 minutes with a repeatedpattern of 5 seconds of sonication followed by 5 seconds of rest to ensure full46dispersion. Next, the lignin solution was combined with the PEO solution in theappropriate amounts and the resultant solution was stirred for an additional 1 hourat 80◦C, again with intermittent periods of heavy agitation with a vortex mixer. Toensure spinnability of the solutions, the ratios of lignin to DMF for each of thethree solutions were as follows: 35% for ECNF(99/1), 32% for ECNF(98/2) and29% for ECNF(97/3). It was found that attempting to electrospin solutions withhigher amounts of PEO lead to poor to no fibre formation during electrospinning,and that attempting to spin without the addition of PEO lead to no fibre formationdue to too little fibre viscosity.The electrospinning setup consisted of a KatoTech Co. Ltd. NEU Nanofi-bre Electrospinning Unit with an electrically grounded and aluminum foil coveredstainless steel rotating drum of diameter 10 cm as the fibre collector. Solutions wereloaded into a 10 ml syringe with a Luer-Lok tip and a 20 guage blunt-tip stainless-steel needle and a built-in syringe pump was employed to dispense the solutionand maintain a distance of 17 cms between the syringe and the drum. Electro-spinning was completed using a positive voltage bias of 17 kV and a flow rate of4.3 ml hr−1. The solution was stirred continuously during electrospinning with ahome-made magnetic stirring apparatus. Atmospheric humidity was kept between28% and 32%. Electrospinning in a non controlled, high humidity atmosphere leadto fibres containing moisture beads.4.1.3 Removal of Sacrificial PolymerPorous Lignin/PEO fibres were prepared by submerging as-spun fibre mats indeionized water and gently agitating for at least 12 hours at a temperature of 80◦Cin order to dissolve the sacrificial PEO polymer from the fibres. It was found that47a soaking time of 12 hours was necessary to remove the all of the PEO from thesurface of the fibres and that soaking for longer than this did not contribute to ad-ditional mass loss from the fibres in a significant way. The fibre mats were thenremoved from the water and dried on a stainless steel rack in a vacuum oven at50◦C for 12 hours.4.1.4 Stabilization and CarbonizationFigure 4.1: Graphical representation of the stabilization and carbonizationprocess.In order to convert the electrospun fibres into conductive carbon nanofibres,fibre mats were placed on stainless steel racks in a Thermo Scientific F21135 fur-nace, with a controllable gas input and a mineral oil bubbler to regulate positive48pressure. Stabilization was done by heating the fibres in air to the stabilizationtemperature of 230-280◦C from room temperature at a rate of 1-2◦C min−1 andholding for 1-2 hours. Next, the fibres were carbonized by heating the fibres fromthe stabilization temperature to a carbonization temperature of 900◦C at a rate of5◦C min−1 and holding for 5 hours followed by natural cooling to room tempera-ture. A constant flow of nitrogen was maintained during carbonization. A graphicalrepresentation of the stabilization and carbonization process is shown in Fig 4.1.The fabrication process for graphitic nanofibre mat electrodes is summarized inFig 4.2.Figure 4.2: Fabrication process for porous lignin-based nanofibre mat elec-trodes: (1) preparation of spin dope from PEO, lignin and DMF, (2)electrospinning of lignin nanofibres, (3) appearance of as-spun ligninnanofibres, (4) hydrothermal treatment step, (5) appearance of car-bonized and porous nanofibres and (6) test cell schematic for electro-chemical testing.494.2 Characterization and evaluationFibre morphology was evaluated with a Zeiss Sigma Scanning Electron Micro-scope. Fibre diameters were determined using the SmartSEM software and bysampling at least 10 different fibres for each sample. Conductivity measurementswere taken via a two probe method using a Tektronix Precision Multimeter withfast drying silver paint being used to ensure good electrical contact between thefibres and the probe and to secure the sample to a glass slide. For this measure-ment, the fibres were cut into 3 mm wide and 20 mm long strips and the resistanceof the fibres was measured for at least 10 different samples. X-Ray diffractionpatterns were acquired with a Bruker D8-Advance X-Ray diffractometer in theBragg-Brentano configuration operated at 40 kV and 40 mA with Copper Kα1 &Kα2 radiation and a Nickel filter to filter the CuKβ radiation. For these XRD mea-surement, ECNFs were milled in a mortar and pestle to create a fine powder whichwas then dropped onto a glass slide. Raman Spectroscopy was performed with aHoriba Jobin Yvon HR-800 spectrometer with a laser wavelength of 633 nm. Poresizes were measured using Quantachrome Autosorb 1-MP gas sorption analyser.TGA curves were acquired with a Q500 thermogravimetric analyzer from TA In-struments. In a typical run, the temperature was increased from room temperatureto 800◦C in an air atmosphere at a rate of 10◦C min−1. The sample was purged byflowing air at a rate of 60 ml min−1.504.3 Results and Analysis4.3.1 Carbonization Parameter OptimizationTo demonstrate the viability of electrospun lignin/PEO nanofibres and to search foroptimal carbonization conditions, we prepared a solution of DMF containing 30%softwood lignin and 0.3% PEO by weight relative to the weight of solvent. Thissolution was successfully electrospun with fibres exhibiting a smooth and uniformmorphology but it was noted beads appeared on the electrospun fibres if the at-mospheric humidity reached above 35%. The fibres, which contain approximately99% lignin and 1% PEO were then stabilized in air for two hours at 230◦C after be-ing heated from room temperature at a rate of 1◦C min−1. Next, a high temperaturecarbonization step was employed. It is in this step that all non-carbon elements areeliminated and graphitic-like structures develop in the fibres which can influencethe final conductivity of the CNFs [96]. For this reason, optimizing the conditionsunder which these structures form is important. In order to determine the optimalcarbonization temperature to produce fibres with high conductivities, fibres werecarbonized in nitrogen for 5 hours at various temperatures (550, 700, 800, 900 and1000◦C). The carbonization heating rate was set to 5◦C min−1 in all cases. Theelectrical conductivity of the fibre mats was calculated from the following equa-tion:σ =Lt wR(4.1)where L is the length of the mat in cm, t is its thickness, w is its width and Ris the electrical resistance in Ω. Here, R was obtained using the aforementioned51two probe method. The prepared fibres exhibit a conductivity of 2.16 S cm−1 for acarbonization temperature of 900◦C and a falling conductivity for lower and highercarbonization temperatures as shown in table 4.1. These results were calculatedfrom at least 5 samples of each type. Fibres carbonized at 900◦C are smooth andhave an average diameter of 443 nm (s.d. = 69 nm).Carbonization Temperature Conductivity (S cm−1)550◦C 1.9*10−7 (s.d. = 1.4*10−8)700◦C 0.26 (s.d. = 0.02)800◦C 1.57 (s.d. = 0.09)900◦C 2.16 (s.d. = 0.16)1000◦C 1.41 (s.d. = 0.13)Table 4.1: Conductivities of lignin/PEO ECNFs(99/1) after stabilization at230◦C and carbonization at various temperatures.It was found that further optimization could be achieved by varying the sta-bilization temperature. Many processes occur during this oxidative stabilizationstep including cross-linking, dehydrogenation, cyclization, aromatization and oxi-dation which lead to the formation of a thermally stable ladder-like structure. Byoptimizing the conditions under which these processes occur, nanofibres with im-proved performance characteristics can be created [97, 98]. Table 4.2 presents themeasured conductivities of fibres prepared with stabilization temperatures from230-280◦C and with the carbonization temperature held constant at 900◦C. Fibresstabilized at a temperature of 250◦C had the highest conductivity of 2.71 S cm−1.The remainder of fibres in this work were fabricated using the optimized stabiliza-tion temperature of 250◦C and carbonization temperature of 900◦C.After optimizing the stabilization and carbonization temperatures, the effect ofstabilization rate and time on the final fibre conductivities was explored. The pre-pared fibres were stabilized using four different combinations of commonly used52Stabilization Temperature Conductivity (S cm−1)230◦C 2.16 (s.d. = 0.16)250◦C 2.71 (s.d. = 0.13)280◦C 1.68 (s.d. = 0.09)Table 4.2: Conductivities of lignin/PEO ECNFs(99/1) after stabilization atvarious temperatures and carbonization at 900◦C.stabilization rates and stabilization times. The fibres were then carbonized as usualand conductivity of the fibres was measured. The results, presented in Table 4.3indicate that a stabilization rate of 2◦C min−1 and a stabilization time of 2 hoursleads to the highest fibre conductivity. The remainder of the fibre mats presented inthis work were prepared under the following optimized conditions: increase tem-perature to the stabilization temperature of 250◦C in air at a rate of 2◦C min−1 andhold for 2 hours and then increase temperature to the carbonization temperature of900◦C at a rate of 5◦C min−1 and hold for 5 hours in a N2 atmosphere.TimeRate 1◦C min−1 2◦C min−11 hour 3.83 (s.d. = 0.28) 4.54 (s.d. = 0.17)2 hour 2.71 (s.d. = 0.13) 7.31 (s.d. = 0.65)Table 4.3: Effect of stabilization time and rate on the conductivities of lign-in/PEO ECNFs(99/1). Conductivity values reported with units S cm− Microstructural ObservationsWith stabilization and carbonization conditions optimized, electrospun carbon nanofi-bres with lignin to PEO ratios of 99/1 (ECNF(99/1)), 98/1 (ECNF(98/2) and 97/3(ECNF(97/3) were spun and submerged in water to remove the PEO before be-ing subjected to the heat treatment regime described above. The resulting porouscarbonized fibres (P-ECNFs) were then characterized. For comparison, charac-53terization of non porous ECNFs(99/1) was also completed. Fibre microstructureswere observed through SEM images after carbonization and the fibre diameterswere measured from the images by sampling at least ten different fibres from eachimage. SEM images for all fabricated ECNFs and P-ECNFs show consistent fibreformation with few to no beads. Fig 4.3 shows the SEM pictures of non porousECNFs(99/1). Fibre formation is very regular and the fibres have a smooth appear-ance with little to no mesopores visible. Additionally, the fibres appear to havefused together at the junctions, a feature that does not appear to as great at extentin fibres that have had the PEO removed before carbonization.Figure 4.3: SEM images of ECNF(99/1) at magnifications of 10k and 40k.Figure 4.4: SEM images of P-ECNFs(99/1) at magnifications of 10k and 40k.Additional images can be found in Appendix C.1.54Fig 4.4 shows SEM images of porous P-ECNFs(99/1). These fibres are againwell formed and of consistent diameter but differ from their non-porous counterparts in that there are numerous breaks in the fibres. These may indicate the pres-ence of pockets of PEO in the fibres which, when dissolved, cause the remainingfibres to become disjointed. Alternatively, these breaks may be an indication ofthe increased brittleness of the fibres caused by the removal of the PEO prior tocarbonization. The images also reveal the presence of obvious pores in the fibres,although these pores are relatively few and widely dispersed along the fibres. Thepresence of pores in the washed fibres indicates that PEO has been successfullyremoved from the fibres prior to heat treatment.Figure 4.5: SEM images of P-ECNFs(98/2) at magnifications of 10k and 40k.Additional images can be found in Appendix C.2.P-ECNFs(98/2), seen in Fig 4.5, are significantly more porous than P-ECNFs(99/1)and appear to fuse to each other to a greater extent than do the P-ECNFs(99/1). Thisis also a feature of the non washed fibres and so may indicate that some PEO re-mains in the fibres prior to carbonization. The P-ECNFs(98/2) appear well formedbut give the appearance of a more fluid and less rigid overall structure, a featurethat again may be explained by the presence of left-over PEO in the fibres prior to55carbonization. The presence of pores in these fibres is much more obvious than inP-ECNFs(99/1) and in the non washed fibres, again confirming that the hydrother-mal treatment step was successful in removing bubbles of PEO that were presenton the surface of the fibres.Figure 4.6: SEM images of P-ECNFs(97/3) at magnifications of 10k and 40k.Additional images can be found in Appendix C.3.Finally, Fig 4.6 show P-ECNFs(97/3) which, compared to the other fibres, areheavily porous with a wide variety of pore sizes that are dispersed all over the fibresalthough not as regularly as the other porous fibres discussed above. The high levelof porosity in these fibres suggests once again that hydrothermal treatment wassuccessful in removing PEO agglomerations on the surface of the fibres. Fibres arebead free and well formed but with periodic breaks in the fibres. P-ECNFs(97/3)show a very fluid and “melted” appearance with crossing fibres fused together attheir junctions which as before, may indicate that some amount of PEO remainsin the fibres prior to carbonization and after the hydrothermal treatment step. Thefused nature of the fibres may suggest high conductivity of the P-ECNFs(97/3)mats.Fibre diameters were measured using the SmartSEM software and are pre-56Fibre Type Fibre Diameter (nm)ECNF(99/1) 439.60(s.d.= 53.71)P-ECNF(99/1) 464.40(s.d.= 52.22)P-ECNF(98/2) 468.88(s.d.= 53.11)P-ECNF(97/3) 428.12(s.d.= 55.87)Table 4.4: Average diameters of different ECNFs and P-ECNFs.sented in Table 4.4. Differences in average fibre diametres between the four dif-ferent types of fibres may be expected, despite the fibres being electrospun underthe same conditions. For one, varying the ratio of PEO to lignin in the spin dopemay effect both the surface tension and the viscosity of the solutions, two factorswhich, as discussed earlier, may affect electrospun fibre diameter. Additionally,increased amounts of PEO in the as-spun fibres may lead to smaller fibres after thehydrothermal treatment step is completed. Despite these two factors, fibre diame-ters remained relatively consistent across the different fibre types, falling between428.12 nm and 468.88 nm. The differences in diameters for the different fibres isnot significant however and does not reveal a clear trend. Additionally, the diam-eter of the porous fibres did not seem to deviate significantly from that of the nonporous fibres with the same precursor ingredient ratios (439.60 nm and 464.40 nmfor ECNFs(99/1) and P-ECNFs(99/1) respectively).4.3.3 Surface Area and PorosityNitrogen adsorption porosimetry was used to measure the specific surface area,pore volume and pore size distribution of each of the fabricated fibre mats. Innitrogen porosimetry, nitrogen is introduced to the sample under cryogenic tem-peratures in tiny doses to saturation pressure. Then, the excess gas is removed andthe relative pressure and quantity of adsorbed gas is measured to give an adsorption57Figure 4.7: N2 adsorption isotherms of the prepared P-ECNFs and ECNFs.isotherm. The adsorption isotherms can be plotted as a straight line in the relativepressure range of 0.05 - 0.35, giving a BET plot. Specific surface area of the mate-rial can be calculated based on the slope (A) and the y-intercept (I) of this plot byway of the following equations:vm =1A+ I(4.2)SBET =vm N sV a(4.3)where vm represents the monolayer adsorbed gas quantity, N is Avogadro’s number,58V is the molar volume of the adsorbate, a is the mass of the sample and s is thecross sectional area of the adsorbing species [99].Fibre Type SSA (m2 g−1) Pore Volume (cm3 g−1)ECNF(99/1) 816 0.42P-ECNF(99/1) 974 0.49P-ECNF(98/2) 1076 0.52P-ECNF(97/3) 1258 0.65Table 4.5: BET specific surface area and pore volume of different ECNFs andP-ECNFs acquired from N2 adsorption isotherms.The isotherms, presented in Fig 4.7 for each of the fibres, show an increasinglevel of adsorption in the low pressure region, followed by continued adsorptionin the middle region and further continued adsorption in the high pressure region.This behaviour is characteristic of type 1 adsorption isotherms under the Brunauerclassification system which is characteristic of materials with high amounts of mi-cropores. The total quantity of adsorbed gas at each relative pressure was highestfor P-ECNFs(97/3) and lowest for P-ECNFs(99/1) with the general trend indicat-ing that those fibres with lower proportions of precursor PEO adsorbed less gas ateach relative pressure. The BET method was used to evaluate the specific surfacearea of each of the fibres and the results are presented in Table 4.5. The results re-veal that P-ECNFs(97/3) have a high specific surface area of 1258 m2 g−1, a valuethat represents a 5-fold improvement over non porous PAN based CNFs [53] anda 54% increase compared to the non porous ECNFs(99/1), which have a specificsurface area of 816 m2 g−1.The density functional theory (DFT) model was used to measure pore sizedistribution and the model used was the cylindrical pore non local DFT equilibriummodel. The results, presented in Fig 4.8, reveal that all fibres contain a large volume59Figure 4.8: Pore size distribution histograms of the prepared P-ECNFs andECNFs.of micropores (<20A˚) as predicted by the isotherm shape, with P-ECNFs(97/3)having the highest volume of micropores and with fibres with less precursor PEOhaving less micropore volume. Relative to micropores, all fibres have a smallervolume of mesopores (20-500A˚) but are still significantly mesoporous. Again P-ECNFs(97/3) have a much higher volume of mesopores compared to the otherfibres, particularly in the pore size range of 40 nm to 80 nm. The majority of themesoporosity induced in the P-ECNFs(98/2) and P-ECNFs(99/1) appears to be inthe 20-40 nm range.The total pore volume of P-ECNFs(97/3) is a high 0.65 cm3 g−1, again a sig-nificant improvement of 55% over the non porous ECNFs(99/1), which have pore60Figure 4.9: Specific surface area of ECNFs and P-ECNFs as a function oftheir precursor PEO content.volume of 0.42 cm3 g−1. The results indicate that removal of PEO from electro-spun fibres prior to carbonization contributes to the formation of both microporesand mesopores in the fibres and that the volume of both micro and mesopores canbe tuned by varying the amount of precursor PEO in the electrospinning solution.Specific surface area and total pore volume of the prepared fibres are visualized inFig 4.9 and Fig 4.10 as a function of PEO content in the electrospun fibres.4.3.4 X-Ray Diffraction AnalysisXRD analysis was performed on all ECNFs in order to further elucidate the mi-crostructure of the porous fibres. In XRD, a concentrated beam of x-rays is directed61Figure 4.10: Total pore volume of ECNFs and P-ECNFs as a function of theirprecursor PEO content.towards the sample using a cathode ray tube and a beam collimator. When inci-dent x-rays have enough energy to dislodge electrons from the targets inner shells,a characteristic x-ray spectra is produced. When conditions satisfy Bragg’s Law,that is when the following equation is satisfied:nλ = 2d sinθ (4.4)where d is the interplanar distance of the crystal lattice, n is a positive integer, θis the scattering of the x-rays and λ is the wavelength of the incident wave, thereflected x-rays will interfere constructively, producing a peak in x-ray intensity62when the diffracted x-rays are detected and counted. By scanning the x-ray beamthrough a range of 2θ angles, all possible diffraction directions of the lattice shouldbe scanned resulting in a characteristic diffraction pattern. This pattern can then becompared to known patterns to better understand the composition of the material[100].Figure 4.11: XRD patterns for different ECNF and P-ECNFs.Fig. 4.11 presents the results of this analysis and reveals a broad diffractionpeak for all P-ECNFs centred around 22-23. This peak is attributed to the (002)crystallographic plane of graphitic structures. A second weaker peak was also ob-served corresponding to the (101) plane of graphite at around 44 and the presenceof which is a strong indication of graphite formation [101]. While these results do63point to the formation of graphitic structures within the fibres, the broad nature ofthe visible peaks indicate that the size of graphitic structures is small and the de-gree of crystallographic order is low [102]. Interestingly, the peaks for the porousfibres appear more pronounced than that of the non-porous fibres, indicating thatgraphitic order in the porous fibres may be higher than in the non porous fibres.4.3.5 Lignin Fibre ConductivityFibre Type Conductivity (S cm−1)ECNF(99/1) 7.31(s.d.= 0.65)P-ECNF(99/1) 19.69(s.d.= 1.02)P-ECNF(98/2) 18.91(s.d.= 0.62)P-ECNF(97/3) 18.39(s.d.= 0.92)Table 4.6: Conductivity of prepared ECNFs and P-ECNFs.Results of conductivity measurement show that P-ECNFs are highly conduc-tive compared to non porous ECNFs with P-ECNFs(99/1) having a conductivityof 19.69 S cm−1 (s.d. = 1.02 S cm−1) compared to ECNFs(99/1) which had a con-ductivity of 7.31 S cm−1 (s.d. = 0.65 S cm−1). P-ECNFs(98/2) had a conductiv-ity of 18.91 S cm−1 (s = 0.62 S cm−1) while P-ECNFs(97/3) had a conductivity of18.39 S cm−1 (s.d. = 0.92 S cm−1). The results are summarized in Table 4.6 andFig 4.12 and indicate that P-ECNFs are between three and four times more con-ductive than conventional PAN based ECNFs (4.90 S cm−1) [103].The differences in conductivities of the different fibres may be explained bythe structural and morphological properties of the different fibres. XRD results re-vealed that differences in the heights of XRD peaks between the fibres with porousfibres having a much higher peak corresponding to the (002) plane of graphiticstructures and a slightly more pronounced peak corresponding to the (101) plane.64Figure 4.12: Conductivity of ECNFs and P-ECNFs as a function of their pre-cursor PEO content.This difference is consistent with the difference in conductivities between thesegroups of fibres indicating that indeed, the presence of graphitic structures in thefibres contributes to the conductivity of the fibre mat.4.3.6 Raman SpectroscopyRaman spectroscopy is another powerful technique with which one can identify themolecular composition of a material. Raman spectroscopy relies on scattering ofmonochromatic light directed onto the sample with the use of a laser. The reflectedlight experiences an energy shift caused by its interaction with the molecular vi-brations present within the material. The Raman spectra is collected by a detector65and the resulting shift can be compared to the Raman spectra of known materialsso identify the molecular make-up of the target material. The Raman shift can becalculated on the basis of the following equation:∆w = (1λ0− 1λ1)×10−7 (4.5)where ∆w is the Raman shift in cm−1, λ0 is the excitation wavelength in nm andλ1 is the Raman spectrum wavelength in nm.Figure 4.13: Raman spectra for different ECNF and P-ECNF.Raman Spectra for all samples are presented in Fig 4.13 and reveal obviouspeaks centred around 1330 and 1590 cm−1. These peaks correspond to the D- andG-band peaks for carbons respectively. It is thought that the D-band is caused by66either defects in curved graphene sheets or is from the vibration of carbon atomswith dangling bonds in crystal lattice structures which may be due to plane termi-nations of disordered graphene [104]. The formation of the G-band on the otherhand is from the vibration of sp2-bonded carbon atoms in graphene sheets andother two-dimensional hexagonal structures [105]. By measuring the ratio of theintensity of the D-band (ID) to that of the G-band (IG), an approximate measure ofthe graphitic quality of fibres can be obtained with a lower ratio indicating fewerstructural defects in the graphitized nanofibres. The ID/IG values were 1.48 forECNFs(99/1) and 1.26, 1.29 and 1.35 for P-ECNFs(99/1), P-ECNFs(98/2) and P-ECNFs(97/3) respectively, indicating a slight decrease in graphitic quality of thefibres as the amount of PEO increased in the precursor electrospinning solution.The results indicate that the formation of ordered graphitic structures duringheat treatment may be disrupted by the addition and then removal of PEO from thefibres prior to the heat treatment step. As seen in Fig 4.12 and Fig 4.14, there is astrong inverse correlation between the ID/IG values of the fibres and their measuredconductivities; fibres with lower ID/IG values (and therefore higher graphitic order)tended to have higher conductivities and vice versa. This correlation suggests onceagain that the graphitic order and quality within the fibres plays a significant rolein determining the conductivity of the fibre mats. Additionally, the elevated ID/IGvalue for the non porous ECNFs(99/1) suggests that the presence of high amountsof PEO in the fibres during the heat treatment step may disrupt graphite formationand conductivity. Again, this result is consistent with the lower conductivity valuesobserved for the non-porous ECNFs(99/1) and with the results of thermogravimet-ric analysis presented below.67Figure 4.14: ID/IG values of ECNFs and P-ECNFs as a function of their pre-cursor PEO content.4.3.7 Thermogravimetric AnalysisThermogravimetric analysis (TGA) is a powerful method that can be used to fur-ther understand and confirm the presence of different polymers in non carbonizednanofibres. In TGA, samples are heated to a specific temperature and at a constantheating rate. During heating, the weight of the sample is measured accurately re-sulting in a graph of percentage weight remaining as a function of temperature. Inthe case of PEO/Lignin fibres, TGA can be used to understand the composition ofthe fibres by taking advantage of the differences in decomposition profiles of thetwo polymers.The significant surface area and porosity of the P-ECNFs suggests that the68Figure 4.15: TGA profiles of P-ECNFs and ECNFs. The inset shows a mag-nification of the heating profile between 700◦C and 800◦C.hydrothermal processing step presented in this work is effective at removing thesacrificial PEO from the lignin fibres to create micro and mesoporosity. TGA re-sults can be further analysed to determine the true extent at which PEO removalwas completed. Thermogravimetric analysis was conducted for all fibres prior toheat treatment to assess the components of the prepared P-ECNFs and ECNFsand to confirm the removal of the sacrificial PEO during the water treatment step.TGA curves in Fig 4.15 with solid lines show data collected from the prepared nonporous ECNFs while dotted lines indicate measurements taken using the P-ECNFs.Additionally, TGA measurements were completed on both untreated lignin andPEO powders for comparison.69The results show that lignin powder begins to decompose at around 200◦Cand continues to decompose slowly until well over 700◦C, eventually leaving over47% weight as residue. This long thermal decomposition process may be dueto the presence of numerous molecular weight fractions in the non fractionatedlignin powder. On the other hand, the PEO powder, which is composed of a singlemolecular weight of PEO, begins to decompose at around 400◦C and finishes de-composing almost completely at around 430◦C. The fabricated non porous fibresfollow a similar decomposition profile to that of the lignin powder. Despite theabsence of a sharp drop in weight around 400-430◦C as would be expected in fi-bres with a much higher proportion of PEO, even the small amounts of PEO causethe TGA profile to shift downwards compared with the lignin powder profile. Thiscan be attributed to the additional loss of weight from the decomposition of thesmall amounts of PEO in the fibres with the profiles of those fibres containing themost PEO experiencing the most downward shift relative to that of lignin powder.As expected, the profiles for the porous fibres (P-ECNFs) are shifted up relativeto the non porous fibres since these fibres have had most of their PEO removedduring the hydrothermal treatment step and so should not expect as much materialdegradation during heating.Interestingly however, the profiles of the three porous fibres are offset fromeach other slightly with P-ECNFs(99/1) and P-ECNFs(97/3) having the most andleast final residue proportions respectively. This may suggest that not all of thePEO was removed from the fibres during the hydrothermal treatment step and thatsome PEO may remain in the fibres and particularly on the inside of the fibreswhich may not have been exposed to water. If all of the PEO contained in the fibrewas removed, we should expect the TGA profile of that fibre to match the profile of70lignin powder. PEO remaining in the fibres may affect the morphology of the fibresand indeed, it appears that the fibres containing the most PEO before carbonizationappear less straight and rigid in the SEM images. Instead, these fibres appear morefluid with rounded turns and more fusing of fibres.71Chapter 5Supercapacitors From Porous Lignin-based ECNF ElectrodesAs a demonstration of performance in electrical devices, supercapacitor deviceswill be used to test electrochemical performance of the prepared P-ECNF elec-trodes. Based on the results of electrochemical testing as well as the previouslyreported structural and morphological results, the prepared electrodes will be pre-sented as a sustainable, low-cost and high performance platform framework onwhich the development of future high performance batteries and capacitors can bebased.5.1 Materials & Methods5.1.1 MaterialsP250-1 ACS Certified Potassium Hydroxide Pellets (KOH) with a molecular weightof 56.11 were obtained from Fisher Scientific. Celgard 2500 microporous mem-brane was purchased from Celgard Company for use as a separator.725.1.2 Electrode PreparationCircular electrodes were stamped from as-carbonized sheets of ECNFs using agasket punch set purchased from General Tools Mfg. Co. and were 6mm in di-ameter. The mass of the circular ECNFs electrodes was measured with a SartoriusCPA225D semimicro balance. A Celgard 2500 microporous membrane was usedas a separator and was stamped into circular pieces of 8 mm diameters. Electrodesand separators where then soaked in electrolyte for at least 6 hours prior to electro-chemical testing. 6M KOH electrolyte was prepared by adding KOH Pellets to theappropriate amount of DI water and stirring for at least 1 hour.5.1.3 Test Cell FabricationFigure 5.1: Schematic representation of the electrochemical testing cell.Two-electrode cells were put together in a lab-designed testing cell (Fig 5.1)that included two stainless steel columns as current collectors surrounded in a PFAair-tight casing. Two electrodes were placed between the two current collectors73and separated by a Celgard separator and then the casing was filled with electrolyteusing a syringe. O-rings were employed in strategic locations to prevent escape ofelectrolyte and to keep the cell closed to the atmosphere. An integrated springsystem ensured that the testing cell applied the same pressure to the two electrodesacross subsequent tests. The testing cell was connected to the testing equipmentvia a small connection port on either side of the test cell.5.2 Characterization and EvaluationCyclic voltammetry (CV), galvanostatic charge/discharge at constant current (CC)and electrochemical impedance spectroscopy (EIS) of the two-electrode ECNF cellwas completed using a Multipotentiostat VMP 300 from Bio-Logic and with a 6MKOH aqueous solution as the electrolyte. Scanning rates used for CV tests were5, 10, 30 and 50 mV s−1 in the voltage window of 0 to 0.8 V. Constant currentgalvanostatic charge/discharge studies were performed at current densities of 400,1000 and 2000 mA g−1 over the same voltage window. EIS measurements wereperformed over the frequency range from 0.1 Hz to 1 MHz.5.3 Results and Analysis5.3.1 Cyclic VoltammetryCyclic Voltammetry (CV) is an electrochemical analysis technique whereby a po-tential difference is applied across the electrodes of a supercapacitor cell and rampedat a continuous rate between two limit potentials. Once the upper limit potentialhas been reached, the voltage is then ramped back down to the lower limit poten-tial to complete one cycle of the scan. By measuring the resultant current flow, a74CV trace can be produced which can yield information about the capacity of theelectrodes, the voltage window of the supercapacitor and the cycle stability of thesystem. In an ideal capacitor, an applied voltage sweep creates a current accordingto the following equation:I =dQdt=CdVdt(5.1)where I is the current, Q is the charge and V is the voltage. Capacitance cantherefore be estimated from CV curves by integrating the current over the voltagewindow and dividing by the voltage window and the scan rate as in the followingequation:C =∫ E2E1 i(E)dE2(E2−E1)mv (5.2)where C is the specific capacitance of the sample in Farads per gram (F g−1), E1and E2 are the maximum and minimum voltages used in the scan in volts (V), i(E)is the current in amperes (A), m is the mass in grams (g) and v is the voltage scanrate in volts per second (V s−1). The result of this calculation multiplied by a factorof four is the capacitance value of an individual electrodes based on CV results fora two-electrode cell [106]. The CV trace of an ideal capacitor will appear as arectangle with a constant current being produced as the voltage is swept acrossthe voltage range. In practice and due to internal cell resistance and other effects,this is rarely the case but the degree of rectangularity of a CV curve can give anindication of the level of ideal supercapacitor behaviour being experience by thesupercapacitor being tested.Each of the porous fibres as well as the non porous ECNFs(99/1) were sub-75Figure 5.2: Cyclic voltammograms of different ECNFs at a scan rate of30 mV s−1.jected to CV testing. Scan rates used were 5, 10, 30 and 50 mV s−1. Of the fourfibres tested and as seen in Fig 5.2 and 5.3, P-ECNFs(97/3) exhibited the mostrectangular shaped curves with the largest areas at the high scan rates of 30 and50 mV s−1. As the fibre porosity decreased, so too did the size and rectangularshape of the CV curves. Additional CV curves for scan rates of 10 mV s−1 and5 mV s−1 can be found in Appendix A. The results indicate that P-ECNFs(97/3)would be the most suitable of the fibres tested for supercapacitor applications sincethey exhibit behaviour most closely approaching what would be expected from anideal capacitor. Additionally, the shape and size of the CV curves for the differentfibres indicates that P-ECNFs(97/3) may have the highest gravimetric capacitance.76Figure 5.3: Cyclic voltammograms of different ECNFs at a scan rate of50 mV s−1.Specific capacitance of the electrodes was calculated from the CV curves basedon Eqn 5.2 and are presented in Table 5.1. Results confirm that P-ECNFs(97/3)do indeed have the highest gravimetric capacitance of the fibres tested with amaximum measured capacitance of 112 F g−1. The demonstrated capacitance val-ues show a clear trend of increasing specific capacitance as the ratio of precur-sor PEO to lignin increases in the fibres. Capacitance values of ECNF(99/1), P-ECNF(99/1), P-ECNF(98/2), P-ECNF(97/3) obtained at the scan rate of 5 mV s−1where measured to be 78, 89, 105 and 112 F g−1 respectively. The increased spe-cific capacitance of the more porous fibres can be attributed to the increased spe-cific surface area of the P-ECNFs as the amount of precursor PEO in the fibres is77increased as well as the higher ionic accessibility of the more porous electrodes.Fibre Type 5 mV s−1 10 mV s−1 30 mV s−1 50 mV s−1ECNF(99/1) 78 70 55 43P-ECNF(99/1) 89 81 61 51P-ECNF(98/2) 102 94 71 54P-ECNF(97/3) 112 105 98 95Table 5.1: Gravimetric capacitance in F g−1 of ECNFs calculated from CVcurves.Next, volumetric capacitance was calculated for each of the prepared electrodesafter carefully measuring the overall volume of the electrodes in the cell and mul-tiplying the calculated gravimetric capacitance by the density of the electrode. Theresults, presented in Table 5.2 demonstrate a similar trend to the gravimetric capac-itance results with the more porous fibres having a higher volumetric capacitancethan the less porous fibres. P-ECNFs(97/3) had the highest gravimetric capacitanceof 35 F cm−3 at a scan rate of 5 mV s−1. Again, the results show that volumetriccapacitance decreases as the scan rate is increased.Fibre Type 5 mV s−1 10 mV s−1 30 mV s−1 50 mV s−1ECNF(99/1) 18 16 13 10P-ECNF(99/1) 23 21 16 13P-ECNF(98/2) 31 29 24 21P-ECNF(97/3) 35 33 31 30Table 5.2: Volumetric capacitance in F cm−3 of ECNFs calculated from CVcurves.With the goal of examining the effects of increased sweep rates, P-ECNFs(97/3)were further scanned at sweep rates up to 200 mV s−1 with the resulting curves(Fig 5.4) revealing good rectangular profiles even at high sweep rates. The vari-ation in capacitance at elevated sweep rates is an important factor to considerwhen designing supercapacitors for use at high rates. Capacitance values for P-78Figure 5.4: Cyclic voltammograms of P-ECNFs(97/3) at scan rates of (start-ing from the center and moving in the direction of the arrow) 5, 10, 30,50, 70, 100, 150, and 200 mV s−1.ECNFs and ECNFs(99/1) at all sweep rates are displayed in Fig 5.5 and show thatP-ECNFs(97/3) maintain high capacitance at elevated sweep rates. P-ECNFs(97/3)retained 84.4%, 78.8% and 71.1% capacitance at high rates of 50, 100 and 200 mV s−1respectively compared to the measured capacitance at 5 mV s−1, indicating excel-lent performance at high sweep rates.This impressive performance at high sweep rates can be attributed to fasterionic transportation and shorter diffusion distances which occur in the more porousnanofibre structure [107]. This was a significant improvement on the P-ECNFs(98/2),P-ECNFs(99/1) and ECNFs(99/1) which could only maintain 65.2%, 57.0% and79Figure 5.5: Specific capacitance of P-ECNFs calculated from CV curves atdifferent sweep rates.55.6% capacitance at a scan rate of 50 mV s−1 when compared to the measuredcapacitance at 5 mV s−1. At sweep rates higher than 50 mV s−1, the CV curvesof P-ECNFs(98/2), P-ECNFs(99/1) and ECNFs(99/1) displayed poor rectangularshape, indicating a loss of reversible supercapacitor behaviour at high sweep rates.The less porous structures in these cases may contribute to increased diffusion re-sistance which would explain the loss of supercapacitor behaviour at higher rates.5.3.2 Galvanostatic Charge/DischargeTo further analyze the electrochemical behaviour of the prepared electrodes, gal-vanostatic charge/discharge testing was conducted on all fibres at varying current80Figure 5.6: Galvanostatic charge/discharge curves of ECNFs and P-ECNFsat a current density of 400 mA g−1.densities. In galvanostatic charge/discharge testing, a constant current is appliedacross the electrodes of a supercapacitor and the resulting voltage is measured asa function of time. Once the voltage in the cell has reached a specific limit volt-age, the current is reversed and the voltage diminishes until it reaches a lowerlimit voltage, at which point the cycle is repeated. The resulting curves for anideal supercapacitor are perfectly triangular in shape with a lower discharge slopeindicating better capacitance. In real capacitors, galvanostatic charge/dischargecurves often exhibit a vertical drop in voltage at the moment the current directionis switched. This drop is referred to as an “IR drop” and is an indication of the levelof internal resistance in the cell. Galvanostatic charge/discharge was performed at81current densities of 400, 1000 and 2000 mA g−1 for all fibres with typical resultspresented in Fig 5.6 for a scan rate of 400 mA g−1. The results show good tri-angular shape with limited IR drop indicating promising reversible supercapacitorbehaviour for all P-ECNFs, with the most triangular and elongated shape belongingto P-ECNFs(97/3). P-ECNFs(97/3) were further scanned at current densities up to6000 mA g−1 with CC curves (Fig 5.7) continuing to maintain a good triangularshape with little IR drop even at high current densities. Additional galvanostaticcharge/discharge results are presented in Appendix B.Figure 5.7: Galvanostatic charge/discharge curves of P-ECNFs(97/3) at cur-rent densities of (starting from the center and moving in the direction ofthe arrow) 6000, 5000, 4000, 3000, 2000, 1000, and 400 mA g−1.825.3.3 Cycle Life TestingCycle No. ECNF(99/1) P-ECNF(99/1) P-ECNF(98/2) P-ECNF(97/3)1 100.0 100.0 100.0 100.0100 101.3 100.3 101.0 100.61000 101.7 101.6 101.9 101.52000 100.3 100.3 100.2 100.23000 100.0 100.0 99.8 100.04000 99.4 99.4 98.5 98.25000 98.9 98.3 95.5 94.96000 98.7 97.4 93.5 91.17000 98.1 96.4 90.5 87.58000 97.8 95.7 87.1 83.0Table 5.3: Capacitance of ECNF electrodes after cycling as a percentage ofthe capacitance measured on the first cycle.Retention of high capacitance values over a long cycle life is an importantproperty of supercapacitors. Galvanostatic discharge/charge from 0 to 0.8 V wasperformed on all fibres for 8000 cycles and, since capacitance is proportional tothe slope of the trace during discharge, the capacitance at various cycles was mea-sured relative to the capacitance calculated on the first cycle with results shownin Table 5.3. Fig 5.8 shows that all fibres maintain excellent cycling stability forthe first 3000 cycles followed by a slow decrease in capacitance afterwards. P-ECNFs(97/3) retained 100% capacitance after 3000 cycles and 98.3%, 91.1% and83.0% capacitance after 4000, 6000 and 8000 cycles respectively. The change inelectrochemical performance of the fibres is apparent in the galvanostatic charge/dis-charge profiles at different cycle numbers with curves at higher cycle numbers ap-pearing more horizontally compressed with a higher slope and showing a more pro-nounced IR drop (inset Fig 5.8). The less porous P-ECNFs(98/2) and P-ECNFs(99/1)maintained 87.1% and 95.7% capacitance after 8000 cycles respectively while thenon porous ECNFs(99/1) maintained 97.8%. The decrease in specific capacitance83may be due to the loss of small amounts of electrode material which is a knownoccurrence in carbonaceous materials and may be exacerbated by nano and mesoporosities which explains the larger capacitance decrease in the more porous fibres[102]. The small increase in specific capacitance on the 1000th cycle may be dueto the liquid electrolyte taking time to slowly penetrate into all of the pores of thefibres during the charge discharge process. This is otherwise known as the wettingeffect.Figure 5.8: Cyclic stability of P-ECNFs and ECNFs(99/1) after 8000 cyclesin 6M KOH. The inset shows the change in the galvanostatic charge/dis-charge profile of P-ECNFs(97/3) after 6000 cycles.845.3.4 Electrical Impedance SpectroscopyElectrical Impedance Spectroscopy (EIS) is a powerful analysis tool that be usedto elucidate the charge transfer and resistance properties of electrode materials. InEIS, a sinusoidal AC current is applied to the capacitor and the resultant voltage ismeasured. From this, the frequency dependant impedance can be calculated basedon the following equation:Zw =EwIw(5.3)where Ew and Iw are the frequency dependant voltage and current respectively. Ina supercapacitor, the measured voltage will be a function of both the resistive be-haviour of the system and the capacitative behaviour of the system and will followthe following relationship:Ew = I R− j Xc I (5.4)where R is the series resistance of the capacitor and Xc is the capacitative reactanceof the system. It should be clear then that the impedance of the system can be bro-ken into a real part (Z’) which dominates at low frequencies and reflects the seriesresistance behaviour of the system, and imaginary part (Z”), which dominates athigh frequencies and reflects the capacitative behaviour of the system. By measur-ing the potential as a function of current in a range of frequencies and by plottingthe imaginary impedance as a function of the real impedance, a Nyquist plot canbe prepared from which important information about the system can be inferred.The Nyquist plots for the prepared nanofibres are presented in Fig 5.9. Each85Figure 5.9: Nyquist plots from EIS analyses of ECNFs(99/1) and P-ECNFs.of the curves show a distinct semicircle in the high frequency range followed bya linear portion in the low frequency range. The intersection of the curve withthe x axis in the high frequency range gives an indication of the electrolyte resis-tance as well as the square of the resistance of the samples in the cell. All foursamples cross the x-axis close to the origin with P-ECNFs(97/3) appearing to havethe lowest sample and electrolyte resistance (having an x-intercept of 0.22) andwith there being a trend of decreasing sample/electrolyte resistance as porosityincreases. P-ECNFs(98/2), P-ECNFs(99/1) and ECNFs(99/1) had x-intercepts of0.89, 1.02 and 2.17 respectively. The semicircular region of the plot correspondsto the charge transfer resistance, with smaller semicircle diameters indicating a86higher migration rate of ions at the electrolyte/electrode interface. The data revealsa trend of increasing semicircle diameter as the porosity of the fibres decreaseswith P-ECNFs(97/3) having a significantly smaller diameter and therefore, a muchsmaller charge transfer resistance than the other fibres. ECNFs(99/1) appear to ex-hibit the highest charge transfer resistance with the EIS curve displaying a largesemi-circle diameter. In the lower frequency region, the finite slope of the samplesgives an indication of the diffuse resistivity of the electrolyte within the electrodepores [108]. The more vertical slope of the P-ECNFs(97/3) line, which approachesthe vertical line expected from an ideal capacitor, is an indication of the more de-sirable ion accessibility within the electrode compared to the other fibres.The above electrochemical tests suggest that porous lignin-based carbon nanofi-bre electrodes are well suited to applications in supercapacitors and energy storagedevices in general. P-ECNFs(97/3) in particular had high gravimetric capacitance,long cycle life and good cell resistance characteristics. These electrodes displayedstrong reversible supercapacitor behaviour, with CV traces appearing rectangularand with limited IR drop apparent on the galvanostatic charge/discharge traces.87Chapter 6Discussion6.1 Structural Suitability of Porous Nanofibre MatsThe suitability of materials for applications in electrodes for energy storage devicesis dependant on a number of factors. Among these factors, high performance elec-trodes for applications in devices like supercapacitors and lithium sulfur batteriesrequire high surface area, high micro and mesoporosity, good electrical conduc-tance and strong physical resilience. The prepared electrodes have demonstratedacceptable physical resilience and can be manipulated during supercapacitor fab-rication without showing signs of breaking or cracking. They can be used as free-standing electrodes without the addition of binding agents which are often a ne-cessity for many other carbon based electrode materials. This property is highlydesirable because it contributes to a high specific capacity and simplifies the fab-rication process. The free-standing nature of the prepared electrodes make themhighly suitable for applications in energy storage devices.In order to further analyze the suitability of lignin-based and porous carbonnanofibre mats for use in high performance energy storage devices, the conductiv-ities of the prepared fibres are considered. As mentioned previously, it has beendemonstrated that the performance of double layer capacitors is heavily influenced88by the conductivity of the materials with which the electrode is prepared [109].Additionally, electrodes with high conductivities perform better in battery appli-cations by reducing internal cell resistance. For these reasons, electrode materialswith high conductivities are very attractive. After spending time optimizing thecarbonization conditions, porous and non porous fibres were prepared that boastedimpressive conductivities, with P-ECNFs having conductivities in the range of 18-20 S cm−1.Fibre Type Conductivity (S cm−1)ECNFs(99/1) 7.31(s.d.= 0.65)P-ECNFs(99/1) 19.69(s.d.= 1.02)P-ECNFs(98/2) 18.91(s.d.= 0.62)P-ECNFs(97/3) 18.39(s.d.= 0.92)PAN[109] 4.9Porous Nafion:PAN 60:40[61] 0.13Porous Fe3O4/PAN[110] 0.0004762Lignin:PVA (85:15)[111] 16.0330%F4 Lignin[112] 530%F4F1−3 Lignin (900◦C)[112] 3030%F4F1−3 Lignin (1000◦C)[112] 5535%F4SKL/PEO Lignin[93] 2−5Table 6.1: Conductivity of different carbon nanofibres as reported in the lit-erature.The results of conductivity testing indicate that P-ECNFs are between threeand four times more conductive than conventional non-porous PAN based ECNFs(4.9 S cm−1) [103] and much more highly conductive than porous PAN based fi-bres. Table 6.1 presents the reported conductivities of a number of other compara-ble nanofibre-based electrodes as reported in the literature. It is interesting to notethat in general, the reported conductivity values of porous PAN based nanofibresare much lower than that of non porous PAN based fibres. This is not the casefor the fibres presented in this work. The high conductivity of the prepared fibres89relative to commonly used nanofibre materials indicates that lignin-based porousnanofibre mats are indeed a highly promising option for developing high perfor-mance electrodes for energy storage devices. It is more difficult to make a directcomparison to other lignin based carbon nanofibre electrodes presented in the liter-ature due to the variability in types, purity and origins of different lignins. The pre-pared P-ECNFs did appear to have comparable conductivities to non porous ligninbased nanofibrous electrodes and the conductivity of the fibre mats presented inthis work fell in the mid range of reported values for lignin based nanofibre mats.A review of the literature did not find any reported conductivity values for porousand lignin-based carbon nanofibres. As a platform for the further development ofenergy storage electrodes, the proposed system has significant room for further in-creases in conductivity. The simplicity of the electrospinning process makes theaddition of additional conductivity increasing nanoparticles and conductive agentspossible. Commercial energy storage devices based on the proposed platform maytake advantage of a wide range of additional doping agents.As well as having a high conductivity, a high surface area can positively influ-ence the capacitance of supercapacitor electrodes and can provide more conduc-tive carbon to interface with the active materials in battery applications, therebyincreasing active material loading. One of the goals of this research was to developa method for introducing increased surface area into electrospun carbon nanofi-bre electrodes. This was achieved by creating porous nanofibres with tuneableporosity via a hydrothermal PEO removal step. Using this process, prepared P-ECNFs(97/3) achieved a specific surface area of 1258 m2 g−1 and a total pore vol-ume of 0.65 cm3 g−1. A review of the literature reveals that the achieved specificsurface area represents a 4-fold improvement over non porous PAN based CNFs90[53]. P-ECNFs(97/3) perform well when compared to other porous non-ligninbased electrode materials despite falling short of the values of over 2000 m2 g−1 forsome PAN based electrode materials [60]. When compared to other Lignin-basednanofibrous electrodes, P-ECNFs have a high specific surface area (Table 6.2).Fibre Type Specific Surface Area (m2 g−1)ECNF(99/1) 816P-ECNF(99/1) 974P-ECNF(98/2) 1076P-ECNF(97/3) 1258PAN/PHMS[53] 302.50PAN/PMMA/TEOS[57] 699PAN/PVA[62] 1232PAN/Nafion[61] 1600PAN/PMMA[60] 2419Lignin[113] 456Lignin/PVA[102] 583Alcell Lignin[114] 1195Table 6.2: Specific surface area of different carbon nanofibres as reported inthe literature.Table 6.3 presents some specific surface areas typical of other common carbonmaterials for electrode applications that are not carbon nanofibres. P-ECNFs(97/3)have a specific surface area that falls in the midrange of these values, an impres-sive achievement for carbon nanofibre based electrodes which, traditionally havesuffered from relatively low specific surface areas. High surface area of electrodecomponents materials is important in as far as it contributes to the capacitance ofthe electrode. From this point of view, P-ECNFs(97/3) stand out from comparableelectrode materials due to their high capacitance. Additionally, a high specific sur-face area may be attractive because it provides more interface for the addition ofperformance enhancing additives and active materials, particularly in battery appli-cations. For this reason, the high specific surface area of P-ECNFs(97/3) make the91prepared electrodes suitable as a platform for enhanced future electrodes for bothsupercapacitors and batteries.Fibre Type Specific Surface Area (m2 g−1)Activated Carbon 1000−3500Functionalized Porous Carbons 300−2200Particulate Carbon from SiC/TiC 100−120Carbon Nanotubes 150−500Templated Porous Carbons 500−3000Activated Carbon fibres 1000−3000Carbon Cloth 2500Carbon Aerogels 400−1000Table 6.3: Specific surface area of different carbon materials as reported inthe literature. Table adapted from ref [115].Importantly, the prepared electrodes exhibit high porosity compared to nonporous carbon nanofibre electrodes. One of the major contributions of this workis the development of a process by which both micro and mesoporosity can beintroduced into electrospun fibres in a controllable fashion. Recently, high profileworks have indicated that mesoporous carbons may be the most promising materialfor high performance Li-S batteries [116, 117]. The prepared mesoporous ECNFshave a total pore volume that compares favourably to values reported in the liter-ature for other mesoporous carbon nanofibres (Table 6.4) and therefore, with theirhigh conductivity in mind, may be suitable for applications in Li-S batteries.The hydrothermal treatment process has been demonstrated as a simple methodof producing micro and mesoporosities in lignin-based carbon nanofibres. Analy-sis of results from SEM imaging, conductivity testing, XRD and Raman analysisand TGA reveal electrode mats that are conducive to being used in both superca-pacitor and battery applications. As discussed earlier, Kraft lignin is a stronglysuitable material for energy storage applications due to its low cost, sustainabil-92Fibre Type Pore Volume ( cm3 g−1)ECNF(99/1) 0.42P-ECNF(99/1) 0.49P-ECNF(98/2) 0.52P-ECNF(97/3) 0.65PVP/TEOS[118] 0.65−1.02PVA/Tin[119] 0.37−0.60PAN/TiO2[120] 0.095-0.225Table 6.4: Pore Volume of different mesoporous carbon nanofibres as re-ported in the literature.ity and natural abundance. Therefore, the ability to develop free standing, highlyporous, highly conductive fibres with competitive surface area from this materialfor applications in energy storage devices is significant and could contribute to thenext generation of high performance energy storage devices including supercapac-itors and Li-S batteries.6.2 Electrochemical Suitability of ECNF Mats forSupercapacitor ApplicationsWhile the structural and morphological properties of the prepared electrodes makethem attractive for applications in energy storage devices, the electrochemical prop-erties of this work should be discussed in the context of electrode material researchin general in order to determine to merit of the method and materials presented.The electrochemical performance of the electrodes was tested by constructing su-percapacitor devices and subjecting them to CV, galvanostatic charge/dischargeand EIS testing. CV testing revealed a maximum specific capacitance of 112 F g−1for P-ECNFs(97/3) at a voltage scan rate of 5 mV s−1. This is an improvement overother similar electrode materials such as activated lignin derived mesoporous car-bon (102.3 F g−1) [121] and non porous lignin and polyvinyl alcohol based nanofi-93bres (64 F g−1) [102]. Table 6.5 compares measured capacitance values with otherelectrode materials reported in the literature.Material Capacitance (F g−1)ECNF(99/1) 78P-ECNF(99/1) 89P-ECNF(98/2) 102P-ECNF(97/3) 112Activated Carbon < 100Functionalized Porous Carbons 100−150Particulate Carbon from SiC/TiC 100−120Carbon Nanotubes < 60Templated Porous Carbons 60−140Activated Carbon fibres 80−200Carbon Cloth 40−80Carbon Aerogels < 80Table 6.5: Specific capacitance of various carbon materials measured in aque-ous electrolytes as reported in the literature. Table adapted from ref[115].P-ECNFs(97/3) exhibit a good specific capacitance relative to other carbonbased electrode materials at a voltage scan rate of 5 mV s−1 meaning porous lignin-based ECNFs are highly suitable for use in supercapacitors. Equally as important,P-ECNFs(97/3) have promising specific capacitance at high scan rates (95 F g−1at a scan rate of 50 mV s−1) and even maintain 71.1% capacitance at rates of200 mV s−1. Even at this high scan rate, the measured capacitance exceeds thatof many different carbon materials and highlights the suitability of this material forsupercapacitors that demand high charge and discharge rates. On the other hand,the measured volumetric capacitance of the prepared fibres may not be as competi-tive. The best result of 35 F cm−3 again was achieved for P-ECNFs(97/3) at a scanrate of 5 mV s−1 which falls below the volumetric capacitance values reported forsome other carbon based electrode materials (Table 6.6). The large spaces in be-94tween the fibres contribute to a low density which is common in carbon nanofibreelectrodes and generally leads to a low volumetric capacitance in these types ofelectrode materials compared to more dense materials.Material Capacitance (F cm−3)ECNF(99/1) 35P-ECNF(99/1) 31P-ECNF(98/2) 23P-ECNF(97/3) 18Activated Carbon 19−611Mesoporous Carbons < 436Graphene 30−488Carbon Nanotubes 11−132Carbide Derived Carbon 75−180Table 6.6: Volumetric capacitance of various carbon materials measured inaqueous electrolytes as reported in the literature. Values taken from ref[122].Despite the relatively low volumetric capacitance of the fibres, he rectangularprofile of the CV trace for all scan rates indicates good reversible supercapacitorperformance, especially for P-ECNFs(97/3). This behaviour is substantiated bythe EIS results which suggest low sample and electrolyte resistance, low chargetransfer resistance and good ion accessibility within the electrodes. In additionto improving the performance of these electrodes in supercapacitor applications,these properties are highly desirable for battery applications as well.Finally, and in addition to high specific capacitance, the prepared P-ECNF elec-trodes exhibit very good cycle stability based on galvanostatic charge/dischargetesting. In particular, ECNFs(99/1) retained 97.8% capacitance after 8000 cycles.Even P-ECNFs(97/3), which displayed the worse cycling stability of the four pre-pared fibre types retained 100% capacitance after 3000 cycles, 91.1% after 6000cycles and 83% capacitance after 8000 cycles. This result is a significant improve-95ment on other lignin based carbon nanofibre electrode materials which maintained90% capacitance after 6000 cycles [102], and on non lignin based porous carbonnanofibre electrode materials which maintained 94% capacitance but after only2000 cycles [55]. The combination of high cycle life and excellent specific capaci-tance, coupled with the low charge transfer resistance and good ion accessibility inthe electrodes make the prepared fibres, and especially P-ECNFs(97/3) highly suit-able for use in supercapacitor applications and give a good indication of promisingpotential use in battery applications.6.3 Commercial ConsiderationsWith promising structural and electrochemical properties having been demonstrated,an analysis of the cost of the proposed system should be undertaken to help un-derstand it’s longer term commercial viability. As previously mentioned, today’scommercial supercapacitor devices usually use activated carbons derived from co-conut husks and can provide about 350 F for a mid-size device. Typical cost forthe electrode material for such a device was calculated in section 3.2 to be around$0.19 USD. With a capacitance of 112 F g−1, an equivalent mid-size commercialdevice would require about 1.56 g of ECNFs. For P-ECNFs(97/3), this translatesto 1.51 g of lignin and 0.05g of PEO. A good estimate of the cost of producinglignin based carbon fibers is around $6.3 USD/kg [123] and PEO costs approx-imately $1.5 USD/kg. This equates to a total cost of electrode material of 1.92cents for a typical device utilizing P-ECNFs(97/3) and a reduction in cost by afactor of about 10 compared to commercial supercapacitors. While the costs ofassembling devices have not been considered in this analysis, it appears that thehigh electrochemical performance, low material cost and absence of additives in96the proposed electrodes may contribute to a significant reduction in cost for super-capacitors utilizing this system compared to commercial capacitors.97Chapter 7Conclusions and Recommendations7.1 Porous Lignin-based Nanofibrous Electrodes forEnergy Storage DevicesIn this study, high performance electrodes were prepared from porous carbon nanofi-bres for applications in energy storage devices. Porous carbon nanofibres werefabricated using a highly scalable process involving electrospinning of a low-cost,naturally abundant and sustainable Kraft lignin precursor followed by a facile hy-drothermal treatment step and a heat treatment step including carbonization at hightemperature. Four different fibre types where prepared, each containing differentratios of lignin to PEO in their precursor solutions. These fibre types were labeledas ECNFs(99/1), P-ECNFs(99/1), P-ECNFs(98/2) and P-ECNFs(97/3). The pre-pared electrodes were used to assemble supacapacitor devices in a lab designedtesting cell with 6M KOH as the electrolyte. Supercapacitors utilizing electrodesprepared from precursor lignin solution containing 97% lignin and 3% PEO (P-ECNFs(97/3)) exhibited the highest gravimetric capacitance of 112 F g−1 at a scanrate of 5 mV s−1.Initially, stabilization and carbonization parameters were optimized in order tomaximize conductivity of carbon nanofibres. The optimized heat treatment condi-98tions were found to be 2 hours of stabilization at a temperature of 250◦C followedby a ramping phase at 2◦C min−1 to a carbonization temperature of 900◦C for 2hours. These parameters were used throughout the rest of this work. The mor-phology and structure of the prepared fibres was first observed from SEM imageswhich revealed well-formed fibres featuring varying degrees of porosity depend-ing on the ratio of lignin to PEO in the precursor solution which was used to makethem. Porosity appeared well distributed in the fibres. Fibre diameter was relativelyconsistent, falling in the range of 428-469 nm for all fibre types, however fibre mor-phology differed slightly between the fibre types, with those fibres containing moreprecursor PEO having a more fluid and less rigid appearance and comprising morefusion points between adjacent fibres. TGA was employed to verify the content ofthe prepared fibres and to confirm the effectiveness of the hydrothermal treatmentstep at removing PEO from the fibres prior to carbonization. It was found that,while the hydrothermal treatment was indeed effective at removing PEO locatedon or near the surface of the fibres, some PEO remained in the fibres, even afterthe washing step was completed and that the relative amount of PEO remainingcorrelated with the proportion of PEO in the precursor solution. The TGA resultsconfirmed the expectation that the washing step would only be able to remove thePEO that was located at the surface of the fibres.Because high surface area and porosity in electrodes may be crucial for theirsuccess in energy storage devices, surface area and porosity were examined usingnitrogen adsorption porosimetry. The BET method revealed that P-ECNFs(97/3)had a very high surface area of 1258 m2 g−1 which represents a 5-fold improve-ment over comparable but non-porous carbon nanofibres based on PAN [53] anda 54% increase in surface area compared to the non porous ECNFs(99/1), which99have a specific surface area of 816 m2 g−1. This dramatic increase in surface areademonstrates the effectiveness of the proposed hydrothermal treatment step duringfabrication at increasing the surface area of the fibres and makes the surface area ofthe prepared fibres competitive with other high surface area carbon materials com-monly used in electrode fabrication. Additionally, the density functional theory(DFT) model was used to measure pore size distribution in all fibres and revealedhigh mesoporosity in the porous fibres compared with the non porous fibres, espe-cially in P-ECNFs(97/3) which contained high pore volume, particularly for poresin the range of 40-80 nm.Conductivity of the prepared fibres was measured using a two point probemethod with results indicating high conductivity for the porous fibres. P-ECNFs(99/1)had the highest conductivity (19.69 S cm−1), almost a 3-fold improvement on thenon-porous ECNFs(99/1). This high conductivity means that porous carbon nanofi-bres prepared from lignin can achieve a conductivity that is four times higher thanfor conventional PAN based nanofibres. Both XRD and Raman spectroscopy whereconducted on the fibres in order to elucidate the molecular and crystal structuresof the fibres. XRD analysis revealed strong peaks around 22 and 44, which werefound to correspond to graphitic structures. As expected, the height of these peakscorrelated well with the conductivity of the fibres, indicating that graphitic struc-tures within the fibres are contributing to the high conductivity of the fibres. Sim-ilarly, Raman spectroscopy results revealed that conductivity of fibres may be af-fected by the degree of graphitic order within the fibres and that PEO containedwithin the fibres during carbonization may disrupt the formation of well orderedgraphitic structures.Results of the structural and morphological characterization tests suggest promis-100ing suitability of porous lignin-based nanofibres for use in energy storage applica-tions. Due to the free standing and binder free nature of the electrodes, their highconductivity and their high surface area and porosity, the prepared electrodes maybe a promising material for use in next generation energy storage and in particularmay be suitable for use in Li-S batteries and supercapacitors. Additionally, theproposed electrodes utilize low-cost, sustainable and naturally abundant precursorlignin as a carbon source and the fabrication process is both simple and highlyscalable. Table 7.1 compares some of the performance parameters of the proposedelectrodes with typical PAN based non-porous electrodes. In order to gage the per-formance of the proposed electrodes in energy storage applications, supercapacitordevices were developed in the lab and electrochemical testing was conducted onall fibre electrodes.Parameter PAN LigninConductivity (S cm−1) 4.9 [103] 18.39Specific Surface Area (m2 g−1) 236.75 [53] 1258Relative Cost 2 [90] 1Sustainability Low HighCapacitance (F g−1) 165 [124] 112Table 7.1: Performance of porous and lignin-based nanofibrous electrodescompared to non porous PAN based electrodes.Cyclic voltammetry testing was completed and CV traces were used to deter-mine the specific capacitance of each of the prepared electrodes with P-ECNFs(97/3)having the highest gravimetric capacitance of 112 F g−1 at a scan rate of 5 mV s−1.This value represents an improvement over many other commonly used carbonmaterials reported in the literature. Even at high scan rates of 50 mV s−1, P-ECNFs(97/3) had a high capacitance of 95 F g−1 indicating that these electrodescould be suitable for energy storage devices with high rate capability and requiring101fast charge and discharge. Additionally, these fibres had very rectangular CV tracescompared to other fibre electrodes which suggests strong reversible supercapaci-tor behaviour. The strong electrochemical performance of the fibres at high rateswas attributed to better ionic accessibility and shorter diffusion distances withinthe electrodes. On the other hand, the fibres exhibited a relatively low volumet-ric capacitance which is common for carbon nanofiber based electrode materials.P-ECNFs(97/3) had the highest volumetric capacitance of 35 F cm−3.Next, cycle stability was measured by way of galvanostatic charge/dischargetesting. The results show that even after 3000 cycles, the prepared fibre elec-trodes maintained 100% of their initial capacity. After 6000 and 8000 cycles,P-ECNFs(97/3) maintained 91.1% and 83.0% capacitance respectively while EC-NFs(99/1) maintained 98.7% and 97.8%. The higher cycle stability of the lessporous fibres was attributed to lower rates of material loss which is a common oc-currence in carbon materials and can be exacerbated by micro and mesoporosities.Lastly, EIS was performed on all fibres and indicated that P-ECNFs(97/3) had thelowest sample, electrolyte and charge transfer resistance of the four samples andthat increasing porosity correlated well with a decrease in these resistances. EISresults confirmed also the desirable ion accessibility characteristics of the moreporous fibres. Electrochemical testing of the prepared fibres reveal strong perfor-mance in supercapacitor applications and highlight the promising nature of lignin-based and porous carbon nanofibre electrodes as a platform for the developmentof other energy storage devices like Li-S batteries. Additionally, with an estimatedcost about 10 times lower than that of today’s commercial supercapacitors, theproposed electrodes could contribute to advancing energy storage technology inthe renewable energy, the transportation and the consumer electronics industries.1027.2 Recommendations for Future WorkAdditional optimization the proposed electrode materials and preparation methodswill be required before porous and lignin-based carbon nanofibre electrodes can bescaled up for commercial purposes. Based on the results presented in this work,which revealed a strong relationship between porosity and surface area of the fibresand supercapacitor performance, the author suggests that further increasing surfacearea may allow the proposed electrode design to achieve even higher performanceresults. Increasing the PEO content of precursor electrospinning solutions mayhelp to achieve this aim but difficulties relating to electrospinnability of high PEOcontent solutions will have to be addressed. Also, while this work has demon-strated how creating porosity in fibres can contribute to increasing surface area inthe fibres, it remains to be seen whether porosity or fibre size is the dominatingfactor in determining surface area. An analysis of the importance of each of thesefactors would be useful.In addition to increasing porosity and as a platform for further electrode de-velopment, the author suggests that conductivity of the electrodes can be furtherincreased through the targeted addition of conductive agents and nanoparticles dur-ing the electrospinning process. Additionally, further testing should be completedto understand the behaviour of the proposed electrodes in other electrolyte systemsincluding organic electrolytes. The behaviour of the system under these conditionswould give additional information as to their potential applications in Li-S batteriesand their battery systems.One of the disadvantages of utilizing carbon nanofibers for electrodes is that theresulting materials end up being significantly less dense than some other common103carbon materials. This can result in a lower volumetric capacitance and indeed,the prepared fibres had relatively low volumetric capacitances. By improving thedensity of the system, improvements in volumetric capacitance could be made.Future efforts could be directed at “filling the gaps” between fibres with highlyconductive and porous structures although care must be taken to maintain good ionaccessibility in the fibres.While increased electrochemical performance is always desired, the proposedelectrodes perform well in this regard as-is. One of the flaws in the proposed ma-terial is the lack of flexibility in these electrodes which, could be caused by the in-troduced pores and defects in the surface of the fibres. An analysis should be doneon the effect of these defects on the over all flexibility of the system to determinewhether these defects contribute in any significant way to the low flexibility of thesystem. While flexibility may not be necessary for many applications, it is indeedhighly sought after in applications like wearable and flexible electronics. The ad-dition on additional nanomaterials may provide an avenue for achieving flexibility.For example, Iron nanoparticles have been shown to increase flexibility in similarcarbon nanofibres and may, as an added benefit, increase the conductivity of thefibres and contribute to better electrochemical performance [112]. Increased flex-ibility of the fibres may also contribute to the physical resiliency of the proposedelectrodes and may further contribute to the already promising cycling stability ofthese electrodes.While the results in this work confirmed the suitability of the proposed elec-trodes for applications in supercapacitor electrodes, one of the motivations behindthe development of these electrodes was the problems currently facing Lithium-Sulfur battery technology. Despite, the electrodes in this work being designed with104LiS batteries in mind, their suitability for this type of energy storage device remainsto be confirmed. In particular, future work should look at the effect of porosity onthe cycle life of Li-S cells to verify the effectiveness of increased porosity in car-bon nanofibres at counteracting the negative effects of the shuttling effect that iscommonly seen in Li-S batteries. Also, one weakness of carbon nanofibre basedelectrodes is the low density of the material. 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PhD thesis, Georgia Institute of Technology, May 2007.115Appendix AAdditional Cyclic VoltammogramsA.1 Cyclic Voltammograms of ECNFs(99/1) at VariousScan Rates.Figure A.1: CV trace of ECNFs(99/1) at a scan rate of 5 mV s−1.116Figure A.2: CV trace of ECNFs(99/1) at a scan rate of 10 mV s−1.Figure A.3: CV trace of ECNFs(99/1) at a scan rate of 30 mV s−1.117Figure A.4: CV trace of ECNFs(99/1) at a scan rate of 50 mV s−1.A.2 Cyclic Voltammograms of P-ECNFs(99/1) at VariousScan Rates.Figure A.5: CV trace of P-ECNFs(99/1) at a scan rate of 5 mV s−1.118Figure A.6: CV trace of P-ECNFs(99/1) at a scan rate of 10 mV s−1.Figure A.7: CV trace of P-ECNFs(99/1) at a scan rate of 30 mV s−1.119Figure A.8: CV trace of P-ECNFs(99/1) at a scan rate of 50 mV s−1.A.3 Cyclic Voltammograms of P-ECNFs(98/2) at VariousScan Rates.Figure A.9: CV trace of P-ECNFs(98/2) at a scan rate of 5 mV s−1.120Figure A.10: CV trace of P-ECNFs(98/2) at a scan rate of 10 mV s−1.Figure A.11: CV trace of P-ECNFs(98/2) at a scan rate of 30 mV s−1.121Figure A.12: CV trace of P-ECNFs(98/2) at a scan rate of 50 mV s−1.A.4 Cyclic Voltammograms of P-ECNFs(97/3) at VariousScan Rates.Figure A.13: CV trace of P-ECNFs(97/3) at a scan rate of 5 mV s−1.122Figure A.14: CV trace of P-ECNFs(97/3) at a scan rate of 10 mV s−1.Figure A.15: CV trace of P-ECNFs(97/3) at a scan rate of 30 mV s−1.123Figure A.16: CV trace of P-ECNFs(97/3) at a scan rate of 50 mV s−1.Figure A.17: CV trace of P-ECNFs(97/3) at a scan rate of 70 mV s−1.124Figure A.18: CV trace of P-ECNFs(97/3) at a scan rate of 100 mV s−1.Figure A.19: CV trace of P-ECNFs(97/3) at a scan rate of 150 mV s−1.125Figure A.20: CV trace of P-ECNFs(97/3) at a scan rate of 200 mV s−1.126Appendix BAdditional Galvanostatic Charge/Dis-charge TracesB.1 Galvanostatic Charge/Discharge Traces forECNFs(99/1) at Various Current Densities.Figure B.1: Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 400 mA g−1.127Figure B.2: Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 1000 mA g−1.Figure B.3: Galvanostatic charge/discharge trace for ECNFs(99/1) at a cur-rent density of 2000 mA g−1.128B.2 Galvanostatic Charge/Discharge Traces forP-ECNFs(99/1) at Various Current Densities.Figure B.4: Galvanostatic charge/discharge trace for P-ECNFs(99/1) at a cur-rent density of 400 mA g−1.Figure B.5: Galvanostatic charge/discharge trace for P-ECNFs(99/1) at a cur-rent density of 1000 mA g−1.129Figure B.6: Galvanostatic charge/discharge trace for P-ECNFs(99/1) at a cur-rent density of 2000 mA g−1.B.3 Galvanostatic Charge/Discharge Traces forP-ECNFs(98/2) at Various Current Densities.Figure B.7: Galvanostatic charge/discharge trace for P-ECNFs(98/2) at a cur-rent density of 400 mA g−1.130Figure B.8: Galvanostatic charge/discharge trace for P-ECNFs(98/2) at a cur-rent density of 1000 mA g−1.Figure B.9: Galvanostatic charge/discharge trace for P-ECNFs(98/2) at a cur-rent density of 2000 mA g−1.131Appendix CAdditional SEM Images132C.1 SEM Images of P-ECNFs(99/1).Figure C.1: SEM images of P-ECNFs(99/1) at various magnifications.133C.2 SEM Images of P-ECNFs(98/2).Figure C.2: SEM images of P-ECNFs(98/2) at various magnifications.134C.3 SEM Images of P-ECNFs(97/3).Figure C.3: SEM images of P-ECNFs(97/3) at various magnifications.135


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