# Impact of non-framework cation mixing on the structure and crystallization behavior of model high-level waste glasses

Nikhila Balasubramanya, Zhenxuan Sun, Mostafa Ahmadzadeh, Saeed Kamali, Daniel R. Neuville, John McCloy, Ashutosh Goel

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Nikhila Balasubramanya, Zhenxuan Sun, Mostafa Ahmadzadeh, Saeed Kamali, Daniel R. Neuville, et al.. Impact of non-framework cation mixing on the structure and crystallization behavior of model high-level waste glasses. Journal of the American Ceramic Society, 2022, 105 (6), pp.3967-3985. 10.1111/jace.18381. hal-04264015

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Impact of non-framework cation mixing on the structure and crystallization behavior of model high-level waste glasses

Nikhila Balasubramanya,<fnref n="1" /> Zhenxuan Sun,<fnref n="1" /> Mostafa Ahmadzadeh,<fnref n="2" /> Saeed

Kamali,<fnref n="3" />,<fnref n="4" /> Daniel Neuville,<fnref n="5" /> John S. McCloy,<fnref n="2" /> Ashutosh Goel<fnref n="1" />,<fnref n="1" />

<footnote n="1">¹ Rutgers, the State University of New Jersey, Piscataway, NJ 08854, USA.</footnote>

<footnote n="2">²School of Mechanical and Materials Engineering and Materials Science and Engineering Program, Washington State University, Pullman, WA 99164, USA.</footnote>

<footnote n="3">³Mechanical, Aerospace and Biomedical Engineering Department, University of Tennessee Space Institute, Tullahoma, TN 37388, USA.</footnote>

<footnote n="4">⁴Department of Physics and Astronomy, Middle Tennessee State University, Murfreesboro, TN 37132, USA</footnote>

<footnote n="5">⁵Institut de Physique du Globe de Paris, CNRS-Géomatériaux, Sorbonne Paris Cité, 1 rue Jussieu, 75005 Paris, France</footnote>

<footnote n="1">¹ Corresponding author
Email: ag1179@soe.rutgers.edu; Ph: +1 848-445-4512</footnote>

<footnote>This is the author manuscript accepted for publication and has undergone full peer review but has not been through the copyediting, typesetting, pagination and proofreading process, which may lead to differences between this version and the Version of Record. Please cite this article as doi: 10.1111/jace.18381.</footnote>

<footnote>This article is protected by copyright. All rights reserved.</footnote>
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Abstract

Spinel crystallization is known to be detrimental to the operation of Joule heated ceramic melters during the vitrification of iron-rich high-level nuclear wastes (HLW) into borosilicate glasses. The literature on this subject focuses on tackling the problem by developing empirical constraints to design compositions, which limit the fraction of spinels formed in the melter or by developing empirical models to predict the settling behavior of spinels in the melter as a function of the glass composition. While these empirical models can predict the behavior of most of the compositions, they are not a failsafe as there are always some compositions, whose behavior is beyond the predictive ability of these models. This can lead to undesirable situations during the vitrification of the nuclear waste, and therefore an in-depth investigation of the chemo-structural descriptors controlling the crystallization behavior in these glasses is warranted. Accordingly, the present study aims to understand the impact of non-framework cation mixing (i.e., Li⁺/Na⁺ and Ca²⁺/Na⁺) on the structure (through Raman spectroscopy and Mössbauer spectroscopy) and crystallization behavior (through XRD, SEM-EDS, and vibrating sample magnetometry) of iron-rich model HLW glasses in the system: (mol.%) x M_yO – (25-x) Na₂O – 9.12 B₂O₃ – 6.4 Al₂O₃ – 51.25 SiO₂ – 7.22 Fe₂O₃ – 0.38 MnO – 0.08 Cr₂O₃ – 0.55NiO (M_yO = Li₂O or CaO).

**Keywords:** nuclear waste; borosilicate; crystallization; structure; spinel
Introduction

The Hanford tank waste is a result of 45 years of plutonium production in support of the U.S. defense programs. The plutonium (²³⁹Pu) was generated by the transmutation of uranium through irradiation of the metallic fuel followed by its extraction using one of the following three primary separation processes: (i) bismuth phosphate process, (ii) REDOX process, and (iii) PUREX (plutonium uranium redox extraction) process. Depending on the process employed to separate ²³⁹Pu, as well as the detailed history of tank filling and blending, several waste streams with significant differences in their compositional chemistry were produced resulting in ~56 million gallons of nuclear waste stored in 177 underground steel tanks.¹ The waste is expected to be vitrified into borosilicate glasses wherein the waste feed will be mixed with the glass forming precursors (not glass frits) followed by melting in Joule Heated Ceramic Melters (JHCMs) at temperatures ≤1150 °C.¹

During the vitrification of the HLW, the non-framework cations, for example, Na⁺, Li⁺ and Ca²⁺, are known to exert a significant impact on the structure and processing of the nuclear waste melt, affecting the crystallization and chemical durability of the resulting glassy waste form. Here the framework comprises of the glass network-forming metals, primarily Si and B, and to an extent Al and Fe, which form the primary backbone or glass network as they bond with oxygen atoms. The problem pertaining to the crystallization of alkali aluminosilicate-based phases during the cooling of the melts (in the steel canisters) and its impact on their chemical durability has been widely reported.²⁻⁴ A similar problem is that of the impact of non-framework cations on the crystallization of (Fe, Ni, Mn, Zn)II(Fe, Cr, Mn)III₂O₄ mixed stoichiometry spinel crystals in the glass discharge riser of the melter during idling (where the melter temperature is lowered to 850–900 °C when not in operation).⁵⁻⁷

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Here it should be noted that the problem of spinel crystallization in nuclear waste melts has been well investigated and reported in the literature, but only at an empirical level, where the effort has been focused on developing constraints to limit the crystallization of spinels from borosilicate glass in the HLW melter.¹, ⁸⁻¹⁰ However, the validity of these empirical constraints is now being challenged.¹¹, ¹² Therefore, a transition from the empirical approaches towards understanding the fundamental science governing the crystallization of spinels in the glass melt is needed.

With respect to the impact of non-framework cations on the devitrification of iron-
rich spinel crystals in the glass melt, Li⁺, Na⁺, and Ca²⁺ constitute the major fraction of these
cations in the HLW. While a large fraction of Na⁺ in the HLW comes from the pre-
vitrification neutralization streams,¹ a small fraction of Ca²⁺ also makes its way through as an
impurity in the caustic NaOH neutralization streams.¹³ Ca²⁺ also enters into the wastes from
the rinse and flush water, and is added in relatively greater amounts for ⁹⁰Sr decontamination
of the wastes.¹³ Li⁺, which is present in insignificant amounts in the waste stream, is
intentionally added to decrease the viscosity and increase the electrical conductivity of the
melt.¹⁴, ¹⁵

It is difficult to investigate the impact of these ions on the structure of iron-rich
borosilicate glasses using conventional solid state nuclear magnetic resonance (s.s. NMR)
spectroscopy due to the paramagnetic behavior of iron. Nevertheless, there is enough
evidence in the literature demonstrating a significant impact of both Li⁺ and Ca²⁺ on the
structure and properties of iron-free sodium borosilicate, sodium aluminosilicate, and sodium
aluminoborosilicate glasses.⁴, ¹⁶⁻¹⁸ For example, the partial substitution of Na⁺ by Li⁺ or Ca²⁺
is known to increase the concentration of trigonal BO₃ units at the expense of tetrahedral
BO₄⁻, ¹⁶, ¹⁷, ¹⁹ and promote the formation of higher-coordinated aluminum species in the
sodium borosilicate and aluminoborosilicate glasses.¹⁶, ¹⁷ Similar behavior of these cations in
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iron-rich glasses will lead to a change in the availability of excess alkali/alkaline-earth cations (left over after charge compensating AlO₄) to charge compensate FeO₄⁻ units,²⁰ thus impacting the redox chemistry of iron and the crystallization behavior of spinels in the glass melt. Despite their significant presence in the nuclear waste melt,¹³⁻¹⁵, ²¹ and their known impact on the redox chemistry of iron,²²⁻²⁴ the role of these ions in altering the crystallization behavior of spinels in borosilicate-based HLW glasses has been largely ignored or not reported in the literature. Therefore, the present study aims at understanding the impact of Li₂O and CaO on the structure and crystallization behavior of model HLW glasses (based on the glass composition of the iron-rich waste present in the AZ-101 tank)²⁵ in the system (mol.%): x M_yO – (25-x) Na₂O – 9.12 B₂O₃ – 6.4 Al₂O₃ – 51.25 SiO₂ – 7.22 Fe₂O₃ – 0.38 MnO – 0.08 Cr₂O₃ – 0.55 NiO (M_yO = Li₂O or CaO), where x varies between 0 - 25 mol.% with Δx = 5 mol.%.

Experimental

Glass design and synthesis

A simplified nuclear waste glass with composition 25.00 Na₂O – 9.12 B₂O₃ – 6.40 Al₂O₃ – 51.25 SiO₂ – 7.22 Fe₂O₃ – 0.38 MnO – 0.08 Cr₂O₃ – 0.55 NiO (mol%), derived from the AZ-101 HLW glass composition,²⁵ has been chosen as the baseline for this work. To understand the impact of Li₂O and CaO on the structure and crystallization behavior of the HLW glasses, two series of glasses with varying Li₂O/Na₂O or CaO/Na₂O ratios in the baseline composition have been synthesized and investigated. Accordingly, x M_yO – (25.00-x) Na₂O – 9.12 B₂O₃ – 6.40 Al₂O₃ – 51.25 SiO₂ – 7.22 Fe₂O₃ – 0.38 MnO – 0.08 Cr₂O₃ – 0.55 NiO (mol.%) represents the batched compositions of the investigated glasses, where M_yO corresponds to Li₂O or CaO, and x varies between 0 to 25 mol.% with Δx = 5 mol.%.
varying CaO/Na₂O ratio will be referred as Ca-x, with x being the molar concentration of Li₂O or CaO in the respective glass.

The homogeneous mixtures of batches corresponding to 100 g oxides were prepared from high-purity powders of Na₂CO₃ (Fischer Scientific: >99.55%), Al₂O₃ (Acros organics, extra pure; 99%), H₃BO₃ (Acros organics; extra pure, 99+%), SiO₂ (Acros organics, 99.5%), Fe₂O₃ (Sigma Aldrich; ≥99%), NiO (Alfa Aesar, 99%), Cr₂O₃ (Fischer Scientific, 99%) and MnO (Fischer Scientific, 99%). The batches were melted in ambient atmosphere in 90% Pt-10% Rh crucibles (covered with a Pt lid) at 1450 °C for 1-2 h in an electric furnace (Carbolite BLF 1800, Carbolite-Gero Limited, Hope Valley, UK). The melts were splat quenched between two copper plates. The amorphous/crystalline nature of the splat-quenched samples was confirmed by powder X-ray diffraction (PANalytical – X’Pert Pro; Cu Kα radiation in the 2θ range of 10-90°, step size of 0.0262° s⁻¹ and a dwell time of 1.94 s per step).

## Raman Spectroscopy

Raman spectra were obtained to study the changes in the structure of as-quenched glasses as a function of their composition. The spectra were measured using a Renishaw inVia Reflex spectrometer in the back-scattering geometry. The flat surface of a glass sample was excited using a solid-state laser emitting at the wavelength of 532 nm. The laser spot was focused onto the sample using an objective of 20× and the backscattered light was analyzed with 1800 grooves/mm diffraction grating. Raman spectra could be reproduced with an accuracy of ±1 cm⁻¹. The power impinging the surface was about 10 mW. Ten accumulations were taken for each scan with an acquisition time of 10 s.

## Differential scanning calorimetry

The glass grains of particle size < 45 µm were used to collect differential scanning calorimetry (DSC) data using a Simultaneous Thermal Analyzer (DSC; NETZSCH Pegasus 404 C) in the temperature range 20-1400 °C at a heating rate 10 °C min⁻¹ under a constant

dip
flow of nitrogen gas in Pt-pans with the sample being referenced to an empty Pt-pan. The temperatures, corresponding to onset of glass transition (T<sub>g</sub>), peak of crystallization (T<sub>c</sub>), and melting (T<sub>m</sub>), were obtained from DSC scans. The DSC data reported for any glass composition are an average of three thermal scans.

## Mössbauer Spectroscopy

Mössbauer spectroscopy was performed to understand the impact of glass composition on the redox chemistry of iron in the as-quenched glasses. Accordingly, the Mössbauer spectra of the as-quenched samples was collected at 20 °C using a constant acceleration spectrometer with a 25 mCi <sup>50</sup>Co source in a Rh matrix. Absorbers were prepared from finely ground samples that were mixed with BN powder to ensure a Mössbauer thickness, t < 1.<fnref n="2" /> The spectra were measured in the velocity range of ±12 mm/s relative to α-Fe and were fitted using the Recoil software package.²⁶ For the samples analyzed, broadened Lorentzian doublets were fitted. While fitting the spectra, it was assumed that the recoil-free fraction ratio f(Fe<sup>3+</sup>)/f(Fe<sup>2+</sup>) = 1.

## Isothermal heat treatment at 950 °C

To understand the crystalline phase evolution in the glass melts upon cooling, the glass frits were melted in 90% Pt-10% Rh crucibles at 1450 °C followed by cooling to 950 °C at 10 °C/s. The glass melts were then isothermally heated at 950 °C for 48 h followed by cooling to room temperature. The heat treatment regime has been primarily chosen considering the processing (idling) constraint that will be faced by the JHCM with respect to spinel formation in glasses, i.e., at a temperature of ≤950 °C, the spinel content in glasses should be ~1 vol.%.

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<footnote n="2">² Mössbauer thickness is the effective thickness of a source ($t_s$) or absorber ($t_a$). For an experiment involving element Z, Mössbauer thickness ($t$) is given by $t = n\sigma af$, where $n$ is the number of atoms of Z per cm², $\sigma$ is the cross-section for resonance absorption, $a$ is the fractional abundance of the Z nuclide which shows the Mössbauer effect being observed, and $f$ is the recoil-free fraction.</footnote>
## X-ray diffraction (XRD) and Rietveld analysis

The qualitative phase analysis of the heat-treated samples followed by their quantitative analysis for determining the concentration of crystalline and amorphous content has been performed by XRD (PANalytical – X’Pert Pro). The samples were crushed to powder (particle size < 45 µm) in an agate pestle and mortar and mixed with 10 wt.% Al₂O₃ (NIST SRM 676A) as an internal standard for Rietveld – reference intensity ratio (RIR) analysis. XRD data were recorded with Cu Kα radiation in the 2θ range of 10-90°, dwell time of 1.94 s per step and step size of 0.0071° s⁻¹. The quantitative crystalline phase analysis was performed using a Jade 6.0 software (MDI, Materials Data Incorporated, Livermore, CA, USA). The crystalline phase fractions were rescaled based on the absolute mass of Al₂O₃ (internal standard) originally added to the samples. The background was fitted with a Chebyshev function with a variable number of coefficients depending on its complexity. The peak profiles were modeled using a pseudo-Voigt function. The lattice constants, phase fractions and coefficients corresponding to sample displacement and asymmetry were refined.

## Vibrating samples magnetometry (VSM) measurements

While the XRD was used as a primary tool to confirm the amorphous/crystalline nature of the quenched and isothermally heat-treated samples, the VSM measurements were performed on powders to detect the presence of any magnetic phases that remained undetected by the XRD. The magnetic measurements were performed using a Lakeshore PMC3900 VSM with a maximum applied field of 1.8 T. The sample's magnetization (M) dependency on the magnetic applied field (H) (i.e., magnetic hysteresis loops) were collected with a maximum applied field of ±1.8 T A/m with a field increment of 2.0 T A/m.

## Scanning electron microscopy and energy dispersive spectroscopy

The isothermally heat-treated samples were polished (using SiC grinding sheets followed by 3 µm diamond suspension) followed by ultrasonic cleaning with deionized

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water. The samples were mounted on metallic stubs and sputter-coated with a ~20 nm thick
gold layer. The microstructure of the heat-treated (and chemically un-etched) samples was
studied in the secondary electron imaging mode using a field-emission secondary electron
microscope (FE-SEM; ZEISS Sigma). The elemental analysis of the observed microstructure
was performed using energy dispersive spectroscopy (EDS; X-Max Oxford Instruments;
Aztec software).

Results

Physical appearance and amorphous nature of the melt-quenched samples

The as-synthesized glasses are dark brown to black in color due to the presence of
transition metal ions, while the XRD analysis depicts their amorphous nature, as shown in
Figure S1. The VSM measurements have been performed to corroborate the XRD results,
wherein the aim is to detect the presence of magnetic crystals, if any, that may have escaped
detection in XRD due to their smaller crystallite size or phase fraction. Accordingly, Figure 1
presents the magnetizations versus magnetic field curves for all the glasses synthesized in the
study. The insets in Figure 1 present the zoomed-in hysteresis loop of Li-0 (Figure 1a) and
Ca-0 (Figure 1b) samples, respectively. The linear behavior observed in the VSM curves is
from the paramagnetic behavior (linear dependency of magnetization on applied magnetic
field) of Fe (also Mn, Cr and Ni whose mol% are << Fe) in the glasses,²⁷ while the weak
hysteresis observed is likely from the clustering of Fe species (also Mn, Cr and Ni species
whose mol% are << Fe).²⁸,²⁹ It is noteworthy that the magnetization observed on the y-axis
due to the possible clustering is similar in the glasses of both the Ca-x series and Li-x series.
According to Bingham *et al.*,³⁰,³¹ the presence of <4 mol.% Fe₂O₃ in alkali/alkaline-earth
silicates results in a non-flat background in the ⁵⁷Fe Mössbauer spectra, a typical
characteristic of hyperfine paramagnetic splitting suggesting the presence of isolated Fe³+
species in the glass structure. On the other hand, the Mössbauer spectra of glasses containing

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>4 mol.% Fe₂O₃ presents a flat background indicating the presence of exchange interactions suggesting the formation of Fe³⁺-O-Fe³⁺ and Fe³⁺-O-Fe²⁺ in the glass structure (clustering). The results in the present study (on glasses containing ~9 mol.% Fe₂O₃) are consistent with the latter literature, where we observe a similar flat background in the Mössbauer spectra of the investigated glasses (Figure S2 (a and b)) suggesting clustering of iron atoms leading to the observed magnetization.

## Compositional dependence of iron redox chemistry in glasses

Figure 2 presents the compositional dependence of Fe³⁺/ΣFe content in the investigated glasses, as measured using Mössbauer spectroscopy. The Mössbauer spectrum of Li-0 glass has been presented in the inset of Figure 2 where the spectrum has been best fitted with three sub-spectra, S₁, S₂, and S₃. The Mössbauer spectra of the remaining samples with the fitted sub-spectra have been presented in the Figure S2 (a and b). Table 1 presents the hyperfine parameters, i.e., centroid shift (δ) with respect to metallic α-iron at room temperature, quadrupole splitting (ΔE_Q), Lorentzian linewidth (Γ) and intensities (I), extracted for each sub-spectra using the least square fit, and Fe³⁺/ΣFe ratio in the glasses.

As evident from Figure 2, the Fe³⁺/ΣFe ratio decreases with increasing Li₂O and CaO content in the glasses. This behavior may be explained based on either decreasing optical basicity (OB) of the glass melts with increasing Li⁺/Na⁺ or Ca²⁺/Na⁺,³²,³³ as shown in Figure 2, or lower availability of Na⁺ to charge compensate tetrahedral FeO₄⁻ units in the glass structure, as discussed later in the article (see Discussion). Although the Fe³⁺/ΣFe ratio follows the OB trend closely, there is a discernible deviation from the predicted OB vs. Fe³⁺/ΣFe trend for the glass Ca-25, as evident from a significant decrease in the Fe³⁺/ΣFe ratio in this glass. While the mean δ values of sub-spectra S₁ and S₂ in both the series of glasses are typical for Fe³⁺ in tetrahedral coordination, an increase in these values with increasing Li₂O and CaO content suggests Fe³⁺ moving towards higher coordination, but still

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not having reached the octahedral Fe³⁺ mean δ values of 0.35-0.55 mm s⁻¹.³⁴, ³⁵ Here it needs to be emphasized that though Fe³⁺ is primarily known to exist in tetrahedral coordination, thus, acting as a network former (FeO₄⁻) being charge compensated by an alkali or an alkaline earth cations,²⁰, ³⁵, ³⁶ five- and six-coordinated Fe³⁺ have also been reported to exist in glasses comprising non-framework cations with high ionic field strength.³³, ³⁷

Further, the mean δ values of sub-spectrum S₃ (in Figure 2 and Figure S2) are typical for Fe²⁺ in the octahedral coordination. A minimal change in the δ of sub-spectra S₃ with increasing Li₂O and CaO content imply a constant Fe²⁺ coordination environment in the investigated glasses.³⁴, ³⁸ According to the literature, Fe²⁺ primarily exists in 5-6 coordination (rarely in 4-coordination), thus acting as a network modifier in silicate and borosilicate glass melts.³⁴, ³⁵

While δ in Mössbauer spectroscopy is an indicator of the oxidation state and the coordination state of the Fe atoms via measurements of the s-electron density at the ⁵⁷Fe nucleus, ΔE_Q is an indication of the symmetry around Fe atoms. It can be seen from Table 1 that the ΔE_Q increases with an increase in the concentration of Li⁺ and Ca²⁺ in the glasses. ΔE_Q reflects the interaction between the quadrupole nuclear moment with surrounding electric field gradient (EFG) leading to splitting of the energy levels. When ΔE_Q is observed, an asymmetrical electric field is produced by an asymmetric electronic charge distribution or ligand arrangement. Hence, in the present study, an increase in ΔE_Q can be attributed to the higher electronegativity of Li⁺ and Ca²⁺ than Na⁺, thus polarizing and inducing greater asymmetry of the ligand oxygen ions around the Fe³⁺ and Fe²⁺.¹⁷

Raman spectroscopy

As deciphered from the results of Mössbauer spectroscopy, the majority of Fe³⁺ exists in tetrahedral coordination in the investigated glasses, though its concentration (Fe³⁺/ΣFe) decreases with increasing Li⁺/Na⁺ or Ca²⁺/Na⁺ ratios. In this section, the aim is to understand
the impact of changes in the redox chemistry of iron (as a function of glass composition) on the short-to-medium range ordering in the structure of the investigated glasses. To this effect, Figure 3 presents the Raman spectra of Li-x (Figure 3a) and Ca-x (Figure 3b) series of glasses. The broad features of the spectra indicate the amorphous nature of the samples being analyzed, thus, corroborating the XRD results (Figure S1).

For the convenience of the readers, the Raman spectra of the investigated glasses has been divided into the following three regions: Region I (250 cm⁻¹ to 800 cm⁻¹; low-frequency), Region II (800 cm⁻¹ to 1200 cm⁻¹; mid-frequency), and Region III (1200 cm⁻¹ to 1600 cm⁻¹; high-frequency). The low frequency region is characterized by a broad band centered at ~500 cm⁻¹ corresponding to the stretching and bending modes of T–O–T units (T = Si, Al).³⁹,⁴⁰ While the band in the 650 cm⁻¹-800 cm⁻¹ in the low-frequency region can be due to the symmetric breathing vibration of triborate, tetraborate, or pentaborate groups,³⁹ the band centered at ~775 cm⁻¹ may also be assigned to the Al–O stretching.⁴⁰ The band centered at ~1350 cm⁻¹ in the high-frequency region has been assigned to the stretching of B–O⁻ bonds attached to large borate groups.³⁹,⁴⁰ This band can also be assigned to the stretching modes of BO₃ units.

The broad intense band between 800 and 1200 cm⁻¹, in the mid-frequency region, is composed of several bands correlated to silicates,⁴¹,⁴² which also include the vibrations of the Fe³⁺–O stretching modes at 980 cm⁻¹.⁴³,⁴⁴ Figure S3 (a and b) presents the normalized and baseline corrected Raman spectra of the glasses in the mid-frequency region. Further, to understand the contribution of different *Q''* units (where, *Q* is the degree of polymerization, and *n* is the number of bridging oxygens), with the help of literature,⁴¹,⁴² this band has been deconvoluted. The details about the methodology used for the deconvolution of Raman spectra have been presented in the supplementary information.

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Based on the previous studies describing the structure of alkali silicate glasses,⁴¹, ⁴² the small band at ~890 cm⁻¹ is attributed to Si–O⁻ stretching, with three non-bridging oxygens (NBOs) per silicon (Q¹). The band at ~930 cm⁻¹ is assigned to the Si–O⁻ stretching in a structural unit with two NBOs per silicon (Q²) and the bands centered between 1050 and 1150 cm⁻¹ are assigned to Si–O⁻ stretching in a structural unit with one NBO per silicon (Q<fnref n="3" />). Further, according to the literature,⁴⁵⁻⁴⁷ the peak at ~1030 cm⁻¹ belongs to the T₂<fnref n="3" /> vibrational modes of the different TO₄ tetrahedra (T = Al, Si). The peak at ~1080 cm⁻¹ has been assigned to Q<fnref n="3" /> units with one NBO. It is noteworthy that the literature ascribes a band at ~1200 cm⁻¹ to the A₁<fnref n="4" /> symmetric stretching of Q<fnref n="4" /> species.⁴⁷, ⁴⁸ However, we did not observe the presence of this band in the Raman spectra of the investigated glass compositions, thus, suggesting the absence of Q<fnref n="4" /> species in the glass network.

The structural speciation in the glasses has been obtained by calculating the ratios of the areas under the fitted bands in the Raman spectra in the region 800 – 1200 cm⁻¹. Table 2 presents the area ratios, while Figure 4 (a and b) shows the variation of the Qⁿ species with the change in the composition. From the area ratios, Qⁿ speciation has been calculated with an assumption that Q<sub>tot</sub> = Q¹+Q²+Q<fnref n="3" />, i.e., Q<fnref n="4" /> = 0 (as discussed above). Further, to deduce the direction in which the 2Q² ⇌ Q<fnref n="3" /> + Q¹ equilibrium shifts, the ratio Q²/(Q<fnref n="3" />+Q¹) has been calculated and plotted in Figure 5 (a and b) along with the Fe<fnref n="3" />+/Q<sub>tot</sub> ratio, which has been calculated for each composition in the two series of glasses. As evident from Figure 5, an increase in Ca²+/Na⁺ and Li⁺/Na⁺ ratios in the glasses result in depolymerization in the silicate network, thus, resulting in higher fraction of Q² units, thus, shifting the 2Q² ⇌ Q<fnref n="3" /> + Q¹ equilibrium towards left. However, the degree of depolymerization is higher in the CaO-

<footnote n="3">³ The T₂ vibrational mode (asymmetric stretching mode) corresponds to the mode, where two oxygen atoms move closer to the central Si atom, while the other two oxygen atoms move away.⁴⁷</footnote>

<footnote n="4">⁴ This vibrational mode (symmetric stretching mode) corresponds to an in-phase motion of the four oxygen atoms towards the Si atom at the centre.⁴⁷,⁴⁸</footnote>
containing glasses in comparison to Li₂O-containing glasses owing to the higher ionic field strength of Ca²⁺ over Li⁺ or Na⁺ (please refer to Discussion for further details).¹⁶, ⁴⁹

## Compositional and structural dependence of glass transition

Figure 6 presents the trends observed for variation in the glass transition temperature (T_g) of the investigated glasses as a function of their composition. The T_g of the glasses have been obtained from the onset of the endothermic dip in their respective DSC scans and presented in Figure S4. The initial substitution (x = 5 mol.%) of Li₂O for Na₂O in the Li-x glasses results in a sharp decrease in the T_g value from 488 °C to 444 °C, while further increase in Li₂O content from 5 mol.% to 15 mol.% results in a subtle decrease in the T_g of these glasses with the glass Li-15 exhibiting a minimum at 428 °C. Further increase in Li₂O content from 15 mol.% to 25 mol.% resulted in a gradual increase in the T_g from 428 °C to 452 °C, thus, resulting in a U- or V-shaped curve, typical for glasses manifesting mixed alkali effect (MAE).⁵⁰⁻⁵²

In the glasses in Ca-x series, the T_g increases almost linearly, from 488 °C to 615 °C, with an increase in their CaO content. The high ionic field strength cations, for example, Mg²⁺, Ca²⁺, and rare-earths (RE³⁺) – hereafter referred as Mⁿ⁺ – when substituted for Na⁺ in the glass structure are known to (1) create highly charged NBOs in silicate and/or borate regions (depending on the glass chemistry) resulting in the formation of Mⁿ⁺–NBO enriched regions;¹⁶, ⁵³ (2) facilitate the conversion of BO₄ to BO₃ by promoting the concentration of negative charges on NBOs in their local coordination environment;¹⁶, ¹⁷, ¹⁹ and (3) promote the formation of AlO₅ units.¹⁶, ⁵⁴ Thus, the increase in T_g of glasses with increasing Ca²⁺/Na⁺ may be attributed to the increasing rigidity of the glass network owing to the formation of (1) stronger Ca²⁺–NBO bonds compared to weaker Na⁺–NBO bonds, and (2) highly coordinated aluminum species in the glass structure. Similar trends with respect to an increase in T_g with

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increasing concentration of $M^{n+}$-NBO regions (in glasses comprising high field strength cations) and five-coordinated aluminum have been widely reported in the literature.⁵⁵⁻⁵⁹

## Crystalline phase evolution as a function of glass composition

XRD was performed on the samples cooled down from the melt followed by heat treatment at 950 °C (for 48 h) to follow their crystalline phase evolution both qualitatively and quantitatively. The XRD data have been corroborated and complemented by the VSM and SEM-EDS results to ascertain the presence of any undetected (in XRD) phases (amorphous/crystalline) in the heat-treated samples.

Figure 7 (a and b) shows the diffraction patterns arising from the heat-treated glass samples investigated in the study. The samples with $Li_2O$ (and $CaO$) ≤ 5 mol.% remain amorphous after the heat-treatment. Further increase in $Li_2O$ (and $CaO$; ≥10 mol.%) results in crystallization of spinel phases with general formula $AB_2O_4$ (space group Fd3m), where the ‘A’ site can be occupied by an $Fe^{2+}$ or $Ni^{2+}$ or $Mn^{2+}$ or $Li^+$ ion, and ‘B’ site can be occupied by $Fe^{3+}$ or $Cr^{3+}$ or $Mn^{3+}$. In the presence of $Mn^{4+}$, the spinel of the form $A_2BO_4$ can form where $Fe^{2+}$ or $Ni^{2+}$ occupies the A sites and the ‘B’ site is occupied by $Mn^{4+}$. However, since the glasses are rich in $Fe_2O_3$ and contain <1 mol.% $NiO$, $Cr_2O_3$ and $MnO$, the spinel phase most likely is close to $Fe_3O_4$ (magnetite) in stoichiometry. The crystallization of spinels is accompanied by the precipitation of hematite ($\alpha-Fe_2O_3$) in the glass Li-25 and all the glasses with $CaO \ge 10$ mol.%. Quantitatively, the amount of crystallinity in all the glass (except Li-25 where the major phase is spodumene, $LiAlSi_2O_6$) is ≤ 6 wt.%, wherein the concentration of spinels varies from 1 wt.% - 5 wt.% and increases with increase in $Li_2O$ or $CaO$ concentrations Figure 8 (a and b).

When comparing the impact of $Li_2O$ and $CaO$ on spinel formation in glasses, $Li_2O$ seems to slightly favor the crystallization of spinels more than $CaO$, as the concentration of spinels increases from 1.4 wt.% to 4.8 wt.% with an increase in $Li_2O$ content (10 - 25 mol.%)

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in the glasses, while it increases from 1.0 wt.% to 3.8 wt.% with increasing CaO content in
glasses on equimolar basis Figure 8 (a and b). These results are in good agreement with
Mössbauer spectroscopy results, as shown in Figure 2, where substitution of Na₂O by Li₂O
has been shown to exhibit higher impact on the OB and Fe³⁺/∑Fe concentration in the
investigated glasses compared to CaO.

The SEM-EDS results of the resultant glass-ceramics corroborate the XRD results by depicting the dendritic morphology of spinel crystals rich in Fe, Ni, Cr and Mn in samples Li-25 (Figure 9a) and Ca-25 (Figure 9b). Further, Figure 9a shows the presence of Al and Si rich crystals most likely confirming the crystallization of spodumene phase in the sample Li-25 (Li remains undetected in the EDS). Based on the shape and morphology of the crystals, as deduced by SEM-EDS, the spinel and/or hematite crystals observed in Li-25 are in the size range of ~ 5-15 µm. In Ca-25, however, the spinel and/or hematite crystals are much smaller in size as can be seen in Figure 9b.

Figure 9c presents the formation Ca-rich regions in the glass Ca-25. These Ca-rich regions were observed all over the surface of the glass ceramic. The Ca-rich regions are in agreement with HLW studies, where anorthite has been reported to precipitate out of the Ca-rich aluminosilicate regions observed under transmission electron microscopy (TEM)⁶⁰ as well as the other Ca rich oxides, like minerals of the types pyroxene and olivine, observed in other studies.⁹, ⁶¹

Figure 10 shows the raw (dashed curve) and slope-corrected (solid curve) hysteresis
loops of both Li and Ca-series glasses heat-treated at 950 °C. The raw hysteresis loops of the
heat-treated glasses (dashed curves) exhibit that most of the samples in the current study are
not fully saturated even at the highest applied field of 1.8 T. The constant slope of the loops
in the raw data at higher magnetic fields is due to paramagnetism arising from the Fe (along
with Cr, Mn and Ni) ions that remain in the residual glassy phase. To correct for the

dipolar
paramagnetic contribution, the slope-corrected hysteresis loops (solid curves) have been obtained by subtracting the slopes at 70% of the maximum applied field. The saturation magnetization ($M_s$), denoted by the point at which the applied magnetic field saturates, represents the magnetic phase content, which arises from the $Fe^{3+}$, $Fe^{2+}$, $Cr^{3+}$, $Mn^{3+}$, $Mn^{4+}$, and $Ni^{2+}$ species entering the crystalline phase field, forming spinel and hematite phases observed in Figure 7 (a and b) and Figure 8 (a and b). The $M_s$ is an absolute parameter that can be obtained from the point where the slope corrected data (solid line) meets the y-axis. As evident from Figure 10, the $M_s$ increases with increase in $Li_2O$ and $CaO$ content suggesting greater formation of spinel and hematite in the samples.

Amongst the phases present, magnetite ($Fe_3O_4$), a ferrimagnetic phase, is represented by high magnetization and low coercivity, whereas antiferromagnetic hematite has small magnetization and high coercivity. It should be noted here that since $Fe_2O_3$ (mol %) >> $Cr_2O_3$ (mol %) or $NiO$ (mol %) or $MnO$ (mol %), although solid solution phases of spinel and hematite are present in the investigated samples, an approximation of the phases to pure spinel and hematite has been made to simply understand the magnetic behavior of such species. Magnetite and hematite, as the most commonly formed iron oxides, have characteristic room temperature saturation magnetizations ($M_s$) of ~92 Am$^2$/kg and ~0.4 Am$^2$/kg respectively.²⁷, ⁶²⁻⁶⁴ Therefore, high $M_s$ values are due to magnetite (ferrimagnetic), while hematite (canted antiferromagnetic) can be responsible for coercivities. However, it has been reported that, when both magnetite and hematite are present, only ~5 wt.% magnetite can mute the coercivity of ~95 wt.% hematite in the magnetic hysteresis loop of the mixture.⁶⁴ This can explain why the samples show relatively low coercivities even when there is significant amount of hematite (i.e., $Ca-x$ series). Further, this also explains the higher $M_s$ observed in the $Li-x$ series compared to that of $Ca-x$ series as due to the small but relatively higher amount of spinel formed in $Li-x$ series as compared to that of $Ca-x$ series.

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It is well known that there exists an empirical relationship between OB and redox
chemistry of iron in oxide glasses wherein increasing basicity favors the upper oxidation state
in Fe²⁺–Fe³⁺ redox couple.²⁰,⁶⁵ When the OB is high, while the O²⁻ ions surrounded by the
non-polarizing cations such as Na⁺, Ca²⁺ etc., are practically uninfluenced by them, and hence
have the highest ability to donate the negative charge, the covalently bonded (polarized)
bridging O²⁻ ions have a decreased ability to donate charges as they are bonded to either two
Si⁴⁺ or one Si⁴⁺ ions. However, in comparison to a bridging oxygen, the non-bridging
oxygens can have greater ‘electron donor power’ as it forms the terminal end of a silicate
network.³³,⁶⁵,⁶⁶ With the presence of a metal cation capable of existing in multiple oxidation
states, for example, +2 and +3 for iron in the present investigation, the basicity, which is a
function of the number of free oxygen ions (non-polarized oxygen ions), bridging oxygens
and the non-bridging oxygens decide the oxidation state of the metal cation; in that, the
higher the basicity, the higher oxidation state is preferred. However, it should be noted here
that though generally acceptable, this relationship may not be true in all cases. For example,
as has been shown by Schreiber et al.,⁶⁷ in a series of sodium silicate glasses containing 1
wt.% Fe₂O₃, the Fe³⁺–Fe²⁺ couple becomes more reduced as the composition becomes more
basic with increasing Na/Si ratio up to about unity (Na/Si = 1) and after that point the redox
couple becomes more oxidized as the composition becomes even more basic. This implies
that the actual dependence of individual redox couples should not be generalized, as it is a
function of the availability of solvation sites for the redox states of the multivalent element in
the melt structure.

From the structural viewpoint, when comparing the role of Li⁺, Na⁺ and Ca²⁺ in the
glass structure, the Na⁺ will preferentially act as a charge compensator for AlO₄⁻, BO₄⁻, and
FeO₄⁻ ions, while the majority of Li⁺ and Ca²⁺ (depending on their concentration in the glass),

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owing to their high ionic field strength, are expected to promote heterogeneity in terms of Si/B mixing, and create non-bridging oxygens (NBOs) in the borosilicate glass network.<sup>16, 17, 68, 69</sup> The formation of NBOs is evidenced by an increase in the forward rate reaction Q<sup>3</sup> + Q<sup>1</sup> &harr; 2Q<sup>2</sup>, which takes place in the presence of Li<sup>+</sup> and Ca<sup>2+</sup>, as discussed in Section 3.3. Based on the literature,<sup>16, 17</sup> and the results from our previous studies,<sup>19, 20</sup> we expect Al<sup>3+</sup>, in the present investigation, to be tetrahedrally coordinated being preferentially charge compensated by Na<sup>+</sup> (at least in glasses where Na<sub>2</sub>O/Al<sub>2</sub>O<sub>3</sub> &ge; 1), while the boron coordination is expected to exhibit a significant decrease from four – to – three, i.e., BO<sub>4</sub><sup>-</sup> &rarr; BO<sub>3</sub> as the concentration of CaO or Li<sub>2</sub>O increases. According to the literature,<sup>16, 17</sup> higher field strength cations (for example, Li<sup>+</sup> and Ca<sup>2+</sup>) favor the formation of highly charged NBOs over the lower charged bridging oxygens that form linkages such as <sup>[4]</sup>B–O–<sup>[4]</sup>Si, thus shifting the equation 1 to the left. The resulting greater concentration of negative charge thus helps stabilize the local coordination environment of the non-framework cation with higher ionic field strength.

$$BO_{3/2} + SiO_{3/2}O^{-} \rightleftharpoons [BO_{4/2}]^{-} + SiO_{4/2} \quad (1)$$

Since FeO<sub>4</sub><sup>-</sup> is known to compete with BO<sub>4</sub><sup>-</sup> for the charge compensation by a non-framework metal cation,<sup>20, 36</sup> the decrease in Na<sup>+</sup> concentration in glasses is expected to lower the concentration of tetrahedral Fe<sup>3+</sup> to octahedral Fe<sup>2+</sup> in the investigated glasses.

## T<sub>g</sub> dependence on mixed alkali effect in Li-Na aluminoborosilicate glasses

The MAE in alkali silicate glasses is a well-known phenomenon.<sup>51, 52, 70-72</sup> However, its structural origins are still debatable, as several models have been proposed to explain the origins of this effect.<sup>50, 73-79</sup> The works of LaCourse,<sup>76</sup> Ingram<sup>74</sup> and Bunde et al.<sup>80</sup> in this direction laid the foundation to the development of the dynamic structural models and form the basis of the effect as agreed upon in literature concerning theoretical models for MAE. As recently as 2015, the thermometer effect, which has been known to be due to MAE was explained by Yu *et al.* on the basis of local instabilities in the structure leading to
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compressive stress exerted by the glass network on the alkali contained within.<sup>50, 79</sup> All these models assume that the ion motion is site-specific, rather than path specific. Simply put, cations occupying specific sites, and another cation “hopping” to occupy the sites with similar coordination atmosphere is believed to be the basis for the MAE. Improvising on the modified random network (MRN) model,<sup>81</sup> Greaves and Ngai<sup>78</sup> via X-ray absorption fine structure in mixed-alkali silicate glasses proposed that the alkali ions tended towards local clustering, although these clusters were distributed homogeneously in the glasses. Contradicting this are the works of Lee et al.,<sup>82-84</sup> and Angeli et al.,<sup>85</sup> using <sup>23</sup>Na and <sup>17</sup>O NMR which support the perturbed cation distribution model which assumes a relatively homogenous distribution of metal cations in the glass network. However, confirming the studies of Greaves and Ngai<sup>78</sup> were the works of Swenson et al.<sup>75</sup> and Swenson and Adams,<sup>70</sup> which rested that the alkali ions mixed randomly. In these models, the alkali modifiers are proposed to diffuse through channels via the cooperative intrachannel hopping mechanism. Le Losq and Neuville attribute the random mixing (resulting in segregation/clustering) of the two alkali cations in the glass structure as a probable cause for the non-linear variation of T<sub>g</sub> in mixed alkali aluminosilicate glasses.

The explanation of MAE becomes more complex in glasses containing multiple network formers, or when the ionic field strength of alkali cations exerts its influence on the glass structure. In the framework of Adam and Gibbs theory of viscous flow, the minima in T<sub>g</sub> observed upon mixing Na<sup>+</sup> and K<sup>+</sup> in silicate melts can be explained by an excess of configurational entropy at T<sub>g</sub>, which originates from a random mixing of pairs of alkali cations.<sup>51, 52</sup> Higher the configurational entropy, lower is the viscosity and higher the ionic mobility. This interpretation supports the cooperative inter-diffusion and mixing of alkali cations in ionic channels percolating through a disrupted silicate network, as discussed in modified random network model (MRN).<sup>78, 81</sup> The high-configurational entropy results
mainly from the differences in the ionic field strength of the two alkali cations; owing to its
high ionic field strength, Li⁺ is expected to introduce heterogeneity in the glass structure. This
has been well supported by the Raman studies on the Na and Li silicate glasses by Matson et
al.,⁸⁶ where silica-rich regions were observed in the presence of Li⁺. Clustering of Li⁺ is also
supported by the ²⁹Si(⁷Li) rotational double echo resonance spectroscopy results from Voigt
et al.⁸⁷ Le Losq and Neuville attribute the random mixing (resulting in segregation/clustering)
of the two alkali cations in the glass structure as a probable cause for the non-linear variation
of T_g in mixed alkali aluminosilicate glasses. It is to be noted here that the ‘random
distribution/randomness’ of the alkali ions is not equivalent to homogeneity, while on the
contrary results in the heterogeneity in the structure.⁷², ⁸⁸ While several studies have assayed
to scrupulously identify the structural causes at the atomic level that leads to high
configurational entropies in mixed alkali silicate glasses,⁵⁰, ⁷⁹, ⁸⁹ the glasses in the present
investigation comprise Al₂O₃, B₂O₃ and Fe₂O₃ along with SiO₂. This adds another layer of
compositional and structural complexity to the glass system, as now one must account for the
impact of ionic field strength of alkali cations on speciation of iron, borate and aluminate
units in the glass structure to identify the structural causes for the MAE. Due to these
limitations, we rest that the minima in T_g arises due to high-configurational entropy resulting
from the random mixing of Na-Li ions and the investigation on structural changes leading to
such high configurational entropy will be the focus of our future studies.

Effect of alkali/alkaline earth cations on the crystallization behavior of glasses

Borosilicate glasses and melts possess short-range order (SRO) with a radius of
influence of ~1.6–3 Å around the central atom and a medium range order (MRO) with the
radius of influence of ~3-6 Å.⁹ In clusters, where the regions are highly ordered, there is an
increase in coordination number and a decrease in the average atomic distance around the
central atom. However, the ordering does not reach the long-range extent as in a crystal. As

dipolar
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evident from Table 1, Fe³⁺ and Fe²⁺ exist in tetrahedral and octahedral coordination, respectively, in the investigated glasses. However, with the addition of high field strength cations, Li⁺ and Ca²⁺, the average coordination of – and distortion around – the respective central atoms (Fe³⁺ and Fe²⁺) increases, as evident from an increase in the Mössbauer δ and ΔE_Q values. These results from Mössbauer spectroscopy point towards a change in the electric field gradient around the Fe ions suggesting the evolution in the SRO in the structure of glasses with the addition of a cation of higher ionic field strength. The possibility of clustering around Fe²⁺ and Fe³⁺ in the investigated glasses is also corroborated by the VSM results (Figure 1 vs. Figure 10). The iron clusters, upon heat-treatment, possibly act as precursors to the crystallization of magnetite and spinel phases in the investigated samples.

Furthermore, the formation of hematite and spinel crystals in the aluminosilicate glasses are known to act as nucleation site for the volume crystallization of pyroxene phases.⁹⁰, ⁹¹ Accordingly, in the present study, spodumene (LiAlSi₂O₆), a clinopyroxene mineral, crystallizes (in volume) as the dominant phase in the sample Li-25. However, it is interesting (but not surprising) that the end-member clinopyroxene minerals corresponding to spodumene, i.e., jadeite (NaAlSi₂O₆) and Ca-Tschermak (CaAl₂SiO₆), did not crystallize from samples Li-0 (or Ca-0) and Ca-25, respectively. This may be attributed to the sluggish crystallization tendency of both the end-member clinopyroxene minerals, as has been discussed in the literature.⁹²

## Conclusions

The various experimental techniques aimed at understanding the compositional and structural drivers resulting in the spinel crystallization in aluminoborosilicate glasses with significant amounts of alkali/alkaline earth ions suggest an increase in spinel crystallization tendency in the presence of a higher field strength cation. In the presence of Li⁺ and Na⁺, it
has been observed in the glass structure that (1) The optical basicity decreases and Fe²⁺/ΣFe increases, (2) The NBOs in the structure increase, and (3) The rigidity of the structure increases as a result of increase in BO₃ and higher coordinated Al units. The extent of depolymerization of the glass structure observed was greater in the presence of Ca²⁺ than in the presence of Li⁺. The resulting changes brought about by the higher field strength cation and the clustering of iron at high iron contents results in the precipitation of spinel and hematite in the glass ceramics. In the glass ceramics it was also seen that the Al units in the presence of Li⁺ interacted closely, resulting in the spodumene in Li-25, while Ca²⁺ at its highest concentration led to Ca-rich regions in Ca-25. In other words, conversely it can be said that Na⁺ improves the glass forming ability and suppresses stabilization of other complexes, thus reducing the tendency of the matrix to phase separate upon heat treatment.

Acknowledgement:

This research is being performed using funding received from the DOE Offices of Nuclear Energy and Environmental Management through the Nuclear Energy University Programs under Grant No. DE-NE0008597.

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<figure><img src="image_3.png" /><figcaption>Figure Captions</figcaption></figure>

**Figure 1 (a and b).** Magnetization versus applied field curves obtained for the as-quenched samples of (a) Li-series and (b) Ca-series. The inset presents the magnified hysteresis loop of the base glass-composition.
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<figure><img src="image_4.png" /><figcaption>**Figure 2.** Optical basicity and Fe³⁺ % as a function of composition for the as-quenched glasses. (inset: Spectra of Li-0/Ca-0 spectra fit with three sub-spectra).</figcaption></figure>

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<figure><img src="image_5.png" /><figcaption>**Figure 3 (a and b).** Normalized, temperature and frequency corrected room-temperature Raman spectra for the as-quenched samples in the (a) Li-series and (b) Ca-series.</figcaption></figure>
<figure><img src="image_6.png" /><figcaption>**Figure 4 (a and b).** Qⁿ quantification for the Li-series obtained from the deconvolution of the silicate region. Dotted lines between points come from linear fitting of the data to serve as guides for the eye.</figcaption></figure>
<figure><img src="image_7.png" /><figcaption>**Figure 5 (a and b).** Q²/(Q³+Q¹) and Fe³⁺/Qₜtot for the Li-series and Ca-series obtained from the deconvolution of the silicate region. Dotted lines between points come from linear fitting of the data to serve as guides for the eye.</figcaption></figure>
<figure><img src="image_8.png" /><figcaption>**Figure 6.** Plot of experimentally obtained T<sub>g</sub> versus the composition for the glasses synthesized in the study. Dotted lines between points serve as guides for the eye.</figcaption></figure>

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<figure><img src="image_9.png" /><figcaption>**Figure 7 (a and b).** (a) X-ray diffraction patterns for Ca-series samples synthesized from melt quenching showing an amorphous nature. (b) X-ray diffraction patterns of the Ca-series samples heat treated at 950°C.</figcaption></figure>
<figure><img src="image_10.png" /><figcaption>**Figure 8 (a and b).** (a) Weight percentage crystalline fraction for the Li-series samples heat treated at 950 °C as obtained from Rietveld analysis with Al₂O₃ as an internal standard. (b) Weight percentage crystalline fraction for the Ca-series samples heat treated at 950 °C as obtained from Rietveld analysis with Al₂O₃ as an internal standard.</figcaption></figure>
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This article is protected by copyright. All rights reserved.

<figure><img src="image_11.png" /></figure>
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This article is protected by copyright. All rights reserved.

<figure><img src="image_12.png" /><figcaption>(b)</figcaption></figure>
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<figure><img src="image_13.png" /><figcaption>**Figure 9.** Secondary electron image and EDS-survey map of (a) a heat-treated Li-25 glass ceramic, (b) heat-treated Ca-25 glass ceramic (c) heat-treated Ca-25 glass ceramic sample showing Ca-rich regions.</figcaption></figure>
<figure><img src="image_14.png" /><figcaption>**Figure 10.** Magnetic hysteresis curves obtained for the samples of Li-series (blue) and Ca-series (red) heat treated at 950 °C. The dashed curves represent the raw data, while the solid curves represent the slope-corrected loops after removing the paramagnetic contribution.</figcaption></figure>
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<table><caption>**Table 1.** Summary of refined Mössbauer parameters for the room-temperature measurements of all glass samples in the Li-series and Ca-series as indicated in the table: centroid shift, δ, quadrupole splitting, ΔE<sub>Q</sub>, Lorentzian linewidth, Γ, and intensities, *I*, of the different components. Estimated errors are in *I* ±2%, in δ and ΔE<sub>Q</sub> ±0.005 mm/s, and in Γ ±0.02 mm/s.</caption><thead><tr><th>Components</th><th>Li-0</th><th>Li-5</th><th>Li-10</th><th>Li-15</th><th>Li-20</th><th>Li-25</th></tr></thead><tbody><tr><td>δ<sub>1</sub> (mm/s)</td><td>0.244</td><td>0.257</td><td>0.271</td><td>0.285</td><td>0.299</td><td>0.303</td></tr><tr><td>ΔE<sub>Q1</sub> (mm/s)</td><td>0.719</td><td>0.759</td><td>0.763</td><td>0.748</td><td>0.778</td><td>0.805</td></tr><tr><td>Γ<sub>1</sub> (mm/s)</td><td>0.43</td><td>0.46</td><td>0.47</td><td>0.45</td><td>0.47</td><td>0.47</td></tr><tr><td>I<sub>1</sub> (%)</td><td>45</td><td>50</td><td>47</td><td>41</td><td>41</td><td>44</td></tr><tr><td>δ<sub>2</sub> (mm/s)</td><td>0.255</td><td>0.263</td><td>0.277</td><td>0.285</td><td>0.291</td><td>0.294</td></tr><tr><td>ΔE<sub>Q2</sub> (mm/s)</td><td>1.187</td><td>1.254</td><td>1.273</td><td>1.265</td><td>1.302</td><td>1.342</td></tr><tr><td>Γ<sub>2</sub> (mm/s)</td><td>0.48</td><td>0.46</td><td>0.48</td><td>0.49</td><td>0.51</td><td>0.49</td></tr><tr><td>I<sub>3</sub> (%)</td><td>43</td><td>35</td><td>38</td><td>39</td><td>41</td><td>36</td></tr><tr><td>δ<sub>3</sub> (mm/s)</td><td>1.051</td><td>1.050</td><td>1.075</td><td>1.089</td><td>1.089</td><td>1.091</td></tr><tr><td>ΔE<sub>Q3</sub> (mm/s)</td><td>1.841</td><td>1.870</td><td>1.904</td><td>1.907</td><td>1.949</td><td>1.927</td></tr><tr><td>Γ<sub>3</sub> (mm/s)</td><td>0.64</td><td>0.68</td><td>0.65</td><td>0.69</td><td>0.66</td><td>0.66</td></tr><tr><td>I<sub>3</sub> (%)</td><td>12</td><td>15</td><td>15</td><td>20</td><td>18</td><td>20</td></tr><tr><td>Fe<sup>3+</sup>/Fe</td><td>88%±2%</td><td>85%±2%</td><td>85%±2%</td><td>80%±2%</td><td>82%±2%</td><td>80%±2%</td></tr></tbody></table><table><thead><tr><th>Components</th><th>Ca-0</th><th>Ca-5</th><th>Ca-10</th><th>Ca-15</th><th>Ca-20</th><th>Ca-25</th></tr></thead><tbody><tr><td>δ<sub>1</sub> (mm/s)</td><td>0.244</td><td>0.254</td><td>0.272</td><td>0.295</td><td>0.307</td><td>0.311</td></tr><tr><td>ΔE<sub>Q1</sub> (mm/s)</td><td>0.719</td><td>0.739</td><td>0.793</td><td>0.857</td><td>0.868</td><td>0.914</td></tr><tr><td>Γ<sub>1</sub> (mm/s)</td><td>0.43</td><td>0.45</td><td>0.47</td><td>0.51</td><td>0.48</td><td>0.50</td></tr><tr><td>I<sub>1</sub> (%)</td><td>45</td><td>43</td><td>46</td><td>47</td><td>41</td><td>36</td></tr></tbody></table>
<table><caption>**Table 2.** Band area fraction (%) and Qⁿ fraction of Raman bands in the 800 cm⁻¹ to 1250 cm⁻¹ region for Li-series and Ca-series.</caption><tr><td>δ<sub>2</sub> (mm/s)</td><td>0.255</td><td>0.261</td><td>0.277</td><td>0.289</td><td>0.295</td><td>0.297</td></tr><tr><td>ΔE<sub>Q2</sub> (mm/s)</td><td>1.187</td><td>1.224</td><td>1.305</td><td>1.398</td><td>1.444</td><td>1.488</td></tr><tr><td>Γ<sub>2</sub> (mm/s)</td><td>0.48</td><td>0.50</td><td>0.50</td><td>0.52</td><td>0.54</td><td>0.54</td></tr><tr><td>I<sub>3</sub> (%)</td><td>43</td><td>45</td><td>41</td><td>38</td><td>45</td><td>41</td></tr><tr><td>δ<sub>3</sub> (mm/s)</td><td>1.051</td><td>1.037</td><td>1.047</td><td>1.053</td><td>1.058</td><td>1.061</td></tr><tr><td>ΔE<sub>Q3</sub> (mm/s)</td><td>1.841</td><td>1.859</td><td>1.901</td><td>1.913</td><td>1.931</td><td>1.910</td></tr><tr><td>Γ<sub>3</sub> (mm/s)</td><td>0.64</td><td>0.63</td><td>0.61</td><td>0.63</td><td>0.62</td><td>0.65</td></tr><tr><td>I<sub>3</sub> (%)</td><td>12</td><td>12</td><td>13</td><td>15</td><td>14</td><td>23</td></tr><tr><td>Fe<sup>3+</sup>/Fe</td><td>88%±2%</td><td>88%±2%</td><td>87%±2%</td><td>85%±2%</td><td>86%±2%</td><td>77%±2%</td></tr></table>

<table><thead><tr><th rowspan="2"></th><th rowspan="2"></th><th>Q<sup>1</sup></th><th>Q<sup>2</sup></th><th>Fe<sup>3+</sup></th><th>T<sub>2</sub></th><th>Q<sup>3</sup></th><th>Q<sub>tot</sub></th></tr><tr><th>(890 cm<sup>-1</sup>)</th><th>(930 cm<sup>-1</sup>)</th><th>(980 cm<sup>-1</sup>)</th><th>(1030 cm<sup>-1</sup>)</th><th>(1080 cm<sup>-1</sup>)</th><th></th></tr></thead><tbody><tr><td rowspan="2">Li-0</td><td>Area fraction</td><td>7.52</td><td>37.96</td><td>20.08</td><td>19.26</td><td>15.16</td><td>60.64</td></tr><tr><td>Qⁿ fraction</td><td>12.40</td><td>62.60</td><td>-</td><td>-</td><td>25.00</td><td>-</td></tr><tr><td rowspan="2">Li-5</td><td>Area fraction</td><td>7.46</td><td>38.54</td><td>17.66</td><td>24.06</td><td>12.25</td><td>58.25</td></tr><tr><td>Qⁿ fraction</td><td>12.81</td><td>66.16</td><td>-</td><td>-</td><td>21.03</td><td>-</td></tr><tr><td rowspan="2">Li-10</td><td>Area fraction</td><td>7.58</td><td>39.21</td><td>18.57</td><td>22.16</td><td>12.45</td><td>59.24</td></tr><tr><td>Qⁿ fraction</td><td>12.80</td><td>66.19</td><td>-</td><td>-</td><td>21.02</td><td>-</td></tr></tbody></table>
<table><tbody><tr><td rowspan="2">Li-15</td><td>Area fraction</td><td>2.75</td><td>47.98</td><td>14.38</td><td>22.91</td><td>11.96</td><td>62.69</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>4.39</td><td>76.54</td><td>-</td><td>-</td><td>19.08</td><td>-</td></tr><tr><td rowspan="2">Li-20</td><td>Area fraction</td><td>3.11</td><td>47.29</td><td>15.29</td><td>22.77</td><td>11.52</td><td>61.92</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>5.02</td><td>76.37</td><td>-</td><td>-</td><td>18.60</td><td>-</td></tr><tr><td rowspan="2">Li-25</td><td>Area fraction</td><td>3.33</td><td>42.50</td><td>14.10</td><td>34.30</td><td>5.74</td><td>51.57</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>6.46</td><td>82.41</td><td>-</td><td>-</td><td>11.13</td><td>-</td></tr><tr><td rowspan="2">Ca-0</td><td>Area fraction</td><td>7.52</td><td>37.96</td><td>20.08</td><td>19.26</td><td>15.16</td><td>60.64</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>12.40</td><td>62.60</td><td>-</td><td>-</td><td>25.00</td><td>-</td></tr><tr><td rowspan="2">Ca-5</td><td>Area fraction</td><td>8.89</td><td>34.56</td><td>22.84</td><td>22.44</td><td>11.26</td><td>54.71</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>16.25</td><td>63.17</td><td>-</td><td>-</td><td>20.58</td><td>-</td></tr><tr><td rowspan="2">Ca-10</td><td>Area fraction</td><td>5.78</td><td>41.04</td><td>19.04</td><td>24.39</td><td>9.73</td><td>56.55</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>10.22</td><td>72.57</td><td>-</td><td>-</td><td>17.21</td><td>-</td></tr><tr><td rowspan="2">Ca-15</td><td>Area fraction</td><td>4.62</td><td>41.31</td><td>17.33</td><td>33.63</td><td>3.09</td><td>49.02</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>9.42</td><td>84.27</td><td>-</td><td>-</td><td>6.30</td><td>-</td></tr><tr><td rowspan="2">Ca-20</td><td>Area fraction</td><td>5.10</td><td>40.63</td><td>19.37</td><td>22.44</td><td>12.43</td><td>58.16</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>8.77</td><td>69.86</td><td>-</td><td>-</td><td>21.37</td><td>-</td></tr><tr><td rowspan="2">Ca-25</td><td>Area fraction</td><td>0.70</td><td>49.41</td><td>11.42</td><td>31.85</td><td>6.59</td><td>56.70</td></tr><tr><td>Q<sup>n</sup> fraction</td><td>1.23</td><td>87.14</td><td>-</td><td>-</td><td>11.62</td><td>-</td></tr></tbody></table>