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Doctoral Thesis
Stress-Relieving Multiscale Design of Silicon Anodes for Lithium-ion Battery
Jaegeon Ryu
Department of Energy Engineering (Battery Science and Technology)
Graduate School of UNIST
2018
Stress-Relieving Multiscale Design of Silicon Anodes for Lithium-ion Battery
Jaegeon Ryu
Department of Energy Engineering (Battery Science and Technology)
Graduate School of UNIST
Stress-Relieving Multiscale Design of Silicon Anodes for Lithium-ion Battery
A thesis/dissertation
submitted to the Graduate School of UNIST in partial fulfillment of the
requirements for the degree of Doctor of Philosophy
Jaegeon Ryu
06/08/2018 Approved by
_________________________
Advisor
Soojin Park
Stress-Relieving Multiscale Design of Silicon Anodes for Lithium-ion Battery
Jaegeon Ryu
This certifies that the thesis/dissertation of Jaegeon Ryu is approved.
06/08/2018
signature
___________________________
Advisor: Soojin Park
signature
___________________________
Prof. Jae Sung Lee
signature
___________________________
Prof. Byeong-Su Kim
signature
___________________________
Prof. Yoon Seok Jung
signature
___________________________
Prof. Jun Hee Lee
Abstract
Replacing current fossil fuel-based power units comes down to the importance of developing high efficiency energy storage system (EES) in various energy related applications, which closely correlate with environmental conservation and renewable energy sources. Accordingly, great technological advances have been made for rechargeable battery systems, while lithium-ion batteries (LIBs) take precedence in current battery market due to their high operation voltage, light-weight and transformable shapes to give a compatibility with forthcoming systems such as electric vehicles (EVs), smart electronics and even grid technology. However, conventional combinations of carbonaceous anode and lithium transition metal oxides cathode could not satisfy the above requirement such as high specific/volumetric energy density or fast-charging ability.
Silicon (Si) anodes could provide a significant energy increase with a low operating voltage, high specific/volumetric capacity and natural abundance as long as we mitigate a mechanical failure which originates from large expansion upon Li insertion. Stress dissipation in Si anodes poses a major challenge in commercialization of Si anodes, thus requiring a rational design to release the stress and achieve robust operation of battery.
Herein, we lay out various designs at a specimen level that greatly affects the electrochemical performances of Si-based anodes toward high energy density LIBs. In each proposed design, different synthetic protocols and evaluation methods in electrochemical and material aspects are covered. From enormous amounts of researches on Si anodes, fundamental principle for material construction builds on the shortened diffusion length of Li-ion and stress-releasing ability facilitated by nanostructuring or interfacial controls that are further characterized by advanced analysis techniques.
In chapter Ⅱ, we introduce a mechanically clamping-layer coated one-dimensional (1D) silicon nanotubes (SiNTs) anode. The preparation process is based on an electrospinning technique followed by metallothermic reduction process. Large volume change of Si anodes can be suppressed by two structural advantages of nanotubular structure and outermost rigid layers of alumina and metal silicide, which guarantee the stable operation of battery and further enhance the interfacial stability. Interestingly, the silicide layer is electrically conductive to improve the electron transfer, thus increasing the rate capability as well.
In chapter Ⅲ, we report an ultrathin nanosheet structure via molten salt induced exfoliation and simultaneous thermal reduction process using natural clays. This one-pot process enables a scalable production way for high purity two-dimensional (2D) and mesoporous Si anodes. The 2D Si have nano-thick dimension in vertical direction and micrometer-scale lateral dimension that increase the material density as well as volume-suppression ability. Also, surface mesoporosity helps to accommodate volume changes in Si anodes for electrochemical stability.
In chapter Ⅳ, three-dimensional (3D) hyperporous Si flakes are reported for high Coulombic efficiency anodes. 2D flake-like natural clays that bear metal oxide interlayers are uniformly bored out via selective reduction and etching process to produce multi-scale pores on its framework. Control of pore types greatly increases the initial Coulombic efficiency as it provides a fast activation process through the macropores, which reversibly accommodate volume change by releasing the internal stress. This unique multi-dimensional structure shows a high durability over long repeated charge/discharge cycles.
In chapter Ⅴ, we report a cost-effective strategy using commercial metal-Si alloy through selective etching and thermal wet oxidation method. The synthesized porous Si frames are uniformly passivated by inactive alumina layers with remained small portion of metallic aluminum cores which act as a structural support. The protecting layers promote formation of stable solid-electrolyte-interphase (SEI) and improves the cycle life of batteries.
Contents
Ⅰ. Introduction --- 1
1.1 Battery choice for ESS --- 1
1.2 Anode Material choice for LIB --- 6
1.3 Potentials and progress of silicon anode --- 10
1.4 Reference --- 14
Ⅱ. Nanotubular structured Si-based multicomponent anodes for high-performance lithium- ion batteries with controllable pore size via coaxial electro-spinning--- 19 2.1 Introduction --- 19
2.2 Experimental --- 22
2.3 Results and discussion --- 24
2.4 Conclusion --- 38
2.5 Reference --- 39
Ⅲ. Synthesis of Ultrathin Si Nanosheets from Natural Clays for Lithium-Ion Battery Anodes --- --- 45
3.1 Introduction --- 45
3.2 Experimental --- 47
3.3 Results and discussion --- 50
3.4 Conclusion --- 64
3.5 Reference --- 65
IV. Multiscale Hyperporous Silicon Flake Anodes for high Initial Coulombic Efficiency and Cycle Stability --- 69 4.1 Introduction --- 69
4.2 Experimental --- 71
4.3 Results and discussion --- 73
4.4 Conclusion --- 88
4.5 Reference --- 89
V. Cost-effective approach for structural evolution of Si-based multicomponent for Li-ion battery anodes --- 93
5.1 Introduction --- 93
5.2 Experimental --- 95
5.3 Results and discussion --- 97
5.4 Conclusion --- 109
5.5 Reference --- 110
List of figures
Figure 1-1. A schematic diagram of energy storage applications.
Figure 1-2. Comparing power sources: energy versus power densities (Ragone plot).
Figure 1-3. Schematic of the working mechanism of a lithium-ion battery.
Figure 1-4. Schematic representation and operating principles of Li rechargeable batteries.
Figure 1-5. Approximate range of average working potentials and specific capacity of all types of electrode materials and those of representative intercalation, conversion and alloying anodes.
Figure 1-6. Charting the number of citations of peer-reviewed articles per years on the use of Si anodes for lithium ion battery and inset figure is three strategies applied for Si anodes per years.
Figure 1-7. General strategy of Si anodes including electrolyte control, size control, doping, functionalization, morphology control and composite formation.
Figure 2-1. (a) Schematic illustration of setup for coaxial electrospinning. The end-spinneret has separate paths for oil (inner) and polymer/inorganic solution (outer). (b) Experimental process for fabricating titanium silicide-coated Si-based multicomponent nanotubes. (c) XRD pattern of nanotubular Si-based multicomponents consisting of Si, Al2O3, and TixSiy. SEM images showing (d) SiO2 precursor-containing PVP nanotubes, (e) pure SiO2 nanotubes, (f) TiO2-coated SiO2
nanotubes, and (g) nanotubular TixSiy-coated Si-based multicomponent.
Figure 2-2. TEM images of four types of nanotubes of (a) SiO2, (b) Si, (c) Si taken from rectangular box seen in (b), (d) SiO2@TiO2, (e) Si@TixSiy, (f) line-scan graph of Si@TixSiy, and (g) EDS mapping images of Si@TixSiy.
Figure 2-3. SEM images of SiO2 tubes fabricated from different feeding rates of heavy oil at a fixed feeding rate (0.9 mL h-1) of polymer solution. The oil feeding rates are (a) 5, (b) 3, (c) 1, and (d) 0.5 μL min-1, respectively. (e) Plots of pore size and shell thickness variation as a function of oil feeding rate.
Figure 2-4. (a) Galvanostatic charge/discharge profiles of different Si-based nanotubes at 0.05C.
(b) Cycling performances and coulombic efficiency of nanotube electrodes at 0.2 C (1st-5th cycles) and at 0.5 C (6th-285th cycles). (c) Rate capability of TixSiy-coated Si nanotube electrode (discharge/charge rates are the same). (d) Electrochemical impedance spectra of carbon-coated and TixSiy-coated Si nanotubes after 1st and 50th cycles.
Figure 2-5. Cross-sectional SEM images of electrodes fabricated from three types of Si nanotubes:
(a-b) bare Si, (c-d) carbon coated Si, and (e-f) TixSiy-coated Si before cycling (a, c, e) and after 100 cycles (b, d, f).
Figure 3-1. Synthesis of ultrathin Si nanosheets. Schematic illustration shows synthetic routes
for preparing Si materials using two different strategies. (a) SEM and (b) TEM images of bare clay minerals. (c) SEM and (d) TEM images of ultrathin Si nanosheets synthesized via simultaneous molten-salt induced-exfoliation and chemical reduction (method II). (e) SEM and (f) TEM images of pre-exfoliated clay minerals. (g) SEM and (h) TEM images of collapsed Si nanosheets synthesized via successive chemical reduction (method I).
Figure 3-2. Characterization of Si nanosheets. (a) TEM image and (b, c) the corresponding bright- field HR-TEM image of SiNSs (inset in Fig. 2c shows FFT diffraction pattern). (d) AFM image and cross-sectional height profile of SiNSs. (e) Nitrogen adsorption-desorption curve of SiNSs (inset: BJH pore size distribution). (f) XPS spectrum and (g) Raman spectrum of ultrathin SiNSs.
Figure 3-3. Characterization of intermediates of SiNSs synthesized by method II. (a) SEM image and corresponding EDS mapping data of intermediates. (b) XRD patterns of bare clays, intermediate, and SiNSs. (c) TEM and (d) corresponding HR-TEM images of as-synthesized SiNS.
Figure 3-4. Electrochemical performances of SiNS anodes. (a) Galvanostatic first cycle discharge/charge voltage profiles at a rate of 0.05 C of three electrodes (N-SiNS, A-SiNS, and C- SiNS). (b) dQ/dV plots of A-SiNS electrode. (c) Rate capabilities of C-SiNS electrode at various C rates. (d) Cycle retentions of N-SiNS, A-SiNS, and C-SiNS electrodes (1st-5th: 0.1 C rate and 6th-200th: 0.2 C rate). (e) Long-term cycling stability of C-SiNS electrodes at three different C- rates (0.2 C, 0.5 C, and 1.0 C).
Figure 3-5. Surface morphologies of SiNSs after cycling. (a) SEM and (b) TEM images of N- SiNS after 200 cycles at a rate of 0.2 C. (c) SEM and (d) TEM images of C-SiNS after 200 cycles at a rate of 0.2 C.
Figure 3-6. Prelithiation of SiNS electrodes. (a) Schematic illustration showing prelithiation process of SiNS electrode on Cu foil by mechanical force. (b) First cycle discharge/charge voltage profiles of prelithiated electrode (for 20 min and 60 min) at a rate of 0.05 C. (c) Cycle retentions of pristine SiNS and prelithiated (20 min and 60 min) SiNS electrodes at a rate of 0.2 C.
Figure 4-1. Schematic illustration of HPSF synthesized from layered talc clay minerals and morphology/phase characterization. (a) Schematic illustration of 3D hyperporous structure.
Bowl-like cavity with 100nm-sized uniform macropores and meso-/micro-pores are formed on the framework. SEM images of HPSF taken at (b) low and (c) high magnification. (d) XRD pattern of HPSF showing phase transition from natural clay to pure silicon.
Figure 4-2. Structural analyses of HPSF. (a) TEM images showing entire structure of HPSF and (b) STEM dark-field image (inset is corresponding FFT diffraction pattern). (c) TEM image showing macro-pore regions of HPSF and (d-f) magnified TEM image showing edge region of HPSF. (g) Nitrogen adsorption-desorption isotherm curve, (h) BJH pore size distribution curve
and (i) Pore volumes distribution of HPSF.
Figure 4-3. Pore evolution mechanism characterization. (a-b) TEM images of 10 min-reacted HPSF showing no clear macropores and porous frame overall. (c-d) TEM images of 30 min- reacted HPSF showing evolution of macropores, but highly porous edges. (e-f) TEM images of 60 min-reacted HPSF showing well-defined thin frame structures with uniform distribution of macropores.
Figure 4-4. Spectral characterization. (a) Deconvoluted high resolution XPS spectra of HPSF.
Atomic silicon peaks are dominant and negligible shoulder peaks assigned to native oxide are detected. (b) Raman spectrum of HPSF showing characteristic crystalline silicon phase without amorphous silicon peaks and amorphous oxides.
Figure 4-5. Electrochemical properties of 3D HPSF and HPSF@C electrodes at initial stages. (a) Galvanostatic first cycle discharge/charge voltage profiles of HPSF and HPSF@C electrodes obtained at a rate of 0.05 C, comparing with carbon-coated bulk porous silicon. (b) Comparison of initial coulombic efficiency of different alloy-based anode materials.
Figure 4-6. Electrochemical properties of 3D HPSF and HPSF@C electrodes for prolonged cycles. (a) Cycle retention of p-Si@C (blue), HPSF (black) and HPSF@C (red) electrode at a rate of 0.2 C. (b) Long-term cycling stability of HPSF and HPSF@C electrode at a rate of 0.2 C (1- 100 cycles), 0.5 C (101-400 cycles) and 1.0 C (401-700 cycles). (c) Rate capabilities of HPSF and HPSF@C electrode at various C rates.
Figure 4-7. Post characterization of HPSF electrodes after cycling. TEM images of cycled (a-c) HPSF and (d-f) HPSF@C. Dashed circles and solid arrows (seen in Figure S16e) indicate macropores and framework, respectively.
Figure 5-1. Schematic illustration of synthetic approach for Si-based multicomponent from commercial Al-Si alloy through selective etching and wet-oxidation.
Figure 5-2. Structural characterization of a series of wet-oxidized etched Al-Si (ASWO). TEM images, STEM-HAADF image and EDX mapping data for (a-d) ASWO-04, (e-h) ASWO-10, (i- l) ASWO-20 and (m-p) ASWO-40 samples in regular order. Dotted lines (as visual aid) indicate amorphous Al2O3 layers on the surface.
Figure 5-3. Physical properties of wet-oxidized etched Al-Si (ASWO). (a) XRD pattern and (b) Raman spectrum of ASWO.
Figure 5-4. Electrochemical properties of ASWO in half cell. (a) Galvanostatic first cycle charge/discharge voltage profiles at a rate of C/20 in the potential window of 0.005-1.5V, (b) cyclic performances for 500 cycles at a rate of C/5, (c) rate capability of ASWO electrodes at various current density and (d) ex-situ impedance analysis of ASWO-10 electrode (pristine, first, second, and third cycled electrodes).
Figure 5-5. Electrode swelling results of ASWO series. SEM images of ASWO (a-c) 04, (d-f) 10, (g-i) 20 and (j-l) 40, respectively for pristine, after 1st lithiation and after 30 cycles.
Figure 5-6. Electrochemical properties of LiCoO2‖ASWO-10/NG full cell. (a) Cycling performance at 1C rate after 200 cycles (inset. first cycle charge-discharge profiles at C/10 rate) and (b) Rate capability of LiCoO2‖ASWO-10/NG full cell.
Nomenclature
0D Si 0-dimensional Si 1D Si 1-dimensional Si 2D Si 2-dimensional Si 3D Si 3-dimensional Si
BET Brunauer-Emmett-Teller analysis C2H2 acetylene gas
CE coulombic efficiency
CMC sodium carboxymethyl cellulose CVD chemical vapor deposition
DEC diethyl carbonate
DI water deionized water
EC ethylene carbonate
EDX(S) energy-dispersive X-ray spectroscopy EIS electrochemical impedance spectroscopy
ESS energy storage system
ETEM environmental transmission electron microscopy
EtOH ethanol
EVs electric vehicles
FEC fluoroethylene carbonate FFT fast Fourier transform FT-IR Fourier transform infrared
HCl hydrochloric acid
HR-TEM high resolution transmission electron microscopy ICE initial coulombic efficiency
LCO LiCoO2
LFP LiFePO4
LIB Lithium-ion battery
LiPF6 lithium hexafluorophosphate
LMO LiMn2O4
LTO Li4Ti5O12
PAA poly(acrylic acid)
PC polycarbonate
PVdF poly(vinylidene fluoride) SAED selected area electron diffraction
SBR styrene-butadiene rubber SEI solid-electrolyte-interface SEM scanning electron microscopy SiNP Si nanoparticle
SiNW Si nanowire
SiNT Si nanotube
SiNS Si nanosheet
STEM scanning/transmission electron microscope TEM transmission electron microscopy
TMS transition metal
XPS X-ray photoelectron spectroscopy
XRD X-ray diffraction
1
1. Introduction
1.1 Battery choice for ESS
Energy storage technologies are one of the most important matters in modern society, as the worldwide energy affairs such as blackout and power outages has brought inconveniences upon billions of consumers1-3. The generated various forms of energies in tremendous amounts from mechanical or natural sources should be efficiently managed with minimizing the loss4-5. Accordingly, most energy conversion and energy-generating systems are at a combined state with different energy storage systems (ESS)6. For a reliable and stable power supply to portable electronic devices, favorable electrochemical systems predominate in global energy market along with a great advancement of smart electronics and electric vehicles unlike grid-scale applications for mechanical or chemical systems7.
Electrochemical ESS realized by simultaneous ionic and electronic transport shows a roaring success among numerous candidates due to its deformable architectures, compactness, high storable energy and compatibility with existing technology, which consists of typically batteries and supercapacitors in rechargeable forms8-10. Lithium-ion batteries (LIBs), in particular, have far reaching effects on our daily life style from miniaturization of most electronic devices to electric load transportation, since the commercial launching in 1991 from SONY11-12. Such technical advancements require high energy and power density, and in this light, utilization of Li, the lightest and the smallest metal in periodic table might well be natural choice13. The single charged lithium ion (Li+) moves fast along the liquid media, delivering high cell potential as well as energies for the low standard reduction potential14. Furthermore, the LIBs almost overcome the memory effect observed in other batteries and have a low self-discharge rate15-16.
Fundamental understanding on LIBs principles begins with catching the four major components; two host electrodes (specifically, anode and cathode), reactive ion medium (electrolyte) and separator. At initial state of factory-released cells, Li+ migrates from Li- containing cathode through the electrolyte, being reduced at anode during charge and followed by reverse reaction during discharge. Continuing this repeated process, the charge storage in the form of Li+ takes place at both electroactive host materials. The amount of Li+ ions that the host materials initially have or could take into unit structure without network collapse determines the capacity of entire cells, in which the materials selection should be based on Li+ storage amount
2
per unit mass or volume and related ion kinetics of responding rate17.
The cathode materials are limiting electrode for overall lithium stoichiometry, given that the usable energies of stored lithium amount from lithium transition metal oxides or polyanions correspond to only maximum reversible energies18-19. Thus, the detailed operation of repeated charge/discharge cycles are configured by cathode materials for their specific applications20. As a commercialized form, LiCoO2 (LCO) provides a layered structure with superior stability and moderate level of specific capacities and potential ranges, which is suitable form mobile and portable devices. In parallel with recent development of EV applications, LiMn2O4 (LMO) and LiFePO4 (LFP) cathode have been extensively explored, while over-stoichiometric cathode (Li- rich, Ni-rich, Mn-rich) also have a great potential for high energy density, high power density, and improved safety, respectively21-24.
For anodes, carbon-based materials primarily including graphite and disordered hard carbons are first adopted for commercialized batteries25-26. The graphite anode theoretically delivers 372 mA h g-1 in a form of LiC6 within the graphene layers and offers good electric conductivity and chemical stability in about 10% expanding during charge27. However, low specific and volumetric capacity of graphite hinders forthcoming applications, which also limits the charging kinetics28. The material choices for anode materials naturally turn to different elements in periodic table that could reversibly alloy with lithium and deliver high capacity. Therefore, developing the anode materials with high specific/volumetric energy density need to take precedence toward future advanced LIBs29-30.
3
Figure 1-1. A schematic diagram of energy storage applications.
4
Figure 1-2. Comparison of power sources: Energy versus power densities (Ragone plot)31.
5
Figure 1-3. Schematic illustration of working mechanism of a lithium-ion battery32.
6 1.2 Anode material choice for LIBs
Anode host materials comprise three different reaction groups upon Li+ intake, where most common classification follows (de)intercalation, conversion, (de)alloying or those of combinations, along with chronological development flows33. Each compartment stands on the preferable electrochemical properties of cycling stability/fast charging ability, cost effectiveness, and high energy density of batteries34.
Archetypal (de)intercalation materials have characteristic direction of Li+ diffusion or inherently given spaces for Li+. For example, graphite, consisting of stacked graphene layers, allows only lateral diffusion of Li+ through the gaps between single graphene sheets and intercalated lithium ions are positioned with electrostatically stabilized by 6-membered carbon atoms, delivering the theoretical specific capacity of 372 mA h g-1 35. Since the Li+ is stored in the pre-reserved interatomic spaces and has the smallest size among the metal cations, structural changes in lateral dimension can be negligible during charging, while its interlayer spacing values increased less than 10%36-37. Although graphite anodes show chemical and mechanical stability with natural abundance and low reduction potential (< 0.1 V vs Li+/Li), sequential intercalation of Li+ through the layers causes sluggish diffusion kinetics, which implies that there is a high risk for direct plating of lithium metals on the electrode surface38. Besides, the use of polycarbonate (PC)-based electrolyte leads to simultaneous intercalation of Li+ and PC itself, causing the exfoliation of graphene sheets and severe battery degradation39. Still, graphite-based anodes are widely used in various electronic devices.
Less or hardly graphitized carbon materials, so-called hard carbon, could provide additional lithium adsorption sites in nanopores that originate from disordered grains40. Although there is a controversy on which reaction takes a precedence between adsorption and intercalations, physisorption to the surface defect sites and diffusion to disordered regions supplement the low Li+ kinetic of graphite anodes. So, high-power tools adapt refined structures of hard carbons41-42.
By extension, several inorganic intercalation materials for high stability and fast Li+ diffusion are suggested for replacing the carbonaceous anodes43-44. Most prominently, Li4Ti5O12 (LTO) anodes have highly reversible intercalation mechanism through the transformation from spinel- like structure to rock-salt LTO that is also known as zero-strain insertion materials with an average two-phase reaction potential of 1.55 V vs Li+/Li45. In principle, LTO anode accommodates 3 Li+ per unit structure which delivers only 175 mA h g-1 of gravimetric capacity, though, recent studies reveal that spinel-LTO can deliver higher capacities of 250-300 mA h g-1 at lower potentials46. However, poor electronic conductivity restricts fast charging applications of LTO anode, while surface modifications or nanostructured systems are demonstrated to realize the LTO anodes47-48.
7
Conversion anodes work based on a displacement reaction of transition metal (TM; Mn, Ni, Fe, Co, Cu, etc.) compounds with lithium ions in a form of oxides, sulfides, nitrides, fluorides, phosphides, and so on33. In other words, anionic atoms fully react with Li+, and then reduced TM nanoparticles are dispersed in its matrix where this complex has a multi-electron transfer systems and consequently high specific capacity of ~1,000 mA h g-1 at high reduction potentials of ~1 V vs Li+/Li. Taking TM oxides as an example49, displacement reaction produces amorphous Li2O media which further reacts with active species for much higher capacity than theoretical calculations. Similar to the case of LTO, typical conversion anodes have poor electronic conductivity, while a large volume expansion during cycles, irregular voltage plateaus and low initial reversibility pose additional challenges to access the high energy LIBs49.
Alloying-type anode could be an extension of subsequent reaction of the above displacement system29-30. Given that the reaction types depend on the coordination of electroactive materials with lithium, alloying host materials differs in that they completely lose their initial structure, and then form a new chemical structure with covalently bonded to lithium50. Therefore, Si, Ge, Sn, Sb, and Al take relatively large amounts of Li+ compared to other materials that belong to intercalation and conversion mechanism51. However, the large uptake naturally entails huge volume change of initial structure. Most of alloying anodes have a low reduction potential less than 0.5 V (vs Li+/Li) and both high specific/volumetric capacities (e.g. 3572 mA h g-1/8334 mA h cc-1 for Si and 1384 mA h g-1/7366 mA h cc-1 for Ge),52 which shed light on practical utilization of those choices. Yet, 300% expansion at a specimen level gives rise to serial degradation mechanism; first, large expansion generates surface crack formation or eventual pulverization53. Second, electrolytes are infiltrated into the slaps in active particles and lead to continuous decomposition of electrolytes54. Finally, inherently ion-conducting solid electrolyte interphase (SEI) layers thicken to block the further diffusion of Li+ and the batteries fail55. This degrading regime has been mitigated by many proposed strategies along with deep-understanding from mechanistic studies. Therethrough, Si-based anodes are at the present time work, considered as most promising choices for future LIBs.
8
Figure 1-4. Schematic representation and operating principles of Li rechargeable batteries33.
9
Figure 1-5. Approximate range of average working potentials and specific capacity of all types of electrode materials and those of representative intercalation, conversion and alloying anodes.
10 1.3 Potentials and progress of silicon anode
Most anode materials have a trade-off between high specific capacity and structural stability, in which the energy factors cannot be ignored to realize the downsized new electronic devices.
Thus, current battery society focused on first seeking out the several high capacity anodes and then reforming their properties and electrochemical behaviors. Among multiple choices, Si has drawn great attentions due to its low working potential, high specific capacity in a form of Li15Si4, natural abundance and mature Si technologies56. However, all these advantages can be guaranteed only if the generated internal stress or poor ionic/electronic kinetics are further addressed. Well- demonstrated strategies for past decades include reducing its size in nanoscale57, interfacial modifications58, morphology controls59, intrinsic property control60, and compositing with current-carrying elements61.
Micro-Si particles preclude the possibility of deep Li+ diffusion into the core, thus leaving behind the crystalline Si at the core-level and limiting the full usage of its capacity. In addition, huge internal stress creates the cracks on the surface during charging and it quickly propagates to pulverize the anodes or lead to delamination from current collectors62. However, rescaling the Si anodes to a nano-size enables the deep lithiation and staves off the structural failure. Interestingly, particle size of Si less than 150 nm neither fractured nor pulverized at initial lithiation stage, which is visually corroborated by in-situ TEM analysis57. Its unbreakable size due to insufficient stored strain energy to drive crack formation was obtained different from previous single-phase modeling whereby it considered only compressive hoop stress at the surface. Even with this valuable observation in the mechanical aspect, intimate problems of Si anodes cannot be solved.
Nano-sized Si anodes also practically suffer from unavoidable deformation in repeated cycles, which suggests that additional interfacial mediation toward a cracked surface or stable formation of interface layers is rather crucial as a second line of anti-pulverization62. Introducing mechanical/conductive layers, small portion of additive in electrolytes and artificial SEI layers on Si anodes are three representative methods to promote its stability. First, various external organic and inorganic components (carbon, metals, metal derivatives) protect the Si surface which act as a barrier between Si surface and electrolyte63. The protection layers could either restrict the expansion itself to reduce the usable capacity or elastically swell together with Si. Otherwise, conductive coating layers provide additional electronic conductivity to Si anodes and certain silicide films facilitate a thermal stability due to the robust SEI layer formation64. More recently, Choi’s group demonstrated molecularly engineered borate additive to protects the Si composite anode as well as cathode by tuning the fluoroethylene carbonate (FEC)-based SEI.
Morphological control of Si anodes represents a major effective method in nanostructured Si.
11
Various dimensional Si materials of 0-dimensional (0D) Si nanoparticle (SiNP), 1D Si nanotube or nanowire (SiNT or SiNW), 2D Si nanosheet (SiNS) and 3D porous Si (p-Si) have been widely investigated in their electrochemical properties59. Several important figure-of-merits in the commercial batteries can be satisfied with each dimensional Si anodes, such as initial Coulombic efficiency (ICE), cycle life, rate capability, swelling-suppression ability and tap density65. One- to-one comparison of their performances are well-condensed in previous review articles65. All these nanostructured Si materials are poor tap density in common and show improved structural integrity over the cycles. Ideally, multi-dimensional Si materials like further assembly of Si nano- building blocks into micrometer-sized secondary particles are highly desirable in successful commercialization of Si anode in practical market66.
Thus, Si-based composite anodes have been appealing to the battery society as well as battery R&D companies more recently. Such composites consist of conventional carbon support and nanostructured Si materials in either single framework or physical blending systems67. The conductive buffers render the composites much suitable to fast kinetics and prevent the Si fracture at a material level. In this regard, the question on how we could satisfy the above battery grade requirement in a single design of materials might be the last question of development stage, while intrinsic low electron conductivity of Si materials still can be addressed by introducing various dopants. However, even this strategy significantly changes the Li+ diffusion mechanism through the doped crystal structure at either reactive or disrupting state.
12
Figure 1-6. Charting the number of citations of peer-reviewed articles per years on the use of Si anodes for lithium ion battery and inset figure is three strategies applied for Si anodes per years68.
13
Figure 1-7. General strategy of Si anodes including electrolyte control, size control, doping, functionalization, morphology control and composite formation69.
14 1.4 Reference
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Chapter Ⅱ. Nanotubular structured Si-based multicomponent anodes for high-performance lithium-ion batteries with controllable pore size via coaxial electro-spinning
2.1 Introduction
Silicon (Si) has attracted much attention as one of the promising anode materials for next- generation lithium-ion batteries (LIBs) owing to its natural abundance and high theoretical specific capacity (3578 mA h g−1, Li3.75Si, at room temperature)1. However, its cycling performance is unsatisfactory for practical use in LIBs because its huge volume expansion (>300%) upon repeated cycles results in severe pulverization, loss of electric contact, and dramatic capacity fading2,3. Also, the poor intrinsic properties of Si such as the low electronic conductivity (∼10−3 S cm−1) and slow lithium diffusion (diffusion coefficient 10−14–10−13 cm2 s
−1) have impeded the utilization of its full capacity and its application to high-power and -energy- density tools4–10.
With the growing need for 1D materials, a great diversity of synthetic methods has been developed. The top-down approach from chemical etching of Si wafers27,28, chemical vapor deposition growth29–31, and electrospinning32–34 are representative examples of methods for synthesizing 1D Si-based materials. As repeatedly shown in the literature, polymeric and/or inorganic 1D nanofibers can be simply produced by means of the electrospinning process.
Recently, Yoo et al. demonstrated the synthesis of Si nanotubes by combining electrospun polyacrylonitrile nanofibers with silica coating followed by magnesiothermic reduction35. They used a well-known method for preparing Si by a magnesiothermic reduction of silicon dioxides.
During the reduction, tremendous amount of heat was generated, leading to structural disruption.
That is why as-prepared Si nanotubes showed crushed and shrunken tubes. Unlike the magnesiothermic reduction, an aluminothermic reduction using an aluminum reductant delivers less amount of heat. In view of the large-scale production, since batch size is directly proportionate to thermo-dynamically generated heat, the aluminothermic reduction is a suitable process to increase the batch size. As another approach, Park et al. reported Si nanotube lithium- ion battery anodes with greatly enhanced electrochemical performances21. However, their process, combining an anodic aluminum oxide template with a chemical reduction of the organic Si precursor, was a very expensive, low yield (<5%), and complicated multi-step process. Along
20
with above technique, there are various methods for preparing a tubular structure, in particular SiO2 nanotubes. Most of fabrications are based on template-assisted or surfactant-aided preparation methods. Using different templates or surfactants such as metal oxides, metals, metal salts, carbon nanofibers, and organic surfactants, high purity and macro-/nano-dimensional SiO2
nanotubes can be fabricated36-41. Furthermore, atomic layer deposition (ALD) can easily yield exceptionally conformal and uniform SiO2 nanotubes by extent of atomic-level. Among them, one of the effective methods to prepare the tubular structure is coaxial electrospinning using dual nozzles with following advantages; (i) it is specialized to synthesize core/shell and hollow tubular organic and/or inorganic fibers in one-pot process42-48, (ii) it has the controllability on the wall and core size of nanotubes by simply tuning the flow rates and also we can easily add other metal oxide precursors to make hybrid materials in the preparation of electrospinning solution, and (iii) electrospinning technique can produce long-dimensional fibers with product yield for further processing. For these reasons, coaxial electrospinning method is one of the effective tools to synthesize tubular metal oxide and/or hybrid metal oxides. To the best of our knowledge, direct fabrication of Si-based nanotubes by the coaxial electrospinning method has not yet been reported.
In addition to nanostructuring of Si, numerous coating materials, including carbon materials49-
54, metal55, metal oxide56, and metal silicide57, 58, have been introduced on the surface of Si-based materials. Most surface treatments have focused on either increasing mechanical property for Si anodes with a large volume change or enhancing electrical conductivity. Recently, Park et al.
demonstrated that a metal silicide coating can significantly enhance both mechanical and electrical properties of micrometer-sized Si particles57. The metal silicide-coated Si anodes showed remarkably improved cycling stability and high rate capabilities. Those enhancements were attributed to increased electrical conductivity, good mechanical property, and a stable solid- electrolyte-interface formation57, 58. Furthermore, Memarzadeh et al. reported ALD method on Si nanotube prepared by sacrificial template. They also showed the effect of various coating materials (e.g., TiO2, TiN, and Al2O3) on the electrochemical properties of Si nanotubes.59 By ALD, coating level can be precisely controllable, thus it has been implemented to thin-film LIBs and also has great advantages for coating irregularly nanostructured materials.60 Among strategies mentioned above, a hybridized approach, combining structural and interfacial effects, is one of the effective methods to make high-performance Si battery anodes. However, nanotubular Si- based materials with both structural integrity and interfacial effect have not been described in other reports.
Herein, we demonstrate a one-pot synthetic route for nanotubular Si-based multicomponents utilizing two different well-known processes, coaxial electrospinning and subsequent aluminothermic reduction reaction. Coaxial electrospinning enables us to make a nanotubular
21
structure of Si-containing materials, while the aluminothermic reaction can successfully convert silica to Si and form multicomponents including Si, alumina, and titanium silicide. The nanotubular Si-based multicomponent anodes show significantly improved electrochemical performances including increased initial coulombic efficiency (88.5%), cycling stability (~75%
after 280 cycles), high rate capability (a specific capacity of 650 mAh g-1 at 5 C rate), and extremely suppressed volume expansion (~14% after 100 cycles) compared to bare Si nanotubes.
22 2.2 Experimental
Coaxial electrospinning process
For the core solution, heavy mineral oil was used (Aldrich). For the sheath solution, poly (vinyl pyrrolidone) (PVP, average Mw ~ 1300K, Aldrich) and tetraethyl orthosilicate (TEOS, Aldrich) were mixed together in ethanol solvent for enough time to form a transparent solution. In a typical process, the solutions were prepared by mixing 2.36 g of PVP, 3.0 g TEOS, 0.5 g Pluronic P123 (average Mn ~ 5,800, Aldrich), and 10 g absolute Ethanol, and then heated up to 80 ℃, simultaneously adding 0.1 g of 0.1 M HCl. It was injected to the outer syringe pump (NE - 1000, New Era Pump Systems) channel, and heavy mineral oil was added to the core syringe pump channel, connected to a metallic needle. These two types of solutions were fed at a certain constant rate. The optimized feeding rate of polymer solution was set at 0.9 mL/h. In addition, the oil rate was varied from 0.5 to 5 μL min-1. Aluminum foil was placed 10 cm away from the needle to serve as a collector. All experiments were conducted at an accelerating voltage of 20 – 23 kV at room temperature in air. As-spun oil@polymer/SiO2 fibers were peeled off from the collector, followed by natural drying at room temperature overnight for complete hydrolysis of TEOS. After drying, the collected fibers were immersed in n-hexane solution (Aldrich) for 24 hours to extract oil. Furthermore, pure SiO2 nanotubes were obtained by calcination of polymer/SiO2 nanotubes at 250 oC (5℃min-1 rise) for 1 h (for stabilization) and 600 oC for 3 h (calcination of polymer and residual oil). In the TiO2 coating system, 2 g of as-prepared SiO2 nanotubes and 0.6 g hydroxypropylcelluose (average Mw ~ 100K, Aldrich) were well-dispersed in 100ml absolute ethanol for 30 min. Then, a certain amount of titanium butoxide (TB, Aldrich) was added to the above solution and it was further stirred for 2 h. With loading 1mL of DI water, TiO2 was well coated on the in/out surface of the SiO2 nanotubes. Finally, by heating these fibers up to 600 oC for 3 h, TiO2 coated SiO2 nanotubes can be obtained (Fig. S1 and Fig. S2b).
Aluminothermic reduction of SiO2-based nanotubes
1 g of TiO2 coated SiO2 nanotubes were mixed with 0.6 g of Al powder (3 – 5 μm) in ethanol so that the powders were uniformly mixed together. After completely drying out the solvent, the mixture was put on an alumina boat and placed in a tube furnace, heated up to 900 oC for 6 h under Ar atmosphere. Then, the product was leached out with
23
concentrated H3PO4 and 0.5 M HF for eliminating undesirable materials such as an excess amount of Al2O3 and SiO2. The fabricating condition for non-coated Silicon nanotubes is equivalent to the above experimental procedure except for the TiO2 coating process. The carbon coating process was conducted by thermal deposition of toluene vapor at 900 oC for 8 min.
Characterizations
Morphology of tubular Si nanotubes and different coating layers were characterized by scanning electron microscope (Verios 460, FEI) at an accelerating voltage of 10 kV, transmission electron microscopy (TEM, JEM-2100) at an acceleration voltage of 120kV, and high – resolution transmission electron microscopy (HR-TEM, JEM-2100F). TEM (JEM-2100) was also used to investigate tube structures before and after cycling. A carbon coating layer on Si nanotube was characterized by Raman spectrometry (Alpha 300s, WITec GmbH) operating with a laser excitation wavelength of 532 nm and elemental analysis. X-ray diffraction (D8 ADVANCE, Bruker) was used to investigate tubular typed silicon nanotubes with titanium silicide coating layers.
Electrochemical test
The battery performance was measured by galvanostatic cycling (WonATech WBCS 3000 battery measurement system) of coin cells with the fabricated titanium silicide coated silicon nanotubes along with carbon coated and non-coated silicon nanotubes as the working electrode and lithium foil as the reference / counter electrode. The working electrodes were prepared using a conventional slurry method with a series of silicon nanotubes, super P, 1:1 mixture of poly (acrylic acid) (PAA, weigh average molecular weight ~100K, Aldrich), and sodium carboxymethyl cellulose (CMC, 4 wt% in H2O, Aldrich) as a binder with a mass ratio of 7:1.5:1.5. The mass loading of active materials, except binder and conducting agent, was ~ 1.59 mg cm-2 and ~ 2.45 mAh cm-2. The potential windows for all cycled cells were between 0.01 and 1.2 V versus Li/Li+. The electrolyte comprised 1.3M LiPF6 in 3:7 v/v ethylene carbonate / diethyl carbonate with 10 wt% fluorinated ethylene carbonate additives to improve the cycling stability.
Polyethylene film (Celgard 2400) was used as a separator. These were all carefully assembled in an Ar-filled glove box. After cycling, each cell was opened in the glove box and washed with dimethyl carbonate to remove residual electrolyte and any other impurities. Then it was dried at room temperature. EIS measurement was carried out between 10000 – 0.1 Hz with an amplitude of 10 mV at a fully-delithiated state (~ 1.2V).
24 2.3 Results and discussion
The typical set up for coaxial electrospinning is illustrated in Fig. 2-1a. Two types of liquids were fed into a metallic dual nozzle with two channels connected to a spinneret.
For the inner core part, heavy mineral oil was fed at a rate ranging from 0.5 to 5 μL min-1. For the outer sheath part, we prepared a polymer solution containing poly(vinyl pyrrolidone) (PVP) and tetraethoxysilane as a silica source in absolute ethanol. Under a moderate electric field of 20-23 kV, core@shell (oil@polymer/SiO2) nanofibers were drawn out. After oil extraction by n-hexane and subsequent calcination, pure SiO2
nanotubes were successfully produced (ESI, Fig. S1 and S2a). Next, titania (TiO2) was uniformly coated on the inner and outer surface of the SiO2 nanotubes by means of coprecipitation (ESI, Fig. S1 and S2b). Then, TiO2-coated SiO2 nanotubes were chemically reduced via aluminothermic reduction reaction. Finally, we synthesized nanotubular Si- based multicomponents consisting of Si/Al2O3 and titanium silicide coating layers (Fig. 2- 1b). The final product, titanium silicide-coated Si-based multicomponents, was firstly confirmed by X-ray diffraction (XRD) analysis (Fig. 2-1c). The XRD patterns showed the presence of Si (85.1 wt%), Al2O3 (9.5 wt%), and two types of titanium silicide (Ti5Si3/TiSi2, 5.4 wt%) in the Si-based multicomponents.
Figure 2-1d-1g show low-magnified and high-magnified (Inset) scanning electron microscopy (SEM) images of samples obtained from each process described in the schematic illustration (Fig. 2-1b). The detailed mechanism of the coaxial electrospinning process will be discussed later. With high stability, heavy mineral oil was stuck in the core part of tubes, which shows clear phase separation from the shell part containing PVP and silica source (Inset of Fig. 2-1d). Pure SiO2 nanotubes were produced with smooth surfaces and dense shells after oil extraction and calcination process in air (Fig. 2-1e). Then, the TiO2 coating process was simply carried out by the co-precipitation method with the help of pre-coated polysaccharide (hydroxyl propyl cellulose, HPC). Firstly, numerous hydroxyl groups of HPC were strongly anchored to the surface of a SiO2 nanotube.
Secondly, a titanium precursor, titanium butoxide, was incorporated into the HPC covered sites in the SiO2 nanotube and subsequent addition of deionized water led to the formation of TiO2-coated SiO2 nanotubes. When the obtained samples were thermally treated in air, the HPC polymer was completely removed while crystalline TiO2-coated SiO2 nanotubes were obtained (Fig. 2-1f and ESI, Fig. S1 and S2b). Crystalline phase of TiO2 was matched with anatase after thermal annealing (ESI, Fig. S1). This will be converted to rutile phase
