Shape-anisotropic cobalt-germanium-borate glass flakes as novel Li-ion battery anodes

Journal Pre-proof Shape-anisotropic cobalt-germanium-borate glass flakes as novel Li-ion battery anodes

Julian D. Esper, Ying Zhuo, Maissa Barr, Tadahiro Yokosawa, Erdmann Spiecker, Dominique de Ligny, Julien Bachmann, Wolfgang Peukert, Stefan Romeis PII:

S0032-5910(19)31024-1

DOI:

https://doi.org/10.1016/j.powtec.2019.11.063

Reference:

PTEC 14945

To appear in:

Powder Technology

Received date:

4 June 2019

Revised date:

11 October 2019

Accepted date:

20 November 2019

Please cite this article as: J.D. Esper, Y. Zhuo, M. Barr, et al., Shape-anisotropic cobaltgermanium-borate glass flakes as novel Li-ion battery anodes, Powder Technology(2018), https://doi.org/10.1016/j.powtec.2019.11.063

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© 2018 Published by Elsevier.

Journal Pre-proof

Shape-Anisotropic Cobalt-Germanium-Borate Glass Flakes as Novel Li-Ion Battery Anodes Julian D. Espera, Ying Zhuob, Maissa Barrb, Tadahiro Yokosawac, Erdmann Spieckerc, Dominique de Lignyd, Julien Bachmannb,e, Wolfgang Peukerta and Stefan Romeisa,* a

Institute of Particle Technology (LFG)

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Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU)

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Cauerstr. 4

Institute for Chemistry of Thin Film Materials

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b

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91058 Erlangen, Germany

Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU)

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Egerlandstr. 1

c

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91058 Erlangen, Germany

Chair of Micro- and Nanostructure Research (IMN) and Center for Nanoanalysis and

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Electron Microscopy (CNEM)

Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU) Cauerstr. 6

91058 Erlangen, Germany

d

Institute of Glass and Ceramics

Friedrich-Alexander-Universität Erlangen-Nürnberg (FAU) Martensstr. 5 91058 Erlangen, Germany

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Institute of Chemistry

Saint-Petersburg State University Universitetskii pr. 26 198504 St. Petersburg, Russia 1

Journal Pre-proof Corresponding Author *

E-mail: [email protected]

Tel.; +49 9131 85-29421 Fax.; +49 9141 85-29402

Highlights

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Keywords: anode, particle design, glasses, shape-anisotropy, lithium-ion battery

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Graphical Abstract

Abstract

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Novel CoO-Li2O-B2O3 and CoO-GeO2-Li2O-B2O3 flake-like glass particles for battery

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applications are produced by an innovative technique: Glasses from melt quenching are processed into micron-sized glass flakes in a stirred media mill. Processing under well-

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controlled conditions yields flake-like particles with high aspect ratios which are suitable for anodes in Lithium-ion batteries. Electrochemical characterization by cyclic voltammetry, electrochemical impedance spectroscopy, galvanostatic charge/discharge cycles and rate capability measurements indicates that the flakes show superior electrochemical properties compared to the bulk glass counterparts with a stable specific capacity of 620 mAh g-1 after 100 cycles. The majority of charge is stored by Li+ ion diffusion inside the material and impedance spectroscopy shows enhanced lithium ion diffusion capabilities of the flakes. The proposed fabrication process is performed under ambient conditions and is simple, costefficient, fully scalable and can be easily transferred to other glass compositions with electrochemically active metal oxides.

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Journal Pre-proof 1. Introduction Lithium-ion batteries (LIBs) offer an unrivaled combination of high energy and power density. This combination makes them the electrochemical storage technology of choice for portable consumer electronics, automotive applications and stationary storage applications. Li-ion batteries have been introduced into the market by Sony in 1991.[1,2] Cell designs based on graphite electrodes with a specific capacity of 372 mAh g-1 have, however, limitations which result in high cost and range restriction of electric vehicles compared to their fuel

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counterparts.[1] Substantial progress has been made in the recent years and new active

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material systems and electrode designs have been introduced.[2] The crystalline materials

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discussed include a wide variety of metal oxides such as SnO[3,4], Fe2O3[5,6], Co3O4[7–11],

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TiO2[12–15] or ZnO[16] and semiconductors (e.g. Si[17], Ge[18]). Typically, crystalline materials suffer from irreversible phase transformations and volume expansion during

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charging and discharging. Amorphization of the material results and leaching of the redox-

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active metal ions into the surrounding electrolyte occurs. The phase transitions also induce a non-uniform volume expansion, which leads to the development of intrinsic strains and

effect.[19,20]

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microcracking.[19] Hence, the volume expansion often leads to a pronounced capacity fading

Various research strategies are being pursued to address the previously mentioned drawbacks and to develop new electrode materials and designs. These can be categorized[21] into: i) Size reduction to tackle mechanical stresses and improve ion/electron transport[22]; ii) formation of composites to add mechanical support and introduce conductive media for increased electron transport[23–25]; iii) morphology control for improved structural stability and improved ion/electron transport.[7,24] With respect to size reduction and shape control for crystalline materials Poizot et. al.[22] and Huang et al.[22,26] showed that a pronounced particle size reduction can lead to an increased particle aggregation during electrochemical cycling. This results in a poor electrochemical performance.[22,26] Therefore, an optimum 3

Journal Pre-proof particle size is essential for the performance of new electrode materials. Yao et. al.[9,27] and Su et. al.[7] reported that altering the particle shape results in different electrochemical performance. Especially crystalline cobalt oxide platelets offered an enhanced electrochemical storage performance. Therefore, particle size, shape and morphology have to be chosen carefully to improve the electrochemical performance of active materials.[7,27] Several groups have proposed various bottom-up synthesis methods for amorphous metal oxide materials suitable for electrochemical storage purposes.[28],[29–34] Uchaker et al.

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reported that amorphous electrode materials can be electrochemically more active than their

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corresponding crystalline phase[33] and currently amorphous or glassy active materials

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generate significant research interest. This is seen to be related to the high mobility of lithium

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ions and atoms in the amorphous structures.[32,35] Based on these considerations, glass systems hold great promise as their open structure may enhance Li-diffusion due to the lack of

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long-range order.[33] Bates et al.[36] showed in 1997 that even the electrolyte can be based

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upon glass systems which lead the way for an all solid state battery. Afyon et. al successfully demonstrated a vanadate-borate system[34] and Hayashi et. al tin oxide-borate glasses[37,38]

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for cathodes. These results demonstrate that a envisioned solid state battery purely based on glass materials is possible.[39] In case of the electrodes, the active material provides the necessary oxidation states for Li storage the amorphous matrix absorbs the accompanying volume expansion. So far, the size of all investigated glass particles lies well in the micrometer size range for melt quenched glasses[34] or slightly below 1 µm for glasses produced by spray pyrolysis.[23] In this work, we present a novel and fully scalable top-down approach to produce glass flakes with lateral diameters of several micrometer in diameter and flake thicknesses below 250 nm from active borate glasses made by traditional glass melting and quenching route. The glasses contain electrochemically active cobalt and germanium oxide. These particulates offer, in addition to the benefits associated with their amorphous nature, short lithium ion diffusional 4

Journal Pre-proof paths, a large electrolyte-surface interface that may improve cell kinetics and a tendency to form dense layers upon electrode formation.[9,27,34] To the best of our knowledge, wet compressed cobalt containing borate-germanate glasses have not been proposed as electrode materials for electrochemical storage applications. The highly promising route is not limited to the presented glass compositions, i.e. also glasses for cathodes or electrolytes could be processed similarly. Moreover, wet comminution is a widely applied, well-understood and scalable unit operation in material science and particle technology for production of micron-

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and even nano-size minerals, ceramics, foods, drugs and for the procession of battery active

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materials.[40–44] With respect to glasses Romeis et al showed the importance of the right

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solvent in order to preserve the chemical composition during grinding. For silica glasses the

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flake formation is attributed to the brittle-to-ductile transition:[45–47] Particles with sizes below the brittle-to-ductile transition (for silica around 1 µm) are deformed plastically without

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further breakage. The proposed glasses are prepared via traditional glass melting and

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quenching followed by stressing in a stirred media mill to obtain shape-anisotropic flake-like particles with high aspect ratios. The effectiveness of the melting technique and wet grinding

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process was assessed by electrochemical studies. X-ray diffraction and Fourier transformed infrared spectroscopy were performed to analyze the structure of the glasses. Nitrogen sorption measurements and He pycnometry as well as scanning electron microscopy were used to characterize the morphology and particle properties. Transmission electron microscopy and energy dispersive x-ray spectroscopy were carried out to validate the homogeneous incorporation of electrochemically active metal oxides. Cyclic voltammetry, electrochemical impedance spectroscopy, galvanostatic cycling and rate capability were used to identify and analyze the electrochemical properties and kinetic of the glass electrode material.

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Journal Pre-proof 2. Experimental Section 2.1 Materials Cobalt (II,III) oxide (Co3O4, 99.5%, Sigma Aldrich, USA), Germanium (IV) oxide (GeO2, ≥99.99%, Sigma Aldrich, USA) and lithium metaborate (LiBO2, Sigma Aldrich, USA) were used without further purification. 1-Pentanol (analytical grade, Merck, Germany) was used as

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dispersion medium.

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2.2 Glass Melting and Comminution Procedure

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Both glasses were obtained via a simple melt quenching technique. Co3O4, GeO2 and LiBO2 powders in predetermined molar ratios were thoroughly mixed in a tumbling mixer (T2F,

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Willy A. Bachofen AG, Germany) for 30 minutes and transferred into a platinum crucible.

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The crucible was heated (100 °C / h) in a muffle furnace to 900 °C (CoO-Li2O-B2O3 glass; denoted as CBO) or 1200 °C (CoO-GeO2-Li2O-B2O3 glass; denoted as CGBO) and kept at

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this temperature for 3 h to obtain a homogeneous and bubble-free melt. The melts of both

listed in Table 1. Table 1

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glasses were quenched between two solid brass plates. Characteristics of the materials are

Chemical identity and bulk properties of the glasses as well as comminuted products were thoroughly checked via He pycnometry, ICP OES and X-ray diffraction. In the next step, the obtained glasses were pre-crushed with pestle and mortar and comminuted in the stirred media mill PE075 (NETZSCH Feinmahltechnik, Germany). The stressing conditions are summarized in Table 2. The temperature was set to 15 °C to minimize the evaporation of n.pentanol at higher temperatures as well as viscosity increase and coupled dampening during milling at lower temperatures. No significant size reduction was observed at longer processing times. Therefore, all further experiments were conducted with samples milled for 4 h. After 6

Journal Pre-proof wet grinding, the materials were washed thoroughly in ethanol to remove pentanol residues and dried at 60 °C for several hours and heated at 400 °C for 2 h. XRD characterization revealed no crystallization of the bulk glass as well as glass flakes below 500 °C (Figure 4b and 4c). Table 2.

2.3 Electrode Preparation

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The working electrodes were prepared by dispersing the prepared shape anisotropic glass

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flakes (68 wt.-%), carbon black (Super C65, Timcal Ltd. Belgium, 17 wt.-%) and Na-alginate

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binder (15 wt.-%) in H2O to form an easy to handle slurry. The slurry was pasted onto a

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cleaned copper foil using a doctor blade (Erichsen Testing Equipment, COATMASTER 510, Germany; coating velocity of 0.1 mm s-1). After drying at room temperature the prepared

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electrode was transferred to a preheated oven and was kept at 80 °C for several hours. Prior to

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the electrochemical characterization circular electrodes with a diameter of 14 mm were punched out and button-type (CR-2032) half-cell configurations were prepared and assembled

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in an argon-filled glove box. Lithium metal foil was used as the counter electrode and Celgard 2400TM as the separator. The electrolyte solution was 1 M LiPF6 dissolved in mixture of ethylene carbonate (EC) and dimethyl carbonate (DEC) in a 50/50 wt-.% ratio. 2.4 Characterization Methods 2.4.1 Scanning Electron Microscopy (SEM) SEM imaging was performed using an Ultra FE-SEM (Gemini Ultra 55, Carl Zeiss Jena, Germany). The accelerating voltage was set to 1.8 kV and an Everhart-Thornley (SE2) detector was used. Samples of the particulates were prepared by suspending dried powders in ethanol on a Si-wafer. SEM image analysis was concluded by measuring a minimum of 250 particles to determine the average Feret diameter as well as the flake thickness.

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Journal Pre-proof 2.4.2 Nitrogen Sorption Measurements Nitrogen sorption measurements at liquid nitrogen temperature were performed using a volumetric gas sorption analyzer (Nova 4200e, Quantachrome, Germany). The mass-specific surface areas (Sm) were determined from dried and degassed (2 h at 200 °C under vacuum) powder samples following the BET (Brunauer-Emmet-Teller) theory.[48] The relevant relative pressure range (0.02 – 0.35 p/p0) was subdivided into 5 evenly spread data points. The

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samples were also checked for micro- and mesoporosity. 2.4.3 Helium Pycnometry

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For measurements of the true powder density, a helium pycnometer (AccuPyc 1330,

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Micromeritics, USA) was used. All dried powders were kept at 60 °C prior to analysis. The

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2.4.4 Powder X-Ray Diffraction (XRD)

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reported density value is the mean of three independent and separate measurements.

Structural analysis was performed using a powder diffractometer in Bragg-Brentano setup

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(Empyrean, Malvern Pananalytical, United Kingdom). The X-ray diffractometer was

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equipped with a GaliPIX3D detector. All shown X-ray patterns were recorded with Cu Kɑ radiation (λ = 0.15405 nm) in the 2theta range 10° to 90° and a step width of 0.014°/step. Background removal was performed by fitting a polynomial function. Kɑ2 was stripped mathematically from the obtained patterns. 2.4.5 Inductively Coupled Plasma Emission Spectroscopy (ICP OES) Quantitative chemical analysis was performed using inductive coupled plasma optical emission spectroscopy (Optima 8300, Perkin Elmer, USA). Typically, 0.1 g of sample was digested in HCl (aq) / HF (aq) / HNO3 (aq) (6/2/2 wt.-%) via microwave digestion. Measurement conditions were: sample flow rate of 1.5 mL min-1, an Argon-plasma power of 1400 W and flow rates of 10 L min-1, 0.6 L min-1 and 0.2 L min-1 for the plasma (Ar)-, nebulizer (Ar)- and auxiliary (N2) gas. Five point calibrations were performed for 8

Journal Pre-proof concentrations up to 100 mg L-1 for all analytes (diluted from standard specifically made for ICP, 1000 mg L-1, Carl Roth, Germany). 2.4.6 Fourier-Transformed Infrared Spectroscopy (FTIR) Infrared spectra (FTS 3100 Spectrometer, Varian, USA) were recorded in the range of 4000 – 400 cm-1 with a spectral resolution of 4 cm-1. Suitable samples for the measurements in transmission geometry were prepared by mixing and tableting about 300 mg dry KBr

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(UVASol, Merck) with 1 mg of the sample. The recorded data was subjected to mathematical

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smoothing and normalization. Occurring vibrational modes of the samples were matched with

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modes from available publications[10,49–52].

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2.4.7 Cyclic Voltammetry (CV)

Cyclic voltammetry measurements were concluded using a potentiostat (Interface 1000,

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Gamry Instruments, USA) in a potential window of 0-3 V (vs. Li/Li+) at a scan rate of

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0.25 mV/s. All electrodes had an area of 1.54 cm2 and were used in a CR2032 coin-type cells. Additionally, two CV cycles at different scan rates (0.25, 0.5, 1, 2, 5, 10, 20, 50, 100 mV s-1)

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were concluded after the initial 10 cycles at 0.25 mV/s. 2.4.8 Electrochemical Impedance Spectroscopy (EIS) Electrochemical impedance spectroscopy measurements were performed using a potentiostat (Interface 1000, Gamry Instruments, USA) over a frequency range of 0.01 – 500 kHz by applying an AC voltage with amplitudes between 0.1-3.0 V. 2.4.9 Galvanostatic Charge-Discharge Measurements Galvanostatic charge-discharge measurements were conducted in the voltage range of 0 – 3.0 V at constant rates of 125 mA g-1 for 100 cycles using a battery test system (LANHE CT2001A, Wuhan LAND Electronics Co., China). Each galvanostatic charging step was followed by a potentiostatic charging step until the resulting current fell below a set minimum value of 25 µA. 9

Journal Pre-proof 2.4.10 Rate Capability Measurements Rate capability measurements were done in sequential blocks of 10 cycles each at rates of 62.5, 125, 250, 500, 1000, 500, 125 and 62.5 mA g-1 using a battery test system (LANHE CT2001A, Wuhan LAND Electronics Co, China). The reported specific capacities were calculated based on the mass of the glass flakes. A conversion to a specific capacity based solely on the amount of cobalt oxide and germanium oxide was also performed.

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2.4.11 Transmission Electron Microscopy (TEM)

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Transmission electron microscopy and scanning transmission electron microscopy with energy-dispersive X-ray spectroscopy (STEM-EDX) were performed by using a double-

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corrected FEI Titan Themis3 300 transmission electron microscope. The microscope was

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operated at an acceleration voltage of 300 kV. The TEM sample was prepared by mixing the

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powder sample with ethanol, dispersing the solution on a copper grid coated with a lacey carbon film and letting it dry overnight before investigation.

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2.4.12 X-ray photoelectron spectroscopy (XPS)

X-ray photoelectron spectroscopy was conducted using a a PHI Quantera II system (ULVAC-

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PHI Inc., Japan) with a base pressure of 1 x 10-9 mbar. A combination of electron and ion netralization was employed to prevent charging. The Co, Ge and O XPS core level spectra were analyzed using a fitting routine that decomposes each spectrum into individual mixed Gaussian-Lorentzian peaks using a Shirley background substraction over the energy range of the fit. Finally, all spectra were shifted to yield a C 1s binding energy position of 284.8 eV. 2.4.13 Differential scanning calorimetry (DSC) The glass transition temperature and crystallization temperature of the bulk glass and glass flakes was determined via a differential scanning calorimeter (Netzsch DSC 404 F1 Pegasus, Germany) at a upscan rate of 20 °C/min up to a temperature of 600 °C. About 15 mg of bulk glass, finely ground prior to DSC characterization, as well as 15 mg of glass flakes, were placed in a Pt crucible. Before each measurement, an empty crucible was heated and cooled 10

Journal Pre-proof to measure the baseline and a calibration measurement was performed with a crucible

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containing 28 mg of sapphire.

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Journal Pre-proof 3. Results and Discussion 3.1 Glass Melting and Platelet Formation The prepared melt-quenched glasses were precrushed with pestle and mortar, which yields shard-like particles with a mean size in the upper micrometer range, to simplify handling of the active material for the subsequent stressing inside the batch-type stirred media mill. To prevent leaching of glass constituents into the comminution solvent, n-pentanol was chosen. [44–46] The processing temperature of 15 °C has been chosen as a good compromise between

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viscosity, which increases with decreasing temperature, and evaporation of the solvent. An

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increased viscosity due to lower processing temperatures would lead to a dampening effect

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and hence hinder the stressing of the glass particles. The grinding time was set to 4 hours, as

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longer times did not lead to further increase of the mass specific surface area and long processing times can lead to significant amounts of unwanted wear of the grinding media.

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After wet grinding the blue color is clearly visible for both glasses. The dark blue color of the

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as-quenched CoO-Li2O-B2O3 (CBO) and CoO-GeO2-Li2O-B2O3 (CGBO) glasses indicate successful cobalt incorporation into the germanium-borate and borate glass matrices

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respectively. Because of the high temperatures of the glass melt, Co3O4 decomposes to CoO and O2, resulting in a strong blue color also known as cobalt blue [53,54]. The crushed glasses as well as scanning electron micrographs of the resulting CBO (a-c) and CGBO (d-f) glass flakes are depicted in Figure 1. After wet compression, the particles are flake-like and exhibit smooth surfaces with thicknesses well under 250 nm. A similar mechanism for structurally much simpler silica particles has been investigated by Romeis et al. who found a brittle-toductile behavior below a particle size around 1 µm.[47] This mechanism holds obviously true for the stressing of borate-germanate-based glasses as well since flake-like particles with lateral dimensions in the micron-range and average thicknesses below 250 nm are formed.

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Journal Pre-proof Figure 1. Comminuted and dried glass flake powders CBO (a) with corresponding SEM micrographs (b) side view (c). Image of CGBO glass flakes (d) with SEM (e) and side (f).

High-angle annular darf-field scanning transmission electron microscopy (HAADF-STEM) analysis (Figure 2a) reveals compact flake-like particles with relatively sharp edges. This is exemplarily shown for the CGBO glass flakes in Figure 2 (CBO glass is shown in Figure S1, Supporting Information). Selected-area electron diffraction (SAED) shows a halo pattern (Figure 2b), confirming the amorphous structure of the material. Elemental mapping by

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STEM-EDX (Figure 2c, 2d 2e and 2f) reveals homogeneously distributed germanium, cobalt,

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oxygen and boron across the sample. No aggregated cobalt clusters were detected. No burning

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or degradation of the material due to the electron beam illumination was observed.

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Figure 2. a) HAADF-STEM image of a CGBO glass flake; b) SAED pattern taken from the glass flake; STEM-EDX elemental mapping of an area of the glass flake; elemental mapping of Ge (c), Co (d), O (e) and B (f).

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SEM image analysis was performed to determine the average Feret diameter and flake thick-

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ness by manually measuring a minimum of 250 particles for each characteristic dimension. The average mean maximum Feret diameter xFe of the plates for both glasses is almost identical and account for approximately 7 µm (Figure S2a, Supporting Information) 6.9 µm and 6.7 µm respectively and is given in Figure 3a. The average flake thickness davg is <250 nm for CBO and CGBO (Figure S2b, Supporting Information). This unique morphology makes the glass flakes suitable as an anisotropic active material for electrochemical cells. The element composition was investigated by X-ray photoelectron spectroscopy. Figure 3 summarizes the results of the fitted XPS spectra. The survey spectrum (Figure 3a) shows the presence of elements like Co, Ge, O, B and C indicating the purity of the sample. The highresolution XPS spectra of C 1s (Figure 3b) indicates the presence of C 1s and can be deconvoluted into 3 peaks located at 284.7, 286.2 and 289.0 eV corresponding to the 13

Journal Pre-proof formation of C=C, C-C and C(O)O. The high-resolution spectrum of O 1s shown in Figure 3c can be deconvoluted in 3 peaks located at 528.5, 531.5 and 533.7 eV. The peak at 32.6 eV can be assigned to Ge 3d indicating the presence of GeO2 (Figure 3d). The two main peaks and their two corresponding satellite peaks with binding energies of 781.4, 786.4, 797.4 and 802.7 can be assigned to Co 2p3/2 and Co 2p1/2 respectively (Figure 3e). XPS spectra for the bulk CGBO glass is given in Figure S3 (Supporting Information) and shows no significant change in elemental composition and oxidation state during the processing in a stirred media mill.

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Nitrogen sorption experiments showed a type III isotherm (Figure 3f) without any indications

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of pores. The mass-specific surface areas SM of the glass flakes were determined to 17.3 m2 g-1

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and 16.0 m² g-1 respectively for CBO and CGBO glass flakes.

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Figure 3. a) XPS survey spectrum of the CGBO glass flakes (a), the high-resolution spectra of the C 1s region (b), O 1s region (c) Ge 3d region (d), Co 2p region (e) and full nitrogen sorption isotherms of CBO and CGBO glass flakes (f).

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Quantitative elemental analysis by ICP OES gave a content of 15 mol.-% CoO and 85 mol.-%

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B2O3-Li2O in the CBO material. For the CBGO glass, 19 mol.-% CoO, 25 mol.-% GeO2 and 56 mol.-% B2O3-Li2O were measured. No changes in the chemical composition or structure were detected due to the processing in a stirred media mill. The structures of the feed materials and the obtained glass powder were assessed using powder X-ray diffraction. The patterns of the two glasses are given in Figure 4. XRD patterns of the starting materials Co3O4 nanoparticles, GeO2 and LiBO2 show the typical material-specific diffraction reflexes (Figure S4, Supporting Information) confirming their crystalline structure. For both glasses, however, no detectable Bragg reflexes are visible, i.e. the materials are amorphous. This vitreous nature is conserved during the shape formation inside the stirred media mill and confirmed by XRD patterns and DSC measurements given in Figure 4. No crystallization was observed for CGBO bulk glass and glass flakes for temperature <500 °C (Figure 4b and 4c). 14

Journal Pre-proof Figure 4d depicts infrared (IR) spectra of the CoO-Li2O-B2O3 and CoO-GeO2-Li2O-B2O3 glass. Clearly, the infrared spectrum of glass differs from the starting materials. Co3O4 shows two distinct bands at 572 and 661 cm-1 which are associated with the OB3 (B represents Co3+ in an octahedral hole) and the ABO (A represents Co2+ in a tetrahedral hole) spinel lattice vibrations.[10,55,56] A small distinct band at ~670 cm-1 can be seen for both glasses. This band can be associated associated with the characteristic band of cobalt blue glass and indicates a succesful incorporation of cobalt in a glass matrix of GeO2 and borate and the

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position of the band is influenced by the glass composition [53]. These two bands disappear

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completely in the quenched glasses indicating the successful incorporation of Co into the

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amorphous glass matrix. The crystalline LiBO2 spectrum (Figure S4, Supporting Information)

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shows absorption peaks in the broad region 1150-1450 cm-1 which can be attributed to B-O bond stretching of trigonal BO3 units while the absorption peaks in the region 800-1100 cm-1

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are assigned to B-O bond stretching of BO4 units.[49–52,57] The addition of alkali ions tends

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to modify the glass network altering trigonal BO3 groups to BO4 groups, therefore increasing the number of structural linkages between B–O–B bonds in the glass network.[34,58] For the

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incorporation of GeO2 inside the CGBO glass matrix, a peak at around 450-500 cm-1 and a shoulder at 1050 cm-1 can be observed. These have been reported by Galeener et al. and can be traced back to the symmetric stretching of bridging oxygens in 6-membered GeO4 rings (450-500 cm-1) and LO split asymmetric stretching of the bridging oxygens.[59] Figure 4e and 4f show the DSC curves for CGBO bulk glass and CGBO glass flakes respectively. The glass transition temperature and the onset of the crystallization can be determined by the DSC curves. The DSC curves prove the vitreous nature of both the CGBO bulk glass and CGBO glass flakes with their respective Tg temperatures at 476 °C and 475 °C. The DSC results show that the processing of the glass in a stirred media mill has almost no influence of the Tg of the glass, yet the onset of the crystallization moves to lower temperatures (591 °C to 550 °C). This result may be attributed to the increased mass specific surface area of the glass 15

Journal Pre-proof flakes compared to the bulk glass. Additionally, the DSC curves confirm the XRD results, showing no crystallizations at temperatures <500 °C and complete crystallization at temperatures of 600 °C.

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Figure 4. a) Powder XRD patterns of CGBO and CBO bulk glass; b) XRD patterns bulk CGBO glass heated at different temperatures; c) XRD patterns of CGBO glass flakes heated at different temperatures; d) FTIR spectra of CGBO and CBO glass; e) DSC curve of CGBO bulk glass and CGBO glass flakes (f) during heating. The determination of Tg and Tc for both bulk and glass flakes is illustrated respectively.

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3.2 Electrochemical Characterization

Electrochemical performance of the glass flakes was examined by assembling button-type

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(CR-2032) half-cells. Figure 5 displays a photo of the coated and dried active material layer

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on the copper foil (a), a punched out electrode (b) and a SEM side view image of the electrode

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(c). A homogeneous film with a high packing density is obtained and neither macro- nor

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microscopically peeling or cracking of the active material layer is observed.

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Figure 5. a) Coated and dried copper foil of CGBO glass; b) punched out electrode of CGBO glass; c) SEM side view of the dried electrode with CGBO glass.

Figure 6 shows the cyclic voltammetry (CV) cycles in a potential region between 0-3 V at a scan rate of 0.25 mV s-1. The CV curves in Figure 6a-b show nearly ideal features with flat baselines and a good separation between the forward and backward scans.[60] For the CBO electrode, the cathodic current is expected to arise from lithium insertion into the active material while the anodic current results from lithium extraction from the cobalt oxide active material.[13] In the first cathodic scan, a small but broad feature between ~ 1.4-0.9 V vs. Li+/Li can be observed. This can be attributed to the decomposition of the electrolyte whereas the cathodic current peak from 0.9-0 V represents the formation of the solid-electrolyte interlayer (SEI) on the electrode surface. This SEI formation is linked to an irreversible 16

Journal Pre-proof capacity fading that can be observed during the first cyclic voltammetry cycles. One prominent anodic feature can be observed around 1.2 V. This peak may be attributed to the In the following cycles, one reduction peak arises at 1.1 V that indicates the reduction of CoO to metallic Co [61]. The anodic feature arising at ~ 1.25 V may be attributed to the oxidation of the glass matrix whereas the small hump around ~ 2.1 V may be assigned to the reformation of cobalt oxide (CoO) [61]. Nonetheless, both the anodic and cathodic currents versus the potential of both glasses are almost featureless in the first cycles. which is characteristic for

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amorphous materials.. The quite constant current density versus potential behavior has been

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shown for amorphous metal oxides and cobalt-containing electrodes before.[12,13,62,63] In

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case of the CGBO glass, during the first cycle, one prominent reduction peak can be observed

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at 0.5 V that can be attributed to the formation of the SEI layer. Small anodic features at 0.7 V, 1.25 V and 2.1 V arise during the first oxidation cycle and can be associated with the

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oxidation of metallic Ge and Co respectively [23,61,64]. More prominent cathodic (0.5 V, 0.9

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V, 1.1 V) and anodic features (at 0.6 V, 1.25 V and 2.1 V) arise in the following cycles (9), where the first two anodic peaks can be attributed to the formation and reformation of

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crystalline Ge-rich nanodomains, which has already been reported by Choi et. al. and Wang et. al. [23,65] and these multiple peaks suggest a multi-step lithium insertion and desertion mechanism inside germanium oxide that may be written as [66]: GeO2 + 4Li  2Li2O + Ge

(1)

Ge + 4.4Li  Li4.4Ge

(2)

Afterwards, the glass matrix stabilizes[64] and the overlapping in the following cycles (10-20) indicates a great reversible lithium insertion and extraction inside the glassy active material. Furthermore, the glass flakes show an improved electrochemical behavior compared to their bulk glass counterparts (Figure S5, Supportin Information). In order to distinguish the contributions of the surface and the volume to the overall Lithium ion storage capacity the 17

Journal Pre-proof current response during cycling voltammetry at different scan rates was recorded for the CGBO glass. This method was proposed by Trasatti et. al to separate the capacitive elements from diffusion-controlled Li+ storage at fixed applied potentials[67–69]. The method gives insight into the charge-storing mechanism of the material and can be divided into a capacitive behavior vs. a diffusion-controlled behavior, where the latter is a function of the diffusion rate and therefore describes the Li+ diffusion into the solid active material. Results for Trasattis

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method on CGBO glass flakes are shown in the Figures 6c, 6d and 6e.

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Figure 6. a) CV curve of the CBO glass flake electrode at a scan rate of 0.25 mV s-1; b) CV curves of the CGBO glass flake electrode at a scan rate of 0.25 mV s-1; c) CV curves of the CGBO glass flake electrode at different scan rates; d) The relationship of the area specific capacitance versus inverse square root of the scan rate; e) The relationship of the inverse area specific capacity vs. the square root of the scan rate, inset shows a magnification at smaller yaxis values.

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At high scan rates, the access of Li+ ions into the inner material becomes limited due to the

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slow diffusion. In this case, the electrochemical response depends exclusively on the “outer” active surface. The outer, capacitive charge is given in the extrapolation plot in Figure 6d and

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results in an electrode capacity of 0.265 mAh cm-2. At low scan rates, the system is given enough time to let Li+ ions diffuse inside the solid material to reach the “inner” material. The total charge of the electrode is given in the extrapolation plot in Figure 6e and the resulting capacity is calculated to be 1.856 mAh cm-2. This capacity includes capacitive surface and diffusional storage. This result shows that 86% of the total charge is based upon Li+ ion diffusion into the active material whereas only 14% of the charge can be traced back to a capacitive storage mechanism of the surface. Therefore, an improved diffusional behavior into the solid electrode material is of importance. The cycle stability of the cells prepared from the novel electrode glass materials was examined at a current density of 125 mA g-1 in a voltage range of 0-3 V which corresponds to the standard voltage range for these types of materials[7,67,70–73]. In total, 100 galvanostatic 18

Journal Pre-proof charge/discharge cycles were carried out. In case of the CGBO glass flake electrode, the first ten cycles were carried out at a current density of 50 mA g-1 to give the electrode enough time to build up Ge rich domains and stabilize (as seen from previous CV measurements). Additionally, pure Li2O-B2O3 (LBO) glass flakes and pure CoO-GeO2 (CGO) glass flakes were produced and measured to show that the amorphous LBO matrix is electrochemically inert, but the addition is important to stabilize the active material during cycling as the capacity of the pure CoO-GeO2 glass flakes fades quickly. The small specific capacities of the

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LBO glass prove that in absence of Ge and Co, the amorphous glass compound is inactive for

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Li+ ion storage. All galvanostatic cycling results are depicted in Figure 7. CBO glass flakes

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deliver a first specific discharge capacity of 400 mAh g-1 that decreases during the first 10

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cycles and stabilizes around 190 mAh g-1 and retains 80 % of its specific capacity after 100 cycles. The CGBO glass flake electrode delivers a first discharge capacity of 572 mAh g-1 that

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settles to a stable specific capacity of 620 mAh g-1 after the formation of presumably

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crystalline Ge nanodomains during the first ten cycles[23] and relaxation of the glass matrix (Figure S6, Supporting Information)[64]. The initial coulombic efficiency (CE) of 77 % and

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linked irreversible capacity loss can be attributed to the formation of the SEI layer and decomposition of the electrolyte. After the second cycle, the CEs gradually increase to approximately 100 %, demonstrating the good charge-discharge reversibility of the CGBO glass material. This result has be observed and reported with other GeO2, Co3O4 and glass anode materials[71,73–75]. The altered shape and reduced size of the CGBO glass and therefore improved surface-to-volume ratio, due to the stressing, results in an enhanced cycle stability with higher capacities compared to its bulk glass counterpart (Figure 7a). It is worth mentioning that pure GeO2 glass and CoO-GeO2 glass also forms flake-like particles (Figure S7, Supporting Information) upon stressing in a stirred media mill, although both glass flakes, show a strong capacity fading during 100 cycles resulting in low specific capacities (Figure 7b). Therefore, the addition of Li2O and B2O3 helps improving the cyclic stability by 19

Journal Pre-proof buffering the detrimental volume expansion. The theoretical specific capacity purely based upon GeO2 and CoO is calculated to be 1250 mAh g-1 (theoretical specific capacity of CoO: 715 mAh g-1 and GeO2: 2200 mA g-1). The CGBO glass offers a reversible specific capacity (based upon CoO and GeO2) of 909 mA g-1 after 100 cycles (Figure 7b). The rate capability of lithium-ion batteries is the major parameter for battery applications which require fast discharge and charge rates. Hence, the rate performance of the CBO and CGBO glass flakes was investigated in detail by cycling glass flake electrodes at various

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current densities (125, 250, 500, 1000 and 1500 mA g-1). The resulting galvanostatic charge-

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discharge rate capability is summarized in Figure 7c. Similar to the 100 galvanostatic charge-

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discharge cycles at 125 mA g-1, the first 15 cycles of the CBO glass flakes are affected by

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irreversible SEI layer formation. The CBGO glass flake electrode shows an increase of the specific capacity during the first 15 cycles until the capacity stabilizes around 620 mAh g-1

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(Figure 7a and 7c).

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with an enhanced capacity stability and rate capability over its feed bulk glass counterpart

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Figure 7. a) Cycling performance of CBO glass flakes and bulk glass as well as CGBO glass flakes and bulk glass 125 mA g-1; b) Cycling performance of CGBO glass flakes, CGO glass flakes, GeO2 glass flakes and LBO glass flakes at 125 mA g-1; c) Rate performance of CGBO glass flakes, CBO glass flakes and their bulk glass counterparts; d) Different cycle chargedischarge profiles of CGBO glass flakes at 50 mA g-1 between 0-3 V; e) Different cycle charge-discharge profiles of CBO glass flakes at 50 mA g-1 between 0-3 V.

As compared to its bulk glass counterpart, the CBO glass flake electrode shows an enhanced rate capability with average discharge capacities of approximately 350 mAh g-1, 219 mAh g-1, 116 mAh g-1, 90 mAh g-1 and 26 mAh g-1 at rates of 125 mA g-1, 250 mA g-1, 500 mA g-1, 1000 mA g-1 and 1500 mA g-1, respectively. The coulombic efficiencies are again close to 100%. For comparison a similar measurement for the pre-crushed bulk glass powder of the CBO glass is given. In this case significantly lower specific capacities were found for all investigated current densities (Figure 7c). The CGBO glass flake electrode shows superior 20

Journal Pre-proof rate capability with average discharge capacities of 610 mAh g-1, 485 mAh g-1, 350 mAh g-1, 335 mAh g-1 and 287 mAh g-1 at rates of 125 mA g-1, 250 mA g-1, 500 mA g-1m 1000 mA g-1 and 1500 mA g-1 respectively. These higher capacities can be traced back to the high amount of GeO2 that provides structural integrity as a glass forming agent as well as being electrochemically active to store Li+ ions. This result indicates that the structural integrity of the glass flake material has been preserved and that the material is tolerant to varying current densities. The improved rate capability of the flake-like particles compared to the bulk glass

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material is probably caused by the poor transport kinetics of the larger glass particles which

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results in a rapid decrease in specific capacity at high current densities (Figure 7c). On the one

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hand, small mass-specific surface areas with no porosity might be beneficial as the amount of

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formed SEI during the first couple of cycles is reduced. On the other hand, however, the surface area interacting with the electrolyte is small. This results in a higher reliance on

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lithium ion bulk diffusion inside the active material.[76] Therefore the shape formation in a

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stirred media mill may be beneficial due to the formation of thin flakes (~200 nm) while exhibiting relatively small mass-specific surface areas <20 m2 g-1.

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Figure 7d and 7e show the first, tenth and twentieth static charge/discharge curves of the CBO and CGBO glass flakes within a potential window of 0-3 V. The slopes of the charge/discharge curves in Figure 7d and 7e show an immediate drop in potential from approximately 3 V to 1.5 V after which it slowly diminishes in a featureless curve with no plateau comparable to reported results for amorphous SnO·CoO and (Co,Mn)3O4.[62,77] The characteristics of the curves do not change upon repeated cycling, indicating a homogeneous phase process of lithium insertion and extraction during the electrochemical cycling.[34] The CGBO glass flake electrode experiences the same potential drop from 3 V to 1.5 V but shows a small plateau between 0.5-0.1 V that can be attributed to the reduction of GeO2 to metallic Ge.[23]

21

Journal Pre-proof Electrochemical impedance spectroscopy (EIS) was carried out to further understand the lithium storage performance and kinetics of the two glass electrodes and determine the alternating current impedance of the assembled cell. Additionally, the Li+ diffusion coefficient was derived from the EIS measurement results. EIS measurements were performed from 50 kHz to 0.01 Hz. Nyquist plots recorded at a potential of 0.5 V of both glass compositions are given in Figure 8 in which Z’ and Z’’ are the real and imaginary part of the impedance, respectively. The data show an intercept at high frequencies that is associated with the Ohmic

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resistance of the cell (Re).[61,78–81,81] For CBO glass (Figure 8a) the two semicircles in the

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medium frequency range are related to the SEI layer impedance (RSEI) and the charge-transfer

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impedance (Rct) which arises due to electrode-electrolyte interface reactions.[61,78–80,82] In

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case of the CGBO glass flakes (Figure 8b), only one depressed semicircle can be observed which means that the SEI layer passing and charge-transfer are running simultaneously. The

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SEI film prevents the solvated Li ion from being stored inside the active material in its

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solvated state, which would lead to an enhanced swelling of the material. The linear incline at low frequency is attributed to the diffusion of Lithium ions into the electrode material and is

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electrically described as a Warburg impedance (RW).[8,61,80] The steeper the straight line in the low frequency region is, the more rapid the Li + ion diffusion inside the electrode material is.

Figure 8. a) Nyquist plot of CBO glass flakes at 0.5 V vs. Li/Li+; b) Nyquist plot of CGBO glass flakes at 0.5 v vs. Li/Li+ c) Variation of pre-exponential factor AW of the Warburg impedance as a function of electrode potential. The CBO electrode investigated exhibits a charge-transfer resistance (Rct) of ~6 Ω at 0.5 V after 20 cycles of CV whereas the CGBO electrode a charge transfer resistance (Rct) of 80 Ω at 0.5 V. This result may be attributed to the increased GeO2 content over Li2O-B2O3 which is in good agreement with reported values of crystalline GeO2 electrodes.[83] Yet, both glasses 22

Journal Pre-proof show good charge-transfer kinetics during lithium insertion and extraction reactions and show improved charge transfer resistances and Warburg impedances towards their feed bulk glass counterparts (Figure S8, Supporting Information). This result may be due to the beneficial active material structure and enhanced lithium diffusivity of the electrode when compared to reported values for crystalline cobalt oxide electrodes.[8,61] Randles circuits, which resemble equivalent electrical circuit models (inset Figure 8a-b), were fitted to the Nyquist plots (Figure 8a-b) and simulated with the commercially available software ZSimWin (Version

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3.20). Additionally, the incline of the linear Warburg impedance diagram at low frequencies

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was investigated (Figure S9, Supporting Information). The Li+ diffusion coefficient is

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proportional to (1/AW)2 and therefore the Warburg impedance can be used as an estimation of

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the Li+ diffusion behavior of the electrode.[84] Figure 8c depicts the different values of the Warburg coefficient AW (in Ω s-1/2) over different applied potentials for the comminuted

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(continuous lines) and the pre-crushed glasses (dashed lines). The AW values obtained show

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that the lithium ion diffusion resistance of the glass flakes decreases rapidly as compared to the feed bulk glass materials. The Warburg factors AW for potentials below 1.0 V are quite

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low for the flake-like particles, indicating highly mobile Li+ ions and a fast diffusion inside the bulk active materials during lithiation. Compared to the Warburg coefficient of the bulk glass material, the Li-ion diffusion resistance is reduced for both glasses by processing in a stirred media mill.

After 100 cycles of galvanostatic cycling, no microscopic detachment of the active material layer was observed. Furthermore, no recognizable particle size reduction or disintegration during electrochemical cycling was observed for either glass. This result indicates that the glass flake material architecture can sustain long-term electrochemical cycling with the accompanied volume expansion during the Li ion insertion (Figure S10, Supporting Information).

23

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4. Conclusions Within this account shape anisotropic glass flakes with high aspect ratios are proposed as promising electrode materials for Lithium ion batteries. The glass flakes are obtained by a combination of conventional glass quenching, size reduction down to the ductile-to-brittle transition with direct, subsequent wet compression in a stirred media mill in 1-pentanol. The

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advantages of the process are its full scalability, simplicity and applicability to other glass

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compositions. Herein, we have demonstrated the proposed approach for two glasses from the

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systems CoO-Li2O-B2O3 and CoO-GeO2-Li2O-B2O3 showing excellent electrochemical

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performance and good cycle stability for Li-ion batteries.

For both glasses XRD analysis proved the amorphous nature of the final glass powders. DSC

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measurements showed that the processing in a stirred media mill has no significant influence

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on the glass transition temperature Tg of both the bulk glass and the glass flakes. ICP OES, XPS and TEM characterization showed the successful and homogeneous incorporation of

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19 mol% of CoO and 25 mol% GeO2 in in a lithium borate glass matrix where GeO2 acted as an electrochemically active site as well as a glass matrix former. Even though Li2O and B2O3 are electrochemically inert, their addition to the glass composition is quite important as they stabilize the active material and buffer the detrimental volume expansion. Electrochemical characterization in the potential range of 3.0–0.0 V (vs. Li/Li+) revealed that the CGBO glass flakes need approximately 10-15 galvanostatic charge-discharge cycles to stabilize, which can be attributed to formation of an SEI layer, the relaxation of the borate glass matrix and formation of crystalline Ge rich nanodomains. The CGBO glass flakes stabilized around a stable specific capacity of 620 mAh g-1 at a current density of 125 mA g-1 with a capacity retention of >90% after 100 cycles. Further improvements to enhance the suitability for

24

Journal Pre-proof practical battery applications may be done via glass composition, electrode engineering as well as slurry processing and composition. The amorphous and inert matrix can be utilized to stabilize the redox active materials and prevent agglomeration during cycling while allowing for volume changes. The results clearly indicate that the studied materials and the corresponding fully scalable processing technique are highly promising triggering further research in the field of amorphous active material

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electrodes as next generation electrodes in LIBs.

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Author Contributions

ACKNOWLEDGEMENT

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approval to the final version of the manuscript.

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The manuscript was written through contributions of all authors. All authors have given

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The authors thank Dr. Jochen Schmidt for performing the ICP OES measurements. Jakob Iser

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is acknowledged for assisting with lab work and sample preparation. SR acknowledges German Research Foundation (DFG) for providing funding through the starting grant of GRK1896 “In situ Microscopy with Electrons, X-rays and Scanning Probes”.

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Journal Pre-proof Table 1. Characteristics of the used materials SSA [m2 g-1]

Density [g cm-3]

Molar Mass [g mol-1]

Co3O4

62.6

6.07

240.79

LiBO2

---

2.22

49.75

GeO2

---

4.23

104.59

Bulk glass (CBO)

0.49

3.03

59.30

Glass flakes (CBO)

19.76

3.05

59.30

Bulk glass (CGBO)

0.53

3.35

81.25

Glass flakes (CGBO)

16.03

3.36

81.25

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Material

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Journal Pre-proof Table 2. Set of parameters for the glass comminution Parameter

Experimental condition 6.50 m s-1

Stirrer speed vtip Grinding bead media

YTZ (95 wt.-% ZrO2 & 5 wt.-% Y2O3)

Grinding bead diameter xGM Solvent

2 mm 1-pentanol

Solids concentration cW

3 wt.-% 4h 15 °C

Grinding bead mass mGM

1.9 kg

Chamber mill capacity

1.5 L

Operation mode

Batch

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Temperature T

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Processing time t

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Journal Pre-proof Highlights 

A two-step method to obtain shape anisotropic CoO-GeO2-Li2O-B2O3 flake-like glass particles for Li-ion batteries.



Improvement of specific capacity, cycle life and rate capability due to shape and size control by wet phase compression. High capacity retention of >90% after 100 cycles.



Increase in specific capacity due to GeO2 as glass network former and improved

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cycling stability due borate glass network.

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Journal Pre-proof Declaration of interests

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☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

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Figure 1

Figure 2

Figure 3

Figure 4

Figure 5

Figure 6

Figure 7

Figure 8