John S Mccloy
email: john.mccloy@wsu.edu

Brian J Riley

Jarrod Crum

José Marcial

Joelle T Reiser

Karen Kruska

Jacob A Peterson

Daniel R Neuville

Deepak S Patil

Muad Saleh

Kristian E Barnsley

John V Hanna

Crystallization study of rare earth and molybdenum containing nuclear waste glass ceramics

Keywords: nuclear waste, nuclear magnetic resonance, rare earths, glass-ceramics, X-ray methods

ou non, émanant des établissements d'enseignement et de recherche français ou étrangers, des laboratoires publics ou privés.

Introduction

A glass-ceramic waste form is being developed through the US Department of Energy Office of Nuclear Energy to immobilize non-fissionable waste streams of alkali (A, e.g., 137 Cs), alkaline-earths (AE, e.g., 90 Sr), rare-earths (RE), and transition metals (most importantly, Mo) generated during a projected scheme of aqueous reprocessing of used nuclear fuel with transuranic extraction (TRUEX plus ). 1,2 Another waste form being considered is a multi-phase crystalline ceramic. [3][4][5] Compared to an alkali borosilicate glass waste form, the glass-ceramic waste form is expected to double waste loading to ~45% and have improved thermal stability. 2 These benefits are realized by the partitioning of the insoluble fission product fraction into a suite of chemically durable crystalline phases through controlled cooling. In practice, waste components plus glass-forming chemicals will be melted at high temperature (~1200-1450°C) using a cold crucible induction melter (CCIM) and poured into a canister to air cool. 1 During cooling, the melt phase separates into Mo-rich and RE-rich droplets within a borosilicate matrix, followed by crystallization. 2 For the specific US case of glass-ceramic waste forms, preliminary research on representative compositions has been conducted with crucible melt studies 2 and a pilot-scale CCIM. 1 Actual nuclear waste glass compositions will be comprised of 30+ components, but simplified 8-oxide systems are being currently studied in more detail, and are discussed here.

It is well-known that Mo has relatively low solubility in borosilicate glass. [6][7][8][9] Some nuclear waste forms target ways to increase Mo solubility in the glass phase by adding RE, 9 or high field strength modifiers like Li, 10 Mg, 11 or V, [START_REF] Mckeown | X-ray absorption and Raman spectroscopy studies of molybdenum environments in borosilicate waste glasses[END_REF] for instance. The importance of RE for the stabilization of Mo in the glass matrix has been well-described, along with the effect of removing RE from the glass by crystallization of oxyapatite. 13- 17 At the same time, when Mo does precipitate in a crystalline phase, it is desirable to avoid the formation of water-soluble alkali molybdates above the Mo solubility limit, by controlling the distribution of Na and Cs, to force the more chemically-durable calcium molybdate (powellite) phase to form preferentially. [18][19][20][21][22] A key barrier to maturation and exploitation of glass-ceramic technologies for nuclear waste forms is the gap in fundamental understanding of the molecular-scale mechanisms of phase separation and crystallization leading to the development of the desired phase assemblage and microstructure, ultimately determining long-term product performance. 1,23 The challenge is in predictably achieving the targeted phase assemblage and microstructure, requiring a detailed understanding of the transformation process as a function of both cooling rate and melt chemistry. This study uses a simplified 8-oxide system representing elements in the main chemical categories of the waste. Using this composition, the effects of cooling profiles on the structure and chemical partitioning is investigated.

It should be noted that there is a trade-off between simplifying a complex glass in order to better understand it and causing behavior to be different than the comparable complex glass. For instance, it is known that some of the alkali behave differently in these glass-ceramic systems, in that Cs tends to stay in the glassy phase or rarely partition to a soluble molybdate, 1,18 Li goes normally to the glass, 1 and Na can go to the glass, powellite, or oxyapatite phase. 17 For alkaline earths, the presence of Ba and Sr cause a separate partitioning of two powellite crystalline phases with different amounts of Ca, 1,24 which would not be reproduced when including only Ca. Finally, it has been shown that the specific rare earth does somewhat influence the rare phases formed, 17 but given that Nd is by far the most prevalent rare earth (see Table 1), Nd was chosen. The 8-oxide composition studied in detail here still provides the basic chemistry of charge compensation (Na), powellite inclusion (mainly Ca), and oxyapatite formation (Nd, with contributions from Ca and sometimes Na), such that useful physical insight can be gained without excessive compositional complexity.

Experimental

Glass synthesis

The starting glass composition for the present study, denoted as "SIMP" (see Table 1), is a simplified version of GC-Mo-5.86 as formulated previously. 1 For this simplified glass, all of the alkalis were combined by mass into Na 2 O, all the alkaline earths into CaO, all the rare earths into Nd 2 O 3 , all the minor components F o r P e e r R e v i e w 4 were eliminated, and the composition was renormalized. After batching and melting oxides at 1350°C for 1 h, the glass was quenched on an Inconel ® plate. For some experiments, the glass was subsequently doped with 0.05 mass% RuO 2 as a nucleating agent, then remelted and quenched (denoted as "Ru-SIMP").

Heat-treatments

In order to capture phase separation and crystallization at different stages, batches of ~3 g of crushed quenched glass were remelted (~1350°C) in a ~5-mL thimble-shaped Pt/10%Rh crucible with a tight-fitting Pt/10%Rh lid and then quickly cooled to the desired soak temperatures (separate samples for each soak, with T soak = 1150 to 1450°C in 25°C increments), held for 30 min at the soak temperature, and then rapidly quenched in a water bath. Note that, in some cases (i.e., 1375-1450°C), these soak temperatures were greater than the original melt temperature of the glass. To facilitate the rapid quenching, each crucible was hung from a thin Pt wire hooked over an alumina rod at the top of the furnace. After holding at the desired temperature, the wire was lifted off the alumina rod and the crucible was dropped through the bottom of the furnace and into a container of deionized water, thus nominally freezing the melt structure (and crystal phases, if any) present at the soak temperature. No water contacted the glass melt due to the crucible geometry.

There was some potential concern about volatilization at elevated temperatures. Thermal analysis experiments on small samples of fine-grained powders in an open crucible indicated that up to 3.5% of the mass of the SIMP glass could be lost in heating from room temperature to 1200°C over a span of 2 hours (see Supplemental Information). Thus, it is possible that there was some volatilization during the soak-andquench heat treatments. However, the tight lid should have reduced preferential element loss, but this was not quantified. Some glasses were cooled slowly (1×SC schedule as described previously 1 ) at the rate anticipated for the centerline of a typical 0.61-m (2-ft) diameter canister, 25 to enable a direct comparison with the microstructures of the rapidly-quenched samples. For these experiments, the simplified glass composition (SIMP), the RuO 2 -doped one (Ru-SIMP), and a more complex representative "centroid" composition (denoted CX, see Table 1) were slow cooled.

Characterization

Thermal analysis

To assess the glass transition (T g ) and crystallization (T x ) temperatures, differential scanning calorimetry (DSC) was performed using a Netzsch STA 449F1. Experiments were performed with ~35-45 mg of crushed glasses, heating from 50 to 1200°C at 10°C/min, then subsequent cooling at 10°C/min to 50°C.

Viscosity

Low-temperature viscosity for the SIMP glass was measured from 585°C to ~660°C with a creep apparatus following the procedures described by Neuville. 26 Glass samples were prepared by water quenching melts from 1300°C, annealing from 500°C to room temperature over 16 hours, and subsequently core-drilling using a high-pressure water jet. Each viscosity datapoint corresponds to an average of between 15 and 40 measurements, made at the same temperature with different stresses, to check for Newtonian behavior of the glass. Expected accuracy is 0.03 log units of viscosity. Data were fit to the Vogel-Fulcher-Tammann (VFT) non-Arrhenius law, and the T g was obtained. This T g value, calculated for viscosity equal to 10 12 Pa•s, was expected to be lower than that obtained by DSC because of the different speed of heating.

High-temperature viscosity for the SIMP glass was measured from 1300°C to ~1050°C using the procedure described in Crum et al. 27 based on ASTM C 965-94. 28 Quenched glass (30 mL container) was loaded into a platinum crucible. Measurements were collected with an FRS1600 rheometer (Anton Paar GmbH, Austria) and a platinum spindle. The temperature program was set at a ramp rate of 8°C min -1 and a 30-min dwell time. Measurements started at 1300°C with 100 data points taken over a 5-min interval for each measurement interval to collect shear stress measurements as a function of shear rate. The temperature was decreased in 50°C increments to the final temperature of 1050°C. The spindle was not rotated during the hold times to avoid impacting any crystals and phase separation in the melts.

Electron microscopy

Scanning electron microscopy (SEM) was used to obtain micrographs on some of the samples after their soak and quench. A JEOL JSM-7001F field emission gun SEM (JEOL USA, Inc., Peabody, MA) with a backscattered electron detector was used for imaging.

Selected samples were investigated with scanning transmission and transmission electron microscopies (STEM/TEM). Here, heat-treated samples were pulverized in a Diamonite TM mortar and pestle and then thinly dispersed on a standard lacey carbon TEM grid. A JEOL ARM 200CF (JEOL USA, Inc., Peabody, MA) equipped with a high-angle annular dark-field (HAADF) detector and a Centurio silicon drift energy dispersive spectrometer (EDS) operated at 200 kV was used for the STEM/EDS analysis. The STEM was operated at 5 µA beam current with a 40-µm condenser aperture (CA) to minimize beam damage during imaging. EDS analysis was carried out at 15 µA with a 70-µm CA. Some degradation of the glass matrix was observed during imaging and EDS analysis, but crystals were found to be relatively stable. Pathfinder (v1.0, Thermo Fisher Scientific Inc.) software was used for recording and analysis of the EDS data.

Selected samples were investigated using electron probe microanalysis (EPMA) using a JEOL JXA-8500F field emission electron microprobe (JEOL USA, Inc., Peabody, MA) equipped with five wavelengthdispersive X-ray spectrometers (WDS), see Supplemental Information for measurement parameters.

Small angle X-ray scattering

Small angle scattering techniques were used to investigate the possible presence of phase separation in the soak-and-quench glasses. Dimensionality accessed in these techniques is referred to in "q space", which is the scattering vector (reciprocal space) magnitude, i.e., q = (4πsinθ)/λ = 2π/d, where θ is the scattered angle, λ is the wavelength, and d is the interplanar spacing. Initially, small angle neutron scattering was attempted (see Supplemental), but small and ultra-small angle X-ray scattering (SAXS and USAXS, respectively) measurements proved more useful.

Overall, one set of experiment allowed feature sizes to be probed from 4π/q max ~ 9 Å to π/q min ~ 10.5 μm, thus covering wide-angle X-ray scattering (WAXS), SAXS, and USAXS ranges. A subset of the quenched samples was measured at the Advanced Photon Source (APS) at beamline 9 ID-C (setup described elsewhere 30,31 ) at Argonne National Laboratory (ANL) in Lemont, IL. A layer of powder ~0.1 mm thick was placed on cellophane tape for X-ray scattering measurements. The incident X-ray energy was 21 keV (λ = 0.5904 Å), and the beam size was 0.8×1.2 mm. The camera system at beamline 9-ID-C was capable of recording SAXS curves with an angular resolution of ~8×10 -4 Å -1 in scattering vector q range from 10 -4 to 1.2 Å -1 . 32 The USAXS instrument ran a slit-smeared Bonse-Hart setup and, in this experiment, the slit length was 0.027843 Å -1 . Slit smearing was included in data analysis models as needed. Each sample was analyzed for five minutes.

USAXS and SAXS data were completely corrected for instrumental variations, reduced, and analyzed with Igor Pro ® (v.7.0.5.2) coupled with Nika (v.1.74) and Irena (v.2.61) packages. 31,33 The samples were analyzed with the unified method developed by Beaucage and Schafer 34,35 using two structural levels to specifically examine the radius of gyration (R g ). From this model, R g characterized the size of the scattered material regardless of shape. 36 For modeling the SAXS and USAXS data, the matrix was assumed to have the composition of the overall glass, and the scatterers were composed of spherically shaped CaMoO 4 , as indicated by TEM (see Results).

X-ray diffraction

In addition to the synchrotron-obtained WAXS data on the soaked and quenched powders, other X-ray diffraction (XRD) data were collected using laboratory diffractometers. All samples were prepared by crushing to fine powders using a Diamonite TM mortar and pestle and using a zero-background silicon holder due to small sample volumes.

Standard room-temperature powder XRD was obtained for all quenched samples (1150-1450°C soaks) on a PANalytical X'pert Pro MPD (the Netherlands) using a Cu Kα source and X'celerator TM detector.

Rietveld refinement was performed using Highscore Plus (PANalytical) software. For selected samples, powders were doped with 10 mass% Al 2 O 3 internal standard for quantification of the amorphous fraction.

XRD was performed on three compositions of slow-cooled samples (1×SC) on a D8 Advance diffractometer (Bruker AXS Inc., Madison, WI) with a Cu Kα source and a LynxEye TM position-sensitive detector. The simplified composition (SIMP), the RuO 2 -doped version (Ru-SIMP), and the complex composition (CX) were all analyzed after 1×SC and mixing with a known quantity of TiO 2 standard (SRM-674b; National Institute of Standards and Technology), using Rietveld refinement and the Topas (Bruker) software, in order to obtain the amorphous fraction. Errors on the quantification of phases are estimated at a maximum of ~0.1 × mass% for major phases (oxyapatite), and 0.25 × mass% for minor phases (powellite), as assessed by similar detailed studies using the same quantitative XRD methods. 37,38 Additionally, hot-stage XRD was obtained on the same D8 Advance diffractometer. A powder of the SIMP glass was suspended in ethanol, placed onto a Pt heater strip, and allowed to dry. Then, the sample was heated from 30°C to 975°C at ~10°C s -1 , held for 24 hrs, then cooled at ~10°C s -1 to 30°C, with scans collected at 30°C, at 975°C, and then every ~1 hr during the hold at 975°C. Data was analyzed using Topas (v4.2, Bruker) software using Rietveld refinement to fit the relative fractions of crystalline phases and obtain lattice parameters. Because of the nature of the experiment, the amorphous fraction could not be obtained. To estimate the total crystal fraction at the beginning and end of this experiment, a sample of SIMP was heated in a furnace to 975°C at ~10°C s -1 , held 2 min, and air quenched; a separate sample was heated to 975°C at ~10°C s -1 , held for 24 h, and air quenched. Samples were mixed with Al 2 O 3 for quantification by XRD (PANalytical X'Pert Pro diffractometer) and Rietveld analysis (Highscore).

Optical absorption

Diffuse reflection was performed on powdered samples at room temperature using an Evolution™ 260 Bio UV-Visible Spectrophotometer (Thermo Fisher Scientific) with installed ISA-220 Integrating Sphere Accessory (Thermo Fisher Scientific). Measurements were performed from 350 nm to 1100 nm with 0.1-s integration time, 2-nm bandwidth, and 1-nm step size; note that a 0.5-nm step size was used for a pure ceramic Ca 2 Nd 8 (SiO 4 ) 6 O 2 ("Ca-Nd oxyapatite") standard made by solid state synthesis described elsewhere. 24 The spectra in the range between 420-440 nm were remeasured at 1-s integration time, 1-nm bandwidth, and 0.2-nm step size; the measurements were repeated five times and averaged. Diffuse reflectance was measured relative to a standard, calibrated Spectralon® (Labsphere, NH, USA) in this case.

(K-M) function [ ] using 𝑟 ∞ 𝑓(𝑟 ∞ ) Equation (1) (1) 𝑓(𝑟 ∞ ) = (1 -𝑟 ∞ ) 2 2 * 𝑟 ∞ = 𝐾 𝑆
where is the absorption coefficient and S is the scattering coefficient. 39 The K-M will simply be referred 𝐾 to as absorption in this paper.

23 Na nuclear magnetic resonance

A subset of SIMP samples were subjected to magic angle spinning (MAS) nuclear magnetic resonance

Thermal analysis

The DSC spectra, after background removal, are shown in Figure S1. Positions of the crystallization peaks show that, on heating, the RuO 2 -doped sample crystallized at lower temperature than the undoped sample, and has an additional higher temperature crystallization peak as well. Measured temperatures from the heating experiment are summarized in Table 2. On heating, there was evidence of two crystallization peaks for both glasses, while on cooling, there was evidence of only one crystallization peak, at lower temperature than for heating.

Viscosity

A summary of the low-temperature and high-temperature viscosities is provided in Table 3. The whole set of data shown in Table 3 was fit using the Vogel-Fulcher-Tammann (VFT) approach as follows:

log  = A + B/(T -T 1 ) (2)
where A, B, and T 1 were determined to be -4.4073±0.514, 4972.3±589, and 565.16±28, respectively, using Igor Pro (v6.37) and the T g was determined as 595.1°C (868.2 K) at 10 12 Pa•s. The data is plotted in Figure 1. A high yield stress is observed in the 1250-1150°C range, which could be related to active phase separation and crystallization. The viscosity behaves normally between 1100 and 1050°C.

Low temperature viscosities around T g were measured with different strains, which confirm a Newtonian behavior and prove that the liquid is crystal-free. 26 For high-temperature viscosities at 1300, 1250, and 1200°C, dispersions are Newtonian, but those at 1150, 1100, and 1050°C are not, suggesting the presence of phase separation and/or crystallites in the liquid. This can be seen in the larger error bars of the 1150-1050°C viscosity measurements as well. The crystal effect noted in the high-temperature viscosity measurements are clearly visible in the SEM images in Figure 2.

Microscopy

Visibly, water-quenched samples from T ≤ 1225°C were cloudy and/or opaque, while those quenched at T ≥ 1250°C were transparent. All samples were lilac to purple colored from the high neodymium As shown, the volume fraction of oxyapatite crystals notably decreased as the soak temperature was increased, in agreement with XRD results (see below). Spherical molybdate domains were found in each sample, with the locations being isolated to the perimeter of oxyapatite crystals at T ≥ 1225°C; this effect is most clearly seen in Figure 2d. At these temperatures, the removal of REs from the melt by crystallization destabilized the melt and resulted in the phase separation of spherical Mo-rich droplets. In general, these phase separation features appeared much less pronounced in this simplified glass compared to the complex glass-ceramic compositions examined previously. 1,2 Final microstructures of heat-treated samples show modifications of the microstructure from the soak-and-quench samples. The isothermally heat-treated sample from the hot-stage XRD showed similarly sized oxyapatite crystals, but in much larger concentrations (Figure 3a), while the slow-cooled SIMP sample showed larger, but not as concentrated, oxyapatite crystals (Figure 3b).

TEM micrographs were obtained for water-quenched samples from 1200, 1250, and 1300°C; some of this data is shown in Figure 4. TEM-EDS maps and line scans showed that these samples were enriched in Ca and Mo and depleted in Si relative to the matrix (see Figure 4d). Morphologically, these spheres looked more like "partial moons," since the volume of the liquid shrank on cooling more extensively than the glass, and the calcium molybdate crystals occupied a smaller volume than that of the spherical liquid phase. This phenomenon had also been observed in previous studies. 41

Small angle scattering

The USAXS and SAXS spectra are shown in Figure 5a. (see Supplemental Information for small angle neutron scattering data). The SAXS/USAXS data may have had a scaling problem, as it could not be determined for certain whether the low q values should have lain on top of one another, in which case the intensity differences were due to slightly different thicknesses or packing factors of the samples (e.g., mounting inconsistencies and particle size differences). However, the intensity differences do not affect the analysis of the scatterer sizes as the Guinier region will remain the same regardless of intensity. Qualitative increased as T soak decreased. Quantitative estimates for R g in Table 4 showed a decreasing trend in R g values as soak temperature increased from 1175 to 1250°C.

In Figure 5a, samples water quenched from 1175, 1200, and 1250°C showed increased scattering in the 0.01 Å -1 region that was not present in 1300 and 1350°C samples. Increased scattering for the 1175°C sample is seen from 0.0001 to 0.03 Å -1 , with the presence of a fairly indistinct Guinier curve with an R g around 0.004 Å -1 . This behavior was likely due to scattering from particles of various sizes (diameters)

and/or shapes. Samples quenched from 1200 and 1250°C showed the presence of more defined Guinier curves (with distinct R g values), which indicated scattering populations are more homogenous in size and/or shape. TEM micrographs showed CaMoO 4 'spheres' which likely caused increased scattering and R g for 1175, 1200, and 1250°C (see Figure 4). The 1300 and 1350°C samples did not show increased scattering in the same q region, rather increased scattering is seen at 0.03-0.2 Å -1 . The R g values for 1300 and 1350°C samples were overly low and unphysical. The other samples likely had the same scattering, but it was overshadowed by scattering due to CaMoO 4 .

Diffraction

WAXS spectra (Figure 5b) obtained at the same time as the USAXS/SAXS showed increase of intensity from q ~1.2 to ~2.2 Å -1 as was also observed in the SAXS data at high q (Figure 5a). Diffraction peaks were visible in the WAXS of 1175, 1200, and 1250°C samples, and some peaks could be observed in the SAXS of 1175 and 1250°C samples, particularly at q = 0.7714, 1.53838, and 1.63355 Å -1 . Quantitative fits were obtained by Rietveld refinements of the WAXS (crystalline fraction only) and corresponding room temperature XRD (quantified amorphous phase) as shown in Table 5. For completeness, all the samples that were cooled from the melt, soaked, then water quenched were subjected to room temperature XRD (Cu X-ray tube) (see Figure 6). The combined data showed that the 1250°C sample was the highest temperature measured with obvious Bragg peaks. As seen in 7b). For the room temperature scan at the beginning of the experiment, the only phase present was Ca-powellite, but Nd-powellite and oxyapatite-1 formed within the first hour at 975°C, with oxyapatite-2 forming beginning at hour 5. The relative crystalline fractions of powellite continued to decrease until ~8 h when it became constant at ~10 mass%. The oxyapatite was the majority phase. Lattice parameters of the oxyapatite phases were observed to decrease as time progressed (Figure 7c,d) possibly suggesting some effects of ion exchange or, more likely, selective volatilization of specific components.

Imaging and chemical analysis of the phases on the Pt heater strip after hot-stage XRD were attempted.

The inability to polish the sample, as well as the small size of the crystals, made obtaining high-quality data difficult. Mo-rich (powellite) crystals were present but were too small (< 0.5 μm) to adequately quantify with the 1-μm beam size of the EPMA. Oxyapatite crystals were likewise small, ~1-μm long and ~0.1-μm in diameter (see Figure 2). To further assess the effect of this different heat treatment, the ex-situ experiments were performed for essentially a t = 0 and t = 24 point for crystallization at 975°C. The conventional XRD results of powdered samples showed a high degree of crystallization at 24 h (~50 mass% amorphous), which was very close to that determined independently on the hot-stage XRD sample by image analysis using ImageJ software (~59 mass% amorphous).

The crystal fractions obtained by ex-situ XRD on the soak-and-quench glasses were 2-6 mass%, with the difference being amorphous residual glass. These crystal fractions were quite low compared to typical slow-cool examples, where crystal fractions were typically 24-36 mass%. 2 To directly compare the crystal fraction and assess compositional effects, 1×SC tests were performed on the simplified SIMP composition, the Ru-SIMP, and the CX complex composition. The crystal fraction was nearly the same in the CX, SIMP, and Ru-SIMP samples, so there was no apparent effect of composition, and the main effect appeared to be the slow cooling.

As discussed in a separate paper, 45 the processes for acquiring reliable and representative diffraction data on glass-ceramic samples using ex-situ and/or in-situ techniques have limitations. The limitations with heat treating samples in sealed capsules include melt-capsule reactions (e.g., if silica, sapphire) or issues with beam attenuation from the capsule wall (e.g., platinum). Additionally, it is difficult to monitor the amorphous fraction for in-situ measurements since a diffraction standard cannot be introduced into the scan data; thus, crystallographic phase distributions can be extracted but quantification is most likely impossible for in-situ diffraction measurement techniques. However, while ex-situ is an option, analysis requires specimens to be cooled prior to analysis, and forcing specimens to undergo cooling results in microstructural changes during the cool-down process and introduces artifacts into the data. The primary artifacts are in the correlations drawn between specimens undergoing cooling and those that already underwent cooling.

Optical absorption

Absorption (K-M) spectra were obtained for selected samples and compared to a single-phase ceramic Ca-Nd oxyapatite (Figure 8). Only the Nd 3+ 4f-4f electronic transitions were seen, and the absorption peaks were relatively broad and convoluted in the 1350°C soak sample. But, as the thermal soak temperature decreased, the peaks appeared sharper and more distinct, showing more resemblance to the features in single-phase Ca-Nd oxyapatite ceramic, indicating that more Nd was in crystalline phase(s). The sharpness of absorption features was best seen in the 4 I 9/2 → 4 G 9/2 , 4 G 7/2 , and 2 K 13/2 (~520 nm); 4 I 9/2 → 4 G 5/2 and 2 G 7/2 (~580 nm); 4 I 9/2 → 4 S 3/2 and 4 F 7/2 (~745 nm); and 4 I 9/2 → 4 F 3/2 (~871 nm) transitions. 46,47 The transitions of 4 I 9/2 → 4 G 5/2, known as the hypersensitive transition due to sensitivities to changes in the Nd 3+ crystal field and symmetry, 13,[START_REF] Quintas | Effect of alkali and alkalineearth cations on the neodymium environment in a rare-earth rich aluminoborosilicate glass[END_REF]49 was of particular interest; however, not much change in the peak position or the shape of that peak was seen in samples soaked between 1350°C and 1250°C, indicating no significant changes in the local environment of the Nd 3+ ion. Below 1250°C soak temperature, the peaks became → 2 P 1/2 , are provided in the Supplemental Information.

Nuclear magnetic resonance

The 23 Na MAS NMR showed the presence of a sharp peak when the soak temperature was 1150°C (Figure 9). This sharp peak was on top of the broad peak of Na in the glass, which was evident in all the other samples. The sharp peak located at -0.7 ppm is indicative of Na in a crystalline phase. It is probable that this Na was in oxyapatite for the reasons described hereafter. The amount of oxyapatite in the glassceramic increased with decreasing soak temperature (Table 5), while the amount of powellite remained about the same. While the 1150°C was not quantified, the oxyapatite concentration should be somewhat more than the quantified 1175°C sample if it followed the trends at higher soak temperatures. Alternatively, the lower soak temperature could have allowed Na to enter the oxyapatite and its composition was somewhat different than at higher temperatures. This was suggested by the hot-stage XRD data where two different oxyapatites (based on differing unit cell parameters) were observed, but it was possible that the phase changes in the hot stage experiment were due to sample composition change during volatilization.

Regardless, it is known that Na can go into the Ca-Nd oxyapatite or form a Na-Nd oxyapatite on its own. 50 The SEM-EDS investigations were not helpful here, as the Na and Nd X-ray lines overlap somewhat (i.e., Na Kα = 1.040 keV and Nd Mα = 0.978 keV); however this was observed in the EPMA data of the 1×SC SIMP sample (see Supplemental Information).

Synthesis and 23 Na measurement of standards, Na 0.5 La 0.5 MoO 4 powellite and NaLa 9 (SiO 4 ) 6 O 2 apatite, confirmed that the most likely assignment of the 23 Na peak in the glass-ceramics was Na in oxyapatite. Note that, for the standards, La was used as the rare earth to avoid any possible problems with paramagnetic Nd, but the Na environment should be very close to the same in the Nd-containing oxyapatite. The NMR showed that it was very unlikely that the Na signature was from Na-containing powellite or Na 2 MoO 4 . To further explore the identity of the crystalline phase containing Na, T 1 measurements were conducted to investigate the relaxation time of the glass versus the crystalline phase. The T 1 of the narrow crystalline phase seen in

F o r P e e r R e v i e w 16
the 23 Na MAS NMR is an order of magnitude less than that of the broader glass phase (0.13 ± 0.02 s the glass versus 0.011 ± 0.001 s for the crystalline phase, see Supplemental Information for fits). Differences in the T 1 values are likely due to the proximity of the paramagnetic Nd nuclei, with closer sites having shortened T 1 values due to the paramagnetic interaction. This result suggested that the Na-bearing crystalline phase contained Nd, which was far more likely for an oxyapatite phase than for a Na molybdate phase.

Summary and Conclusions

Various heat treatment protocols were used to study phase separation and crystallization of a simplified glass-ceramic composition, being investigated as a waste form for aqueous wastes generated during used nuclear fuel reprocessing. The techniques used during this study included scanning, scanning transmission, and transmission electron microscopies; small angle neutron scattering; wide-angle, small-angle, and ultrasmall angle X-ray scattering; room temperature and hot-stage X-ray diffraction; optical absorption; and 23 Na MAS NMR.

Taken together, the results for the SIMP sample suggest that the 1300-1250°C temperature range is the critical soak temperature where amorphous melts can be maintained and, if quenched fast enough, can yield a product without phase separation and crystallization in this composition. Soak temperatures above this threshold maintain relatively homogeneous glass upon quenching (per TEM) and no evident crystallization (by XRD and WAXS). SAXS analysis for 1350°C and 1300°C show Guinier R g values < 2 nm, suggesting there are no significant scattering phases in these samples at room temperature. The 1250°C soak sample showed spherical (~20-nm diameter) calcium molybdates in TEM, and small fractions of powellite and oxyapatite (~0.5 and ~1.0 mass%, respectively) in XRD and WAXS, superimposed on a broad glassy background. SAXS-obtained R g values of the 1250°C sample are 7.3 nm (radius), which correspond well to the ~20 nm particles observed in TEM in this sample. At T soak ≤ 1225°C, quenched samples were visibly cloudy, and showed increasing fractions of crystals, although still very small amounts (~6 mass% total at 1175°C), but large enough to cause evident changes in the optical absorption spectra of Nd 3+ . ). This paper shows that crystallization studies using nucleation and growth, such as the quenching and heat treating at constant high temperature of 975°C, do not produce a microstructure and crystal fraction representative of melts that are slow-cooled and crystallized from high temperature. In the future, these results should be confirmed by in-situ diffraction from high temperatures in a sealed container, so that volatilization does not cause compositional changes and different phase evolution as suggested by laboratory hot-stage XRD.

Thus, while the goal of the current work was to track the microstructural evolution in glass-ceramic samples, several data gaps still exist. The main issues for in-situ diffraction analysis (like hot-stage XRD) are to contain volatile samples during the experiment and identify a sample container that is passive enough not to interfere with the experiment, i.e., it does not react with the melt and does not (significantly) attenuate the interrogation beam. Regarding ex-situ analysis, if samples could be cooled quickly enough, it is possible that they could be interrogated with ex-situ analysis and the microstructure would be representative of the high-temperature structure. This was the overall approach of the current study. Solving the in-situ

diffraction issues would open the possibilities of characterizing a wide range of waste-form materials for optimization of microstructure, chemistry, phase distribution, and chemical durability.

From a waste form perspective, controlling crystallization in a predicable way, using primarily compositional control, is highly beneficial when the precipitated crystals are chemically durable. The durability of oxyapatite and powellite crystalline phases, as well as the nominal residual glass composition, has been recently demonstrated. 24,51 For conventional borosilicate waste glasses, crystallization is generally unwanted, as it usually results in melter processing problems or poor durability of the residual glass. 52,53 Defense waste glasses (such as those planned for Hanford) are so compositionally diverse 54 that it would be difficult to formulate compositions for making glass-ceramic waste forms with predictable crystalline phases that do not cause either problems in melter processing or in product performance due to

F o r P e e r R e v i e w
18 residual glass durability. For the particular fuel reprocessing waste stream considered here, the problem is easier because the phases form upon cooling so they do not cause melter processing problems, the phases that form are chemically durable, and the residual glass is also durable. The primary gap that remains with this glass-ceramic waste form is to refine the control of microstructure and phase distribution as a function of cooling rate and composition, which can be achieved in order to eliminate any chemical durability issues induced by interfacial stresses between the glass and crystalline phases. Figure 9. 23 Na MAS NMR spectra of SIMP. The sharp peak in the 1150°C sample is located at -0.7 ppm. The Na-containing oxyapatite peaks at 1.5 ppm, the Na-containing powellite has a quadrupolar lineshape peaking around 12.5 ppm, and the Na-molybdate has a complex lineshape with the main peak around -1 ppm.

Introduction

A glass-ceramic waste form is being developed through the US Department of Energy Office of Nuclear Energy to immobilize non-fissionable waste streams of alkali (A, e.g., 137 Cs), alkaline-earths (AE, e.g., 90 Sr), rare-earths (RE), and transition metals (most importantly, Mo) generated during a projected scheme of aqueous reprocessing of used nuclear fuel with transuranic extraction (TRUEX plus ). 1,2 Another waste form being considered is a multi-phase crystalline ceramic. [3][4][5] Compared to an alkali borosilicate glass waste form, the glass-ceramic product waste form is expected to double waste loading to ~45% and have improved thermal stability. 2 These benefits are realized by the partitioning of the insoluble fission product fraction into a suite of chemically durable crystalline phases through controlled cooling. In practice, waste components plus glass-forming chemicals will be melted at high temperature (~1200-1450°C) using a cold crucible induction melter (CCIM) and poured into a canister to air cool. 1 2 For the specific US case of glass-ceramic waste forms, preliminary research on representative compositions has been conducted with crucible melt studies 2 and a pilot-scale CCIM. 1 Actual nuclear waste glass compositions will be comprised of 30+ components, but simplified 8-oxide systems are being currently studied in more detail, and are discussed here.

It is well-known that Mo has relatively low solubility in borosilicate glass. [6][7][8][9] Some nuclear waste forms target ways to increase Mo solubility in the glass phase by adding RE, 9 or high field strength modifiers like Li, 10 Mg, 11 or V, [START_REF] Mckeown | X-ray absorption and Raman spectroscopy studies of molybdenum environments in borosilicate waste glasses[END_REF] for instance. The importance of RE for the stabilization of Mo in the glass matrix has been well-described, along with the effect of removing RE from the glass by crystallization of oxyapatite. 17 At the same time, when Mo does precipitate in a crystalline phase, it is desirable to avoid the formation of water-soluble alkali molybdates above the Mo solubility limit, by controlling the distribution of Na and Cs, to force the more chemically-durable calcium molybdate (powellite) phase to form preferentially. [18][19][20][21][22] A key barrier to maturation and exploitation of glass-ceramic technologies for nuclear waste forms is the gap in fundamental understanding of the molecular-scale mechanisms of phase separation and crystallization leading to the development of the desired phase assemblage and microstructure, ultimately determining long-term product performance. 1,23 The challenge is in predictably achieving the targeted phase assemblage and microstructure, requiring a detailed understanding of the transformation process as a function of both cooling rate and melt chemistry. This study uses a simplified 8-oxide system representing elements in the main chemical categories of the waste. Using this composition, the effects of cooling profiles on the structure and chemical partitioning is investigated.

It should be noted that there is a trade-off between simplifying a complex glass in order to better understand it and causing behavior to be different than the comparable complex glass. For instance, it is known that some of the alkali behave differently in these glass-ceramic systems, in that Cs tends to stay in the glassy phase or rarely partition to a soluble molybdate, 1,18 Li goes normally to the glass, 1 and Na can go to the glass, powellite, or oxyapatite phase. 17 For alkaline earths, the presence of Ba and Sr cause a separate partitioning of two powellite crystalline phases with different amounts of Ca, 1, 24 which would not be reproduced when including only Ca. Finally, it has been shown that the specific rare earth does somewhat influence the rare phases formed, 17 but given that Nd is by far the most prevalent rare earth (see Table 1), Nd was chosen. The 8-oxide composition studied in detail here still provides the basic chemistry of charge compensation (Na), powellite inclusion (mainly Ca), and oxyapatite formation (Nd, with contributions from Ca and sometimes Na), such that useful physical insight can be gained without excessive compositional complexity.

Glass synthesis

The starting glass composition for the present study, denoted as "GCSIMP" (see Table 1), is a simplified version of GC-Mo-5.86 as formulated previously. 1 For this simplified glass, all of the alkalis were combined by mass into Na2O, all the alkaline earths into CaO, all the rare earths into Nd2O3, all the minor components were eliminated, and the composition was renormalized. All B2O3 was replaced with 11 B2O3 to avoid excessive neutron absorption in neutron probe experiments. After batching and melting oxides at 1350°C for 1 h, the glass was quenched on an Inconel ® plate. For some experiments, the glass was subsequently doped with 0.05 mass% RuO2 as a nucleating agent, then remelted and quenched (denoted as "Ru-GCSIMP").

Heat-treatments

In order to capture phase separation and crystallization at different stages, batches of ~3 g of crushed quenched glass were remelted (~1350°C) in a ~5-mL thimble-shaped Pt/10%Rh crucible with a tight-fitting Pt/10%Rh lid and then quickly cooled to the desired soak temperatures (separate samples for each soak, with Tsoak = 1150 to 1450°C in 25°C increments), held for 30 min at the soak temperature, and then rapidly quenched in a water bath. Note that, in some cases (i.e., 1375-1450°C), these soak temperatures were greater than the original melt temperature of the glass. To facilitate the rapid quenching, each crucible was hung from a thin Pt wire hooked over an alumina rod at the top of the furnace. After holding at the desired temperature, the wire was lifted off the alumina rod and the crucible was dropped through the bottom of the furnace and into a container of deionized water, thus nominally freezing the melt structure (and crystal phases, if any) present at the soak temperature. No water contacted the glass melt due to the crucible geometry.

There was some potential concern about volatilization at elevated temperatures. Thermal analysis experiments on small samples of fine-grained powders in an open crucible indicated that up to 3.5% of the mass of the SIMP glass could be lost in heating from room temperature to 1200°C over a span of 2 hours (see Supplementary Information). Thus, it is possible that there was some volatilization during the soakand-quench heat treatments. However, the tight lid should have reduced preferential element loss, but this was not quantified.

Some glasses were cooled slowly (1×SC schedule as described previously 1 ) at the rate anticipated for the centerline of a typical 0.61-m (2-ft) diameter canister, 25 to enable a direct comparison with the microstructures of the rapidly-quenched samples. For these experiments, the simplified glass composition (GCSIMP), the RuO2-doped one (Ru-GCSIMP), and a more complex representative "centroid" composition (denoted CX, see Table 1) were slow cooled.

Characterization

Thermal analysis

To assess the glass transition (Tg) and crystallization (Tx) temperatures, differential scanning calorimetry (DSC) was performed using a Netzsch STA 449F1. Experiments were performed with ~35-45 mg of crushed glasses, heating from 50 to 1200°C at 10°C/min, then subsequent cooling at 10°C/min to 50°C.

Viscosity

Low-temperature viscosity for the GC SIMP glass was measured from 585°C to ~660°C with a creep apparatus following the procedures in described by Neuville. 26 Glass samples were prepared by waterquenching glass melts from 1300°C, annealing from 500°C to room temperature over 16 hours, and subsequently core-drilling using a high-pressure water jet. Each viscosity datapoint corresponds to an average of between 15 and 40 measurements, made at the same temperature with different stresses, to check for Newtonian behavior of the glass. Expected accuracy is 0.03 log units of viscosity. Data were fit to the Vogel-Fulcher-Tammann (VFT) non-Arrhenius law, and the Tg was obtained. This Tg value, calculated for viscosity equal to 10 12 Paꞏs, is was expected to be lower than that obtained by DSC because of the different speed of heating.

High-temperature viscosity for the GC SIMP glass was measured from 1300°C to ~1050°C using the procedure described in Crum et al. 27 based on ASTM C 965-94. 28 Quenched glass (30 mL container) was loaded into a platinum crucible. Measurements were collected by usingwith an FRS1600 rheometer (Anton Paar GmbH, Austria) and a platinum spindle. The temperature program was set at a ramp rate of 8°C min - was not rotated during the hold times to avoid impacting any crystals and phase separation in the melts.

Electron microscopy

Scanning electron microscopy (SEM) was used to obtain micrographs on some of the samples after their soak and quench. A JEOL JSM-7001F field emission gun SEM (JEOL USA, Inc., Peabody, MA) with a backscattered electron detector was used for imaging.

Selected samples were investigated with scanning transmission and transmission electron microscopies (STEM/TEM). Here, heat-treated samples were pulverized in a Diamonite TM mortar and pestle and then thinly dispersed on a standard lacey carbon TEM grid. A JEOL ARM 200CF (JEOL USA, Inc., Peabody, MA) equipped with a high-angle annular dark-field (HAADF) detector and a Centurio silicon drift energy dispersive spectrometer (EDS) operated at 200 kV was used for the STEM/EDS analysis. The STEM was operated at 5 µA beam current with a 40-µm condenser aperture (CA) to minimize beam damage during imaging. EDS analysis was carried out at 15 µA with a 70-µm CA. Some degradation of the glass matrix was observed during imaging and EDS analysis, but crystals were found to be relatively stable. Pathfinder (v1.0, Thermo Fisher Scientific Inc.) software was used for recording and analysis of the EDS data.

Selected samples were investigated using electron probe microanalysis (EPMA) using a JEOL JXA-8500F field emission electron microprobe (JEOL USA, Inc., Peabody, MA) equipped with five wavelengthdispersive X-ray spectrometers (WDS), see Supplemental Information for measurement parameters.

Small angle X-ray scattering

Small angle neutron and X-ray scattering techniques were used to investigate the possible presence of phase separation in the soak-and-quench glasses. These are described below in more detail. At times, the dimensionality Dimensionality accessed in of these techniques is referred to in "q space", which is the scattering vector (reciprocal space) magnitude, i.e., q = (4πsinθ)/λ = 2π/d, where θ is the scattered angle, λ is the wavelength, and d is the interplanar spacing. Initially, small angle neutron scattering was attempted (see Supplementary), but small and ultra-small angle X-ray scattering (SAXS and USAXS, respectively) measurements proved more useful. Overall, one set of experiment allowed feature sizes to be probed from 4π/qmax ~ 9 Å to π/qmin ~ 10.5 μm, thus covering wide-angle X-ray scattering (WAXS), SAXS, and USAXS ranges.

Small angle neutron scattering

Small-angle neutron scattering (SANS) was obtained at room temperature at beamline CG-2 (Cold Guide #2), the general purpose small angle neutron scattering (GP-SANS) beamline at the High Flux Isotope Reactor (HFIR) at Oak Ridge National Laboratory (ORNL) in Oak Ridge, TN. [START_REF] Wignall | The 40 m general purpose small-angle neutron scattering instrument at Oak Ridge National Laboratory[END_REF] The twodimensional ~1 m 2 area detector was placed ~7.8 m away from the sample to accept small angles; single wavelength cold neutrons of 4.72 Å were used to perform the small angle scattering for "short" and "medium" q ranges, while 12 Å neutrons were used for the "long" range. These three q ranges (0.05658-1.030021 Å -1 , 0.007423-0.176563 Å -1 , and ~0.000409-0.028031 Å -1 ) were measured with some overlap to allow for data stitching. The approximate size of viewable objects was obtained (in nm) by π/qmin/10 to 4π/qmax/10, corresponding to objects ~1.22 -768.05 nm in diameter.

Approximately 0.5-1.0 g of coarse ground powder was placed within a titanium ring between two sapphire plates in an aluminum housing. The overall path length of the material tested was ~2 mm. Six samples were placed in a magazine that allowed for semi-automated data collection. Samples were not sieved, thus the packing fraction was uncontrolled. Collected data was processed using Igor Pro (WaveMetrics, Inc., Oswego, OR, USA), and q ranges were stitched for fitting.

Small angle X-ray scattering

To assess whether X-ray contrast would be better than neutron contrast for seeing scattering, aA subset of the quenched samples was also measured at the Advanced Photon Source (APS) at beamline 9 ID-C (setup described elsewhere 30,31 ) at Argonne National Laboratory (ANL) in Lemont, IL. A layer of powder ~0.1 mm thick of powder was placed on cellophane tape for X-ray scattering measurements. The incident X-ray energy was 21 keV (λ = 0.5904 Å), and the beam size was 0.8×1.2 mm. Beamline 9-ID-C is equipped with an advanced-design Bonse-Hart camera for ultra-small angle X-ray scattering (USAXS) analysis integrated with fixed-length pinhole SAXS (500 mm) cameras. 32 The camera system at beamline 9-ID-C is was capable of recording SAXS curves with an angular resolution of ~8×10 -4 Å -1 in scattering vector q range from 10 -4 to 1.2 Å -1 .. The USAXS instrument runs ran a slit-smeared Bonse-Hart setup and, in this experiment, the slit length was 0.027843 Å -1 . Slit smearing was included in data analysis models as needed. Each sample was analyzed for five minutes.

USAXS and SAXS data were completely corrected for instrumental variations, reduced, and analyzed with Igor Pro ® (v.7.0.5.2) coupled with Nika (v.1.74) and Irena (v.2.61) packages. 31,33 The samples were analyzed with the unified method developed by Beaucage and Schafer. 34,35 The unified method models the SAXS distribution as a hierarchical structure where the small structures build into larger structures. The unified equation for a hierarchical structure with multiple levels is given in Equation ( 1) 34 𝐼 𝑞 ∑ 𝐺 exp 𝐵 exp erf 𝑞𝑅 /6 / /𝑞

where n is the structure level, q is described above, Rgi is the radius of gyration of the i-th level, Pi is the power-law exponent, Gi is the Guinier prefactor, and Bi is the prefactor specific to the type of power-law scattering. All samples were analyzed using the unified method using two structural levels to specifically examine the radius of gyration (, Rg. ). From this model, Rg characterizes characterized the size of the scattered material regardless of shape. 36 However, Rg relates to the dimension of the scattered material differently depending on the shape of particle. For a sphere of radius r, Rg is defined as seen in Equation (2) 36 as an example.

Rg = 𝑟 ≈ 0.77*r (2)
Since the shape of the scattering material is unknown, Rg values are presented rather than r values of a predetermined shape. Additionally, the power-law slope parameter (P) describes the fractal dimension of the ensemble of scattering objects. By Porod's law, structures with smooth surfaces have P values near -4, rough surfaces have P values between -3 and -4, and mass fractals have P values between -1 and -3. 36 The G and B parameters are not reported as they change as intensity is modified.

Overall, the experiment allowed feature sizes to be probed from at 4π/qmax ~ 9 Å to π/qmin ~ 10.5 μm, thus covering wide-angle X-ray scattering (WAXS), small-angle X-ray scattering (SAXS), and ultra-small

X-ray dDiffraction

In addition to the synchrotron-obtained WAXS data on the soaked and quenched powders, other X-ray diffraction (XRD) data was were collected using laboratory diffractometers. All samples were prepared by crushing to fine powders using a Diamonite TM mortar and pestle and using a zero-background silicon holder due to small sample volumes.

Standard room -temperature powder X-ray diffraction (XRD) was obtained for all quenched samples (1150-1450°C soaks) on a PANalytical X'pert Pro MPD (the Netherlands) using a Cu Kα source and X'celerator TM detector. Rietveld refinement was performed using Highscore Plus (PANalytical) software.

For selected samples, powders were doped with 10 mass% Al2O3 internal standard for quantification of the amorphous fraction.

XRD was performed on three compositions of slow-cooled samples (1×SC) on a D8 Advance diffractometer (Bruker AXS Inc., Madison, WI) with a Cu Kα source and a LynxEye TM position-sensitive detector. The simplified composition (GCSIMP), the RuO2-doped version (Ru-GCSIMP), and the complex composition (CX) were all analyzed after 1×SC and mixing with a known quantity of TiO2 standard (SRM-674b; National Institute of Standards and Technology), using Rietveld refinement and the Topas (Bruker) software, in order to obtain the amorphous fraction. Errors on the quantification of phases are estimated at a maximum of ~0.1 × mass% for major phases (oxyapatite), and 0.25 × mass% for minor phases (powellite), as assessed by similar detailed studies using the same quantitative XRD methods. 37,38 Additionally, hot-stage XRD was obtained on the same D8 Advance diffractometer. A powder of the GC SIMP glass was suspended in ethanol, placed onto a Pt heater strip, and allowed to dry. Then, the sample was heated from 30°C to 975°C at ~10°C s -1 , held for 24 hrs, then cooled at ~10°C s -1 to 30°C, with scans collected at 30°C, at the beginning and end and at 975°C, and then every ~1 hr during the hold at . Data was analyzed using Topas (v4.2, Bruker) software using Rietveld refinement to fit the relative fractions of crystalline phases and obtain lattice parameters. Because of the nature of the experiment, the amorphous fraction could not be obtained. To estimate the total crystal fraction at the beginning and end of this experiment, a sample of GC SIMP was heated in a furnace to 975°C at ~10°C s -1 , held 2 min, and air quenched; a separate sample was heated to 975°C at ~10°C s -1 , held for 24 h, and air quenched. Samples were mixed with Al2O3 for quantification by XRD (PANalytical X'Pert Pro diffractometer) and Rietveld analysis (Highscore).

Optical absorption

Diffuse reflection was performed on powdered samples at room temperature using an Evolution™ 260

Bio UV-Visible Spectrophotometer (Thermo Fisher Scientific) with installed ISA-220 Integrating Sphere Accessory (Thermo Fisher Scientific). Measurements were performed from 350 nm to -1100 nm with 0.1s integration time, 2-nm bandwidth, and 1-nm step size; note that a (0.5-nm step size was used for a pure ceramic Ca2Nd8(SiO4)6O2 ("Ca-Nd oxyapatite") standard made by solid state synthesis described elsewhere. 24 ) The spectra in the range between 420-440 nm was were remeasured at 1-s integration time, 1-nm bandwidth, and 0.2-nm step size; the measurements were repeated five times and averaged. Diffuse reflectance was measured relative to a standard, calibrated Spectralon® (Labsphere, NH, USA) in this case.

Measured relative reflectance (𝑟 ) was then converted to the Kubelka-Munk (K-M) function [𝑓 𝑟 ] using Equation ( 31)

𝑓 𝑟 * ( 31
)
where 𝐾 is the absorption coefficient and S is the scattering coefficient. 39 The K-M will simply be referred to as absorption in this paper.

23 Na nuclear magnetic resonance

A subset of GC SIMP samples were subjected to magic angle spinning (MAS) nuclear magnetic resonance (NMR) measurements. The 23 Na MAS NMR measurements were acquired on a 14.1 T (B0) (δiso = 7 ppm with respect to NaCl(aq) 0.1 M, δiso = 0 ppm). To compare the Na environment in typically expected crystals in the glass-ceramic, single pulse experiments were conducted on commercially available Na2MoO4 andas well as in-house produced Na0.5La0.5MoO4 powellite and NaLa9(SiO4)6O2, which were synthesized using a previously described process (see Supplementary Information). 24 43 Results and discussion

Thermal analysis

The DSC spectra, after background removal, are shown in Figure S1. Positions of the crystallization peaks show that, on heating, the RuO2-doped sample crystallized at lower temperature than the undoped sample, and has an additional higher temperature crystallization peak as well. Measured temperatures from the heating experiment are summarized in Table 2Table 2. On heating, there was evidence of two crystallization peaks for both glasses, while on cooling, there was evidence of only one crystallization peak, at lower temperature than for heating.

Viscosity

A summary of the low-temperature and high-temperature viscosities is provided in Table 3Table 3.

The whole set of data shown in Table 3Table 3 was fit using the Vogel-Fulcher-Tammann (VFT) approach as follows:

log  = A + B/(T -T1) ( 42
)
where A, B, and T1 were determined to be -4.4073±0.514, 4972.3±589, and 565.16±28, respectively, using Igor Pro (v6.37) and the Tg was determined as 595.1°C 868.2 K (595.1°C868.2 K) at 10 12 Paꞏs. The data is plotted in Figure 1Figure 1. A high yield stress is observed in the 1250--1150°C range, which could be related to active phase separation and crystallization. The viscosity behaves normally between 1100-and 1050°C.

Low temperature viscosities, around Tg, were measured with different strains, which confirm a Newtonian behavior and prove that the liquid is crystal-free. 26 For hHigh -temperature viscosities at 1300, 1250, and 1200°C, dispersions are Newtonian, but those at 1150, 1100, and 1050°C are not, suggesting the presence of phase separation and/or crystallites in the liquid. This can be seen in the larger error bars of the 1150--1050°C viscosity measurements as well. The crystal effect noted in the high-temperature viscosity measurements are clearly visible in the SEM images in Figure 2. increased, in agreement with XRD results (see below). Spherical molybdate domains were found in each sample, with the locations being isolated to the perimeter of oxyapatite crystals at T ≥ 1225°C; this effect is most readily clearly seen in Figure 2d. At these temperatures, the removal of REs from the melt by crystallization destabilized the melt and resulted in the phase separation of spherical Mo-rich droplets. In general, these phase separation features appeared much less pronounced in this simplified glass compared to the complex glass-ceramic compositions examined previously. 1,2 Final microstructures of heat -treated samples show modifications of the microstructure from the soak-and-quench samples. The hot stage XRD isothermally heat-treated sample from the hot-stage XRD showed similarly sized oxyapatite crystals, but in much larger concentrations (Figure 3a), while the slow-cooled GC SIMP sample showed larger, but not as concentration concentrated, oxyapatite crystals (Figure 3b).

TEM micrographs were obtained for water -quenched samples from 1200, 1250, and 1300°C; some of this data is shown in Figure 4. TEM-EDS maps and line scans showed that these samples are were enriched in Ca and Mo and depleted in Si relative to the matrix (see Figure 4d). Morphologically, these spheres looked more like "partial moons," since the volume of the liquid shrank on cooling more extensively than the glass, and the calcium molybdate crystals occupied a smaller volume than that of the spherical liquid phase. This phenomenon had also been observed in previous studies. 41 Small angle scatteringThe spectra were compared, with USAXS and SAXS spectra ( are shown in Figure 5a) . compared to the SANS spectra (Figure 5b) for the same samples (see Figure S4 for the ~0.01-0.1 Å -1 SANS spectra for all samples Supplementary Information for small angle neutron scattering data).

For the SANS data, the low q region showed a flattening of intensity for all samples, and was likely a result of multiple scattering events due to thick samples. Both SANS and The SAXS/USAXS data may have had a scaling problem. , as It it could not be determined for certain whether the low q values should have lain on top of one another, in which case the intensity differences were due to slightly different thicknesses or packing factors of the samples (e.g., mounting inconsistencies and particle size differences). However, the intensity differences do not affect the analysis of the scatterer sizes of the scatters as the Guinier region will remain the same regardless of intensity. However, for the USAXS data, the intensities appeared to scale with Tsoak, with the 1175°C sample showing the most scattering throughout the majority of the available q range, consistent with a higher density of interfaces. Qualitative locations for the Guinier region (indicating size of the scatterer) showed that the size (radius of gyration) increased as Tsoak decreased. Quantitative estimates for Rg in Table 4Table 4 showed a decreasing trend in Rg values as soak temperature increased from 1175 to 1250°C. and Rg for 1175, 1200, and 1250°C (see Figure 4).

All Porod P values in the q region directly higher than Rg for 1175-1250°C samples were about -3, indicating rough surfaces for the CaMoO4 scatterers. The 1300 and 1350°C samples did not show increased scattering in the same q region, rather increased scattering is seen at 0.03-0.2 Å -1 . The Rg values for 1300 and 1350°C samples were overly low and unphysical. , and were probably due to gaps between individual particles, considering that the samples were powders as solid coupons were not obtainable using the quench procedure. The other samples likely had the same scattering, but it was overshadowed by scattering due to CaMoO4.

Diffraction

WAXS spectra (Figure 65b) obtained at the same time as the USAXS/SAXS showed increase of intensity from q ~1.2 to ~2.2 Å -1 as was also observed in the SAXS data at high q (Figure 5a). Diffraction peaks were visible in the WAXS of 1175, 1200, and 1250°C samples, and some peaks could be observed in the SAXS of 1175 and 1250°C samples, particularly at q = 0.7714, 1.53838, and 1.63355 Å -1 .

Quantitative fits were obtained by Rietveld refinements of the WAXS (crystalline fraction only) and corresponding room temperature XRD (quantified amorphous phase) as shown in Table 5Table 5. For completeness, all the samples that were cooled from the melt, soaked, then water quenched were subjected to room temperature XRD (Cu X-ray tube) (see Figure 76). The combined data showed that the 1250°C sample was the highest temperature measured with obvious Bragg peaks. As seen in Table 5Table 5, it can be seen that crystalcrystal fraction increases increased only slightly as the soak temperature is was decreased. Imaging and chemical analysis of the phases on the Pt heater strip after hot-stage XRD were attempted.

The inability to polish the sample, as well as the small size of the crystals, made obtaining high-quality data difficult. Mo-rich (powellite) crystals were present but were too small (< 0.5 μm) to adequately quantify with the 1-μm beam size of the EPMA. Oxyapatite crystals were likewise small, ~1-μm long and ~0.1-μm in diameter (see Figure 2). To further assess the effect of this different heat treatment, the ex-situ experiments were performed for essentially a t = 0 and t = 24 point for crystallization at 975°C. The conventional XRD results of powdered samples showed a high degree of crystallization at 24 h (~50 The crystal fractions obtained by ex-situ XRD on the soak-and-quench glasses was were 2--6 mass%, with the difference being amorphous residual glass. These crystal fractions were quite low compared to typical slow-cool examples, where crystal fractions is were typically 24-36 mass%. 2 To directly compare the crystal fraction and assess compositional effects, 1×SC tests were performed on the simplified GC SIMP composition, the Ru-GCSIMP, and the CX complex composition. The crystal fraction was nearly the same in the CX, GCSIMP, and Ru-GC SIMP samples, so there was no apparent effect of composition, and the main effect appeared to be the slow cooling.

As discussed in a separate paper, 45 the processes for acquiring reliable and representative diffraction data on glass-ceramic samples using ex-situ and/or in-situ techniques have limitations. The limitations with heat treating samples in sealed capsules include melt-capsule reactions (e.g., if silica, sapphire) or issues with beam attenuation from the capsule wall (e.g., platinum). Additionally, it is difficult to monitor the amorphous fraction for in-situ measurements since a diffraction standard cannot be introduced into the scan data; thus, crystallographic phase distributions can be extracted but quantification is most likely impossible for in-situ diffraction measurement techniques. However, while ex-situ is an option, analysis requires specimens to be cooled prior to analysis, and forcing specimens to undergo cooling results in microstructural changes during the cool-down process and introduces artifacts into the data. The primary artifacts are in the correlations drawn between specimens undergoing cooling and those that already underwent cooling.

Optical absorption and NMR

Absorption (K-M) spectra were obtained for selected samples and compared to a single-phase ceramic Ca-Nd oxyapatite (Figure 98). Only the Nd 3+ 4f-4f electronic transitions were seen, and the absorption peaks were relatively broad and convoluted in the 1350°C soak sample. But, as the thermal soak

F o r P e e r R e v i e w
25 temperature decreased, the peaks appeared sharper and more distinct, showing more resemblance to the features in single-phase Ca-Nd oxyapatite ceramic, indicating that more Nd was in crystalline phase(s). The sharpness of absorption features was best seen in the 4 I9/2 → 4 G9/2, 4 G7/2, and 2 K13/2 (~520 nm); 4 I9/2 → 4 G5/2 and 2 G7/2 (~580 nm); 4 I9/2 → 4 S3/2 and 4 F7/2 (~745 nm); and 4 I9/2 → 4 F3/2 (~871 nm) transitions. 46,47 The transitions of 4 I9/2 → 4 G5/2, known as the hypersensitive transition due to its sensitivitiesy to changes in the Nd 3+ crystal field and symmetry, 13,[START_REF] Quintas | Effect of alkali and alkalineearth cations on the neodymium environment in a rare-earth rich aluminoborosilicate glass[END_REF]49 was of particular interest; however, not much change in the peak position or the shape of that peak was seen in samples soaked between 1350°C and 1250°C, indicating no significant changes in the local environment of the Nd 3+ ion. Below 1250°C soak temperature, the peaks became sharper, likely due to the crystallization of Nd-containing phases. More details, especially regarding the 4 I9/2 → 2 P1/2, are provided in the Supplementary Information. The 23 Na MAS NMR showed the presence of a distinct sharp peak when the soak temperature was 1150°C (Figure 109). This sharp peak was on top of the broad background peak of Na in the glass, which is was evident in all the other samples. The sharp peak is located at -0.7 ppm, and is indicative of Na in a crystalline phase. It is probable that Probably this Na was in oxyapatite for the following reasons described hereafter. The amount of oxyapatite in the glass-ceramic increased with decreasing soak temperature (Table 5Table 5), while the amount of powellite remained about the same. While the 1150°C was not quantified, the oxyapatite concentration should be somewhat more than the quantified 1175°C sample if it followed the trends at higher soak temperatures. Alternatively, the lower soak temperature could have allowed Na to enter the oxyapatite and its composition was somewhat different than at higher temperatures. This was suggested by the hot-stage XRD data where two different oxyapatites (based on differing unit cell parameters) were observed, but it was possible that the phase changes in the hot stage experiment were due to sample composition change during volatilization. Regardless, it is known that Na can go into the Ca-Nd oxyapatite or form a Na-Nd oxyapatite on its own. 50 The SEM-EDS investigations were not helpful here, as the Na and Nd X-ray lines overlap somewhat (i.e., Na Kα = 1.040 keV and Nd Mα = 0.978 keV); however this was observed in the EPMA data of the 1×SC GC SIMP sample (see Supplemental Information).

Synthesis and 23 Na measurement of standards, Na0.5La0.5MoO4 powellite and NaLa9(SiO4)6O2 apatite, confirmed that the most likely assignment of the 23 Na peak in the glass-ceramics is was Na in oxyapatite.

Note that, for the standards, La was used as the rare earth to avoid any possible problems with paramagnetic Nd, but the Na environment should be very close to the the same in the Nd-containing oxyapatite. The NMR should showed that it is was very unlikely that the Na signature was from Na-containing powellite or Na2MoO4. To further explore the identity of the crystalline phase containing Na, T1 measurements were conducted to investigate the relaxation time of the glass versus the crystalline phase. The T1 of the narrow crystalline phase seen in the 23 Na MAS NMR is an order of magnitude less than that of the broader glass phase (0.13 ± 0.02 s the glass versus 0.011 ± 0.001 s for the crystalline phase, see Supplemental Information

Summary and Conclusions

Various heat treatment protocols were used to study phase separation and crystallization of a simplified glass-ceramic composition, being investigated as a waste form for aqueous wastes generated during used Slow cooling for the GC SIMP sample produced a significantly larger fraction of crystals (~30 mass%).

It appears that the powellite phase separation happens upon cooling at quite high temperatures, corresponding to an estimated viscosity of ~0.04 Paꞏs per the VFT model, well above the glass transition temperature (Tg ~ 609°C). This study paper shows that crystallization studies using nucleation and growth, such as the quenching and heat treating at constant high temperature of 975°C, do not produce a microstructure and crystal fraction representative of melts that are slow-cooled and crystallized from high temperature. In the future, these results should be confirmed by in-situ diffraction from high temperatures in a sealed container, so that volatilization does not cause compositional changes and different phase evolution as suggested by laboratory hot-stage XRD. isare to contain volatile samples during the experiment, and identifying a sample container that is passive enough not to interfere with the experiment, i.e., it does not react with the melt and does not (significantly) attenuate the interrogation beam. Regarding ex-situ analysis, if samples could be cooled quickly enough, it is possible that they could be interrogated with ex-situ analysis and the microstructure would be representative of the high -temperature structure. This was the overall approach of the current study. Solving the in-situ diffraction issues would open the possibilities of characterizing a wide range of waste-form materials for optimization of microstructure, chemistry, and phase distribution, and chemical durability.

From a waste form perspective, controlling crystallization in a predicable way, using primarily compositional control, is highly beneficial when the precipitated crystals are chemically durable. The durability of oxyapatite and powellite crystalline phases, as well as the nominal residual glass composition, has been recently demonstrated. 24,51 For conventional borosilicate waste glasses, crystallization is generally unwanted, as it usually results in melter processing problems or poor durability of the residual glass. 52,53 Defense waste glasses (such as those planned for Hanford) are so compositionally diverse 54 that it would be difficult to formulate compositions for making glass-ceramic waste forms with predictable crystalline phases that do not cause either problems in melter processing or in product performance due to residual glass durability. For the particular fuel reprocessing waste stream considered here, the problem is easier because the phases form upon cooling so they do not cause melter processing problems, the phases that form are chemically durable, and the residual glass is also durable. What The primary gap that remains with this glass-ceramic waste form is to refine the understanding so that control of microstructure, as well as and phase distribution, as a function of cooling rate and composition, which can be achieved in order to eliminate any interface stress-induced chemical durability issues induced by interfacial stresses between the glass and crystalline phases. Figure 109. 23 Na MAS NMR spectra of SIMP. The sharp peak in the 1150°C sample is located at -0.7 ppm. The Na-containing oxyapatite peaks at 1.5 ppm, the Na-containing powellite has a quadrupolar lineshape peaking around 12.5 ppm, and the Na-molybdate has a complex lineshape with the main peak around -1 ppm. Note that at 1150, 1100, 1050°C, the melt is not Newtonian. The melt is more like a Bingham plastic, meaning there is a yield stress that must be overcome to cause flow (non-zero intercept) followed by a linear response to increasing shear rate. This behavior is explained by the presence of crystals. The crystals resist flow until the yield stress is overcome, then they probably realign with the direction of flow, and it shows Newtonian behavior with increasing shear rate. The key is the intercept or offset. Small-angle neutron scattering (SANS) was obtained at room temperature at beamline CG-2 (Cold Guide #2), the general purpose small angle neutron scattering (GP-SANS) beamline at the High Flux Isotope Reactor (HFIR) at Oak Ridge National Laboratory (ORNL) in Oak Ridge, TN. S1 The twodimensional ~1 m 2 area detector was placed ~7.8 m away from the sample to accept small angles; single wavelength cold neutrons of 4.72 Å were used to perform the small angle scattering for "short" and "medium" q ranges, while 12 Å neutrons were used for the "long" range. These three q ranges (0.05658-1.030021 Å -1 , 0.007423-0.176563 Å -1 , and ~0.000409-0.028031 Å -1 ) were measured with some overlap to allow for data stitching. The approximate size of viewable objects was obtained (in nm) by π/q min /10 to 4π/q max /10, corresponding to objects ~1.22 -768.05 nm in diameter.

Approximately 0.5-1.0 g of coarse ground powder was placed within a titanium ring between two sapphire plates in an aluminum housing. The overall path length of the material tested was ~2 mm. Six samples were placed in a magazine that allowed for semi-automated data collection. Samples were not sieved, thus the packing fraction was uncontrolled. Collected data was processed using Igor Pro (WaveMetrics, Inc., Oswego, OR, USA), and q ranges were stitched for fitting.

For the SANS data, the low q region showed a flattening of intensity for all samples, and was likely a result of multiple scattering events due to thick samples.

TEM

Imaging was performed, and relative concentration profiles were obtained by line scans using Energy Dispersive Spectroscopy (EDS). Absolute concentrations were not obtainable due to the dissolution of particles under the electron beam. A few maps were obtainable.

XRD

The soak-and-quench crystal fractions were quite low compared to the to- is not observed in any tested compositions, S2 nor was it observed here. However, since the simplified composition in the current study was based on the GC-5.86 composition, S2 it was initially expected that some lanthanide borosilicate would be observed, but it was not. It was unclear whether this was due to the borosilicate phase normally forming at longer times or at slower cooling rates or if there was a subtle composition effect that is not reproduced in the simplified glass.

It may be that the long cooling time is necessary to develop the larger crystal fractions; however, this possibility is counter to the results for the composition studied previously as a function of cooling rate (GC-6.25-Li), where crystal fraction varied very little (~20-23% total) when modifying cooling rate from 0.25×SC-4×SC. S3 In that study, the main difference in results for different cooling rates was the size and morphology of the crystals, and to some extent the type of crystal, while the overall fraction remained quite constant.

Rietveld refinements on repeated centroid glasses

Na-La Powellite

Sodium-lanthanum powellite (Na 0.5 La 0.5 MoO 4 ) was synthesized in an alumina crucible using Na 2 CO 3 , La 2 O 3 , and MoO 3 , the masses are listed in the Table below. Initially the powders were hand mixed and then fired in an alumina crucible using the following heat treatment: ramp to 1500°C at 5°C/min, hold at 1500°C for 0.5hr, then ramp to 1400°C at 1°C/min where the sample was held at 1400°C for 1hr and then finally ramped back down to room temperature at 1°C/min. The sample was chipped from the crucible, milled in a tungsten carbide mill for 30second, and then re-melted in a different alumina crucible with the same heat treatment as above. The sample was then chipped from this second alumina crucible. Some of the sample was milled by hand in a diamonite mortar and pestle and some of the powder was used for powder XRD.

Na-La Powellite oxide and carbonate masses (in g) for batching

Na-La Powellite Na 2 CO 3 2.200 La 2 O 3 6.762 MoO 3 11.951

During synthesis, the sample did wet the alumina crucible walls, and the sample had to be chipped from the walls of the crucible. It is possible during this process that some small alumina pieces got into the sample, although the alumina pieces were separated visually as carefully as possible. Pictures of the first and second heat treatment are shown below along with a view of one of the crystals in the alumina crucible after the first heat treatment.

Na-La Apatite

A Na-containing oxyapatite was synthesized with a solution based process, NaLa 9 (SiO 4 ) 6 O 2 . The Table below gives the masses and volumes used for synthesis. Synthesis was performed by weighing out the nitrates and then dissolving the nitrates by adding EtOH and acetic acid. After mixing this solution for 24 hr, TEOS was added and mixed for 1 hr. This solution was then placed into a drying oven at 80°C for 3 days. In an alumina crucible the samples were heat treated at 600°C for 4 hr, and then the sample was milled by hand in a diamonite mortar and pestle. The sample was then heat treated at 1000°C for 2 hr, followed by a 1550°C heat treatment for 8 hr. The ramp rates were 5°C/min for the 600°C and 1000°C heat treatments, while the 1550°C heat treatment had ramp rates up and down of 1°C/min.

Before the 1550°C heat treatment, a pellet the oxyapatite was prepared. This pellet was treated at the same time and in the same furnace as the powder for the 1550°C heat treatment. The Na-La pellet and powder had some light brown "specks" or "dots" on the surface after the heat treatment shown.. Also, for both samples, the powder seemed to have shrunk together some by pulling away from the walls, although the sample was still very powdery when removed. Pictures showing the whole process are below. The XRD data for Na-La oxyapatite seems to have no secondary phases shown by XRD. The unit cell parameters and space group for Na-La oxyapatite and Ca-Yb oxyapatite is shown below.

Masses and Volumes for batching of Oxyapatites

Unit cell parameter and space group for

NMR

The fits to the 23 Na T1 data are shown below. The T1's of the two peaks were calculated by fitting saturation recovery data, as can be seen in the plots. The values found for the two phases were T1(crystalline) = (0.011 ± 0.001) s and T1(glass) = (0.13 ± 0.02) s. Therefore, the crystalline phase was found to relax a factor of 10 times faster. The function used in the fitting was I(t) = I0(1 -e -t/T1 ).

Post hot-stage XRD

One sample, that created after hot-stage X-ray diffraction (XRD) (see below) was investigated using electron probe microanalysis (EPMA) using a JEOL JXA-8500F field emission electron microprobe (JEOL USA, Inc., Peabody, MA) equipped with five wavelength-dispersive X-ray spectrometers (WDS). Because the sample was a crystallized glass on the platinum heater strip from the hot-stage XRD, the sample could not be polished and imaging conditions were not ideal.

  The target crystalline phase assemblage includes powellite [(RE,A,AE)MoO 4 ], oxyapatite [(RE,A,AE) 10 (SiO 4 ) 6 O 2 ], cerianite [(Ce,Zr)O 2 ], and sometimes lanthanide borosilicate (RE 3 BSi 2 O 10 ), pollucite (CsAlSi 2 O 6 ), and/or celsian [AEAl 2 Si 2 O 8 ].

  reflectance ( ) was then converted to the Kubelka-Munk

(

  NMR) measurements. The23 Na MAS NMR measurements were acquired on a 14.1 T (B 0 ) Bruker Avance II+ Spectrometer operating at a 23 Na Larmor frequency (ν 0 ) of 158.6 MHz. A Bruker 3.2-mm probe was used to provide 20 kHz MAS. A secondary solid reference NaCl(s) (δ iso = 7 ppm with respect to NaCl(aq) 0.1 M, δ iso = 0 ppm) was used for the23 Na chemical shift and pulse time calibration. A non-selective (solution) π/2 pulse of 12 μs was measured on solid NaCl which represented a selective (solid) π/2 pulse of 6 μs. For all23 Na MAS NMR measurements, a selective π/6 pulse of 2 μs was implemented to achieve quantitative conditions for the excitation of the central (½, -½) transition. For single pulse experiments, 2000 co-added transients were used. The23 Na T 1 measurements were performed using a saturationrecovery pulse sequence which employed a saturation train of 300 π/2 pulses for the saturation of the magnetization. All acquired data were simulated using the QuadFit 40 software package, and the integrated intensities were plotted against the recovery delays to elucidate the individual T 1 values. To compare the Na environment in typically expected crystals in the glass-ceramic, single pulse experiments were conducted on commercially available Na 2 MoO 4 and in-house produced Na 0.5 La 0.5 MoO 4 powellite and NaLa 9 (SiO 4 ) 6 O 2 , which were synthesized using a previously described process (

  few representative SEM micrographs obtained on the glass-ceramics are shown in Figure2.

  Guinier region (indicating size of the scatterer) showed that the size (radius of gyration)

  due to the crystallization of Nd-containing phases. More details, especially regarding the 4 I 9/2

  the SIMP sample produced a significantly larger fraction of crystals (~30 mass%). It appears that the powellite phase separation happens upon cooling at quite high temperatures, corresponding to an estimated viscosity of ~0.04 Pa•s per the VFT model, well above the glass transition temperature (T g ~ 609°C

Figure 1 .

 1 Figure 1. Viscosity of SIMP sample from 588°C -1300°C (861 -1573 K) shown as log(η, Pa•s) as a function of both inverse temperature (K -1 ).

Figure 2 .

 2 Figure 2. Comparison of SEM micrographs at (a-d) 1 k× magnification and (e-h) ≥ 10-25 k× magnifications for samples water quenched after heat treatments at (a,e) 1175°C, (b,f) 1200°C, (c,g) 1225°C, and (d,h) 1250°C. Scalebars are provided for each sample.

Figure 3 .

 3 Figure 3. SEM micrographs of (a) hot-stage XRD sample after 975°C for 24 hrs and (b) 1×SC slow-cooled sample.

Figure 4 .

 4 Figure 4. (a) STEM HAADF micrograph of 1250°C glass shard with lacy carbon grid in the background and other shard or Nd-rich area on top; (b) close-up showing spherical Mo-rich regions showing darker areas from cooling and crystallization of the molybdates. (c) STEM HAADF image of 1200 °C glass shard and (d) STEM-EDS maps of molybdate particles in the marked region in (c).

Figure 5 .

 5 Figure 5. (a) SAXS/USAXS data; arrows in the SAXS/USAXS plot show the approximate position of the Guinier region indicating particle size, which becomes larger as T soak decreases; (b) Room temperature WAXS of different samples (Tsoak = 1175-1350°C), normalized to the peak intensity and offset for clarity

Figure 6 .

 6 Figure 6. Room temperature standard XRD of SIMP quench series, normalized to the peak and offset for clarity

Figure 7 .

 7 Figure 7. Hot stage XRD of SIMP; (a) time dependence of diffraction pattern; (b) crystalline phase distribution as a function of time with Pt normalized out (see text); hours "1" and "26" are at room temperature, while all others are at 975°C; lattice parameters (c) a-axis and (d) c-axis of the two oxyapatite phases identified.

Figure 8 .

 8 Figure 8. Optical absorption spectrum of selected soak-and-quench SIMP glasses compared to Ca-Nd oxyapatite ceramic standard showing transitions of the Nd 3+ ion. Shown is the full collected spectrum along with close-ups of some hypersensitive transitions.

Figure 1 .Figure 2 .Figure 3 .Figure 4 .Figure 5 .Figure 6 .Figure 7 .Figure 8 .Figure 9 .

 123456789 Figure 1. Viscosity of SIMP sample from 588°C -1300°C (861 -1573 K) shown as log(η, Pa•s) as a function of both inverse temperature (K-1).125x97mm (300 x 300 DPI)

  temperature was decreased in 50°C increments to the final temperature of 1050°C. The spindle

32

 32

  ray scattering (USAXS) ranges. For modeling the SAXS and USAXS data, the matrix was assumed to have the composition of the overall glass, and the scatterers were composed of spherically shaped CaMoO4 of spherical shape, as indicated by TEM (see Results).

  Spectrometer operating at a23 Na Larmor frequency (ν0) of 158.6 MHz. A Bruker 3.2mm probe was used to provide 20 kHz MAS. The experiments were calibrated on theA secondary solid reference NaCl(s) (δiso = 7 ppm with respect to NaCl(aq) 0.1 M, δiso = 0 ppm) was used for the23 Na chemical shift and pulse time calibration. A non-selective (solution) π/2 pulse of 12 μs was measured on solid NaCl which represented a selective (solid) π/2 pulse of 6 μs. For all23 Na MAS NMR measurements, a selective π/6 pulse of 2 μs was implemented to achieve quantitative conditions for the excitation of the central (½, -½) transition. The single pulse experiment consisted of a 'non-selective' (solid) π/3 pulse of 2.0 μs. For eachsingle pulse experiments, 2000 co-added transients were used in each experiment.The23 Na T1 measurements were performed using a saturation-recovery pulse sequence which employed a saturation train of 300 π/2 pulses for the saturation of the magnetization. All acquired data were simulated using the QuadFit 40 software package, and the integrated intensities were plotted against the recovery delays to elucidate the individual T1 values. The experiments were calibrated on the secondary reference NaCl(s)

Formatted: Font color: Accent 5 F 13 Table 2 .

 5132 Critical values from thermal analysis and viscosity measurements. All parameters are from DSC heating measurements except as noted. Extrapolated Tg from low temperature viscosity and VFT equation (b) DSC on cooling

Formatted: Font color: Accent 5 Formatted

 5

Figure 1 . 2 .

 12 Figure 1. Viscosity of GC SIMP sample from 588°C -to through 1300°C (861 -1573 K) shown as log(η, Paꞏs) as a function of both inverse temperature (K -1 ).

16 Figure 2 .

 162 Figure 2. Comparison of SEM micrographs at (a-d) 1 k× magnification and (e-h) ≥ 10-25 k× magnifications for samples water quenched after heat treatments at(a,e) 1175°C, (b,f) 1200°C, (c,g) 1225°C, and (d,h) 1250°C. Scalebars are provided for each sample.

Figure 3 . 17 Figure 4 .

 3174 Figure 3. SEM micrographs of (a) hot-stage XRD sample after 975°C for 24 hrs and (b) 1×SC slowcooled sample.

Figure 5 .

 5 Figure 5. (a) SAXS/USAXS data; and (b) SANS data for the same 5 samples, with q ranges aligned. Aarrows in the SAXS/USAXS plot show the approximate position of the Guinier region indicating particle size, which becomes larger as Tsoak decreases.; (b) Room temperature WAXS of different samples (Tsoak = 1175-1350°C), normalized to the peak intensity and offset for clarity

21 Figure 6 .

 216 Figure 6. Room temperature WAXS of different samples (Tsoak = 1175-1350°C), normalized to the peak intensity and offset for clarity

22 Figure 76 .

 2276 Figure 76. Room-temperature standard XRD of GC SIMP quench series, normalized to the peak and offset for clarity. This data was taken prior to introduction of alumina standard used for quantification.

Figure 8a 7a shows

  Figure8a7a shows the evolution of the hot-stage XRD pattern versus time, where peaks become sharper with time at constant temperature of 975 °C. Only powellite and oxyapatite were identified, but two types of each phase were found with unique unit cell parameters. The powellite types fit well with Nd0.67MoO4 (ICSD 23675, 42 I41/a space group) and CaMoO4 (ICSD 62219, 43 I41/a space group) and the oxyapatites were both Ca-Nd oxyapatites (ICSD 62174,44 P63/m space group) but with different lattice parameters. The Pt sample holder/heating strip was identified in most scans and normalized out in the crystalline phase distribution graph (Figure8b7b). For the room temperature scan at the beginning of the experiment, the only phase present was Ca-powellite, but Nd-powellite and oxyapatite-1 formed within the first hour at 975°C, with oxyapatite-2 forming beginning at hour 5. The relative crystalline fractions of powellite continued to decrease until ~8 h when it became constant at ~10 mass%. The oxyapatite is was the majority phase. Lattice parameters of the oxyapatite phases were observed to decrease as time

Figure 87 .

 87 Figure 87. Hot-stage XRD of SIMP; (a) time dependence of diffraction pattern; (b) crystalline phase distribution as a function of time with Pt normalized out (see text); hours "1" and "26" are at room temperature, while all others are at 975°C; lattice parameters (c) a-axis and (d) c-axis of the two oxyapatite phases identified.

  , which was very close to that determined independently on the hot-stage XRD sample by image analysis using ImageJ software (~59 wtmass% amorphous).

Figure 98 .

 98 Figure 98. Optical absorption spectrum of selected soak-and-quench GC SIMP glasses compared to Ca-Nd oxyapatite ceramic standard showing transitions of the Nd 3+ ion. Shown is the full collected spectrum along with close-ups of some hypersensitive transitions.

Formatted:Formatted

  Justified, Indent: First line: 0.25", Line spacing: Double . Differences in the T1 values are likely due to the proximity of the paramagnetic Nd nuclei, with closer sites having shortened T1 values due to the paramagnetic interaction. This result suggested that the Na-bearing crystalline phase contained Nd, which was far more likely for an oxyapatite phase than for a Na molybdate phase.

Figure 109 .

 109 Figure 109. 23 Na MAS NMR spectra of SIMP. The sharp peak in the 1150°C sample is located at -0.7 ppm. The Na-containing oxyapatite peaks at 1.5 ppm, the Na-containing powellite has a quadrupolar lineshape peaking around 12.5 ppm, and the Na-molybdate has a complex lineshape with the main peak around -1 ppm.

  . The techniques used during this study included scanning, scanning transmission, and transmission electron microscopies; small angle neutron scattering; wide-angle, small-angle, and ultrasmall angle X-ray scattering; room temperature and hot-stage X-ray diffraction; optical absorption; and23 Na MAS NMR.Taken together, the results for the GC SIMP sample suggest that the 1300--1250°C temperature range is the critical soak temperature where amorphous melts can be maintained and, if quenched fast enough, can yield a product without phase separation and crystallization in this composition. Soak temperatures above this threshold maintain relatively homogeneous glass upon quenching (per TEM) and no evident crystallization (by XRD and WAXS). SAXS analysis for 1350°C and 1300°C show Guinier Rg values < 2 nm, suggesting there are no significant scattering phases in these samples at room temperature. The 1250°C soak sample showed spherical (~20-nm diameter) calcium molybdates in TEM, and small fractions of powellite and oxyapatite (~0.5 and ~1.0 mass%, respectively) in XRD and WAXS, superimposed on a broad glassy background. SAXS-obtained Rg values of the 1250°C sample are 7.3 nm (radius), which correspond well to the ~20 nm particles observed in TEM in this sample. At Tsoak ≤ 1225°C, quenched samples were visibly cloudy, and showed increasing fractions of crystals, although still very small amounts (~6 mass% total at 1175°C), but large enough to cause evident changes in the optical absor ptions spectra of Nd 3+ .

  the goal of the current work was to track the microstructural evolution in glass-ceramic samples, several data gaps still exist. The main issues for in-situ diffraction analysis (like hot-stage XRD)

Figure 61 .

 61 Figure captionsFigure 61. Viscosity of GC SIMP sample from 588°C -through 1300°C (861 -1573 K) shown as log(η, Paꞏs) as a function of both inverse temperature (K -1 ).

Figure 72 .

 72 Figure 72. Comparison of SEM micrographs at (a-d) 1 k× magnification and (e-h) ≥ 10-25 k× magnifications for samples water quenched after heat treatments at (a,e) 1175°C, (b,f) 1200°C, (c,g) 1225°C, and (d,h) 1250°C. Scalebars are provided for each sample.

Figure 83 .

 83 Figure 83. SEM micrographs of (a) hot-stage XRD sample after 975°C for 24 hrs and (b) 1×SC slowcooled sample.

Figure 94 .

 94 Figure 94. (a) STEM HAADF micrograph of 1250 °C glass shard with lacy carbon grid in the background and other shard or Nd-rich area on top; (b) close-up showing spherical Mo-rich regions showing darker areas from cooling and crystallization of the molybdates. (c) STEM HAADF image of 1200 °C glass shard and (d) STEM-EDS maps of molybdate particles in the marked region in (c).

Figure 5 .Figure 6 .

 56 Figure 5. (a) SAXS/USAXS data and (b) SANS data for the same 5 samples, with q ranges aligned. Arrows in the SAXS/USAXS plot show the approximate position of the Guinier region indicating particle size, which becomes larger as Tsoak decreases. Figure 10. (a) SAXS/USAXS data; and (b) SANS data for the same 5 samples, with q ranges aligned. Aarrows in the SAXS/USAXS plot show the approximate position of the Guinier region indicating particle size, which becomes larger as Tsoak decreases.; (b) Room temperature WAXS of different samples (Tsoak = 1175-1350°C), normalized to the peak intensity and offset for clarity

Figure 98 .

 98 Figure 98. Optical absorption spectrum of selected soak-and-quench GC SIMP glasses compared to Ca-Nd oxyapatite ceramic standard showing transitions of the Nd 3+ ion. Shown is the full collected spectrum along with close-ups of some hypersensitive transitions.

Figure S1- 1 .

 1 Figure S1-1. Thermal analysis on (a) heating and (b) cooling of simplified GC and RuO 2 -doped GC. Backgrounds are removed and power values are normalized and offset for clarity. Large peak evident on cooling is T g .

Figure S1- 2 .

 2 Figure S1-2. Thermal analysis raw data showing TGA and volatile losses for heating of SIMP glass.

Figure S2 .

 S2 Figure S2. Viscosity of GC sample from 1050°C through 1300°C shown as ln(η, Paꞏs) as a function of both (a) temperature (K) and (b) inverse temperature (K -1 ). In (b), insets are shown of both datasets with a linear fit.

Figure S3 .3

 S3 Figure S3. High temperature stress versus strain curves for high temperature rotating spindle viscosity measurements.

Figure S4 .

 S4 Figure S4. SANS spectra for all measured samples (different T soak ) taken from the medium Q range

Figure S6 .Figure

 S6 Figure S6. EDS linescan and map through the region of a spherical particle; note that both the linescan and EDS image of Mo (lower left box) shows these areas are concentrated in Mo.

Figure

  Figure S8. EDS map and linescan, showing possible Ca-Nd oxyapatite crystal

Figure S9 .

 S9 Figure S9. XRD of all soak-and-quench samples. Note that alumina standard for quantification is not present in these measurements shown.

Figure S10 :

 S10 Figure S10: (left) After first heat treatment at 1400C; (center) Closer view of crystals after first 1400C heat treatment; (right) After second heat treatment at 1400CPowder XRD results are shown below. There are a few unfitted peaks that could not be identified using EVA. The refined unit cell parameters are listed in the Error! Reference source not found.below and the refinement is shown, both of these were obtained in TOPAS.Na-La powellite space group and unit cell parametersPhase Name SG a (Å) c (Å) Na-La powellite I41/aZ 5.344276 11.73823

Figure S11 :

 S11 Figure S11: Na-La oxyapatite (left to right): after 80°C oven; after 600°C heat treatment; after 1000°C heat treatment; after 1550°C heat treatment; pellet

7 Absorption

 7 Slight blue shifts in the 4 I 9/2 → 2 P 1/2 transition were seen (431.3-430.7 nm) with increasing the soak temperature from 1175°C to 1350°C, indicating an increase in the Nd-O bond distance (nephelauxetic effect) with decreasing soak temperature S4, S5, S6 (the 2 P 1/2 level does not split with the crystal field, so it is typically monitored to infer information about Nd-O bond distance and Nd distribution). Increase of the Nd-O bond distance caused electrons to move from the oxygen ligands to the central Nd ion, reducing the bond covalency. This lead to expansion in the partially filled f-shell electron toward the oxygen ligand, reducing the distance and in turn interaction between the electrons, leading to a lowering in the energy of the 2 P 1/2 level relative to the ground 4 I 9/2 level. S4, S7 A shoulder could be clearly seen in the Ca-Nd oxyapatite phases due to Nd 3+ occupying two sites, i.e., 9-coordinated 6h and 7-coordinated 4f sites; S5, S8 the low energy shoulder absorption is due to Nd in the 6h site. The shoulder could be seen less clearly in the 1175°C and 1200°C samples, indicating the crystallization of Nd-containing oxyapatite phase; the 4 I 9/2 → 2 P 1/2 transition was broad and was shifted closer to the site 4f in oxyapatite in the 1250°C to 1350°C samples, indicating a more amorphous nature; however, this might not have been the only reason for not seeing a shoulder. It is important to note that due to the measurements being performed at room temperature, the higher energy Stark levels of the 4 I 9/2 ground state were partially populated, leading to broadening and asymmetry in the 4 I 9/2 → 2 P 1/2 transition.S4

Figure S12: T 1 23

 1 Figure S12: T 1 23 Na measurements of crystalline species (left) and glass species (right)

Figure S13 .

 S13 Figure S13. Pt heater strip with crystallized glass after 24 h at 975°C: BSE image taken with EPMA.

  Per the chemical analyses below, point 1 is powellite, points 2 and 3 are oxyapatite.

Table 5

 5

	, crystal fraction increased only slightly as the soak

Table 1 . Composition of starting glasses used in these studies shown in descending order of abundance by mass.

 1 soak and quench glasses, Hua Chen for viscosity sample preparation, and Scott Burroughs for electron microprobe help. DA, Gan H, Pegg IL. X-ray absorption and Raman spectroscopy studies of molybdenum environments in borosilicate waste glasses. J Nucl Mater. 2017;488:143-9. A, Majérus O, Lenoir M, Caurant D, Klementiev K, Webb A. Effect of alkali and alkalineearth cations on the neodymium environment in a rare-earth rich aluminoborosilicate glass. J Non-Cryst Solids. 2008;354(2-9):98-104.

	51, 55

McKeown

Quintas

Table 2 . Critical values from thermal analysis and viscosity measurements. All parameters are from DSC heating measurements except as noted.

 2

	Temperature	SIMP	Ru-SIMP
	T g (K)	882, 869 (a)	881
	T c (K)	1042, 923 (b) 993, 1097, 923 (b)
	T p (K)	1060	1024, 1111
	(a) Extrapolated T g from low temperature viscosity and VFT equation
	(b) DSC on cooling

Table 3 . Viscosity measurements for SIMP glass. Data are presented in degrees Kelvin due to the need for fitting viscosity in these units. Temperatures in Celsius, η and ln η are provided in the Supplementary for model fitting purposes.

 3

	T (K) log (η, Pa•s)
	861.0 12.46
	871.6 11.95
	879.3 11.41
	890.3 10.89
	897.4 10.61
	902.1 10.31
	912.7	9.86
	923.3	9.37
	933.4	9.00
	938.5	8.81
	1323.2	2.52±1.40
	1373.2	2.14±1.35
	1423.2	1.38±0.69
	1473.2	0.77±0.02
	1523.2	0.54±0.34
	1573.2	0.47

Table 4 . R g values for SIMP samples from SAXS analysis. Errors are determined from fitting errors associated with Irena software. 33

 4

	T soak (°C)		R g (Å)
	1175	330	± 31
	1200	205	± 19
	1250	73	± 12
	1300	13.40 ± 0.85
	1350	13.40 ± 0.89

Table 5 . Quantitative fit results for crystalline phase fraction (mass%), based on room-temperature (RT) laboratory XRD and WAXS. Amorphous fraction could only be determined for the XRD, since for the WAXS, there was no internal standard. Errors on the crystal fraction values are estimated as described in the text. Values listed as "n.m." indicate that they were not measured.

 5

		WAXS			RT XRD
	Sample	Oxyapatite Ca 2 Nd 8 Si 6 O 26	Powellite CaMoO 4	Oxyapatite Ca 2 Nd 8 Si 6 O 26	Powellite CaMoO 4	Other	Amorphous fraction
	1350	0	0	0	0	0	100.0
	1300	0	0	0	0	0	100.0
	1250	56.6	43.4	1.0	0.6	0	98.4
	1200	85.8	14.2	3.0	0.6	0.9 (Nd 0.67 MoO 4 )	95.6
	1175	89.8	10.2	5.6	0.4	0	94.0
	975-2 min	n.m.	n.m.	22.0	7.5	0	70.5
	975-24 h	n.m.	n.m.	37.1	12.3	0	50.6
	1×SC (SIMP) #	n.m.	n.m.	18.1	8.0	0	73.9
	1×SC (RuO 2 -SIMP) #	n.m.	n.m.	19.2	8.5	1.8 (ZrO 2 )	70.4
	1×SC (CX) #	n.m.	n.m.	17.6	9.0	0.6 (BaMoO 4 )	72.8
	# sample slow cooled in Pt crucibles with lids in air; analyzed with Bruker XRD; TiO 2 internal standard

  13-

	F o r
	P e
	e r
	R e
	v
	i e
	w
	Field Code Changed
	Field Code Changed
	Field Code Changed
	3

Table 1 . Composition of starting glasses used in these studies shown in descending order of abundance by mass.

 1

	Oxides	GC SIMP mass% mole%	mass%	CX	mole%
	SiO2	35.51	48.02	31.00		47.68
	Nd2O3	20.43	4.93	9.09		2.50
	B2O3	9.73	11.36	8.53		11.32
	CaO	8.91	12.91	5.12		8.44
	Na2O	8.80	11.54	2.73		4.07
	Al2O3	6.13	4.89	4.18		3.79
	MoO3	5.95	3.36	6.59		4.23
	ZrO2	4.54	2.99	5.03		3.77
	Ce2O3	0	0	5.23		1.47
	Cs2O	0	0	4.85		1.59
	BaO	0	0	3.72		2.24
	La2O3	0	0	2.67		0.76
	Pr2O3	0	0	2.44		0.68
	Sm2O3	0	0	1.81		0.48
	SrO	0	0	1.65		1.47
	TeO2	0	0	1.11		0.64
	Y2O3	0	0	1.06		0.43
	Li2O	0	0	1.05		3.25
	Rb2O	0	0	0.71		0.35
	RuO2	0	0	0.33		0.23
	Eu2O3	0	0	0.29		0.08
	Ag2O	0	0	0.19		0.08
	CdO	0	0	0.19		0.14

Table 3 . Viscosity measurements for GC SIMP glass. Data are presented in degrees Kelvin due to the need for fitting viscosity in these units. Temperatures in Celsius, η and ln η are provided in the Supplementary for model fitting purposes.

 3

	T (K) log (η, Paꞏs)
	861.0 12.46
	871.6 11.95
	879.3 11.41
	890.3 10.89
	897.4 10.61
	902.1 10.31
	912.7	9.86
	923.3	9.37
	933.4	9.00
	938.5	8.81
	1323.2	2.52±1.40
	1373.2	2.14±1.35
	1423.2	1.38±0.69
	1473.2	0.77±0.02
	1523.2	0.54±0.34
	1573.2	0.47

Table 4 . Rg values for GC SIMP samples from SAXS analysis. Errors are determined from fitting errors associated with Irena software. 33 Tsoak (°C)

 4

			Rg (Å)
	1175	330	± 31
	1200	205	± 19
	1250	73	± 12
	1300	13.40 ± 0.85
	1350	13.40 ± 0.89

2 Viscosity Viscosity measurements for GC glass in various units T (°C) T (K) 1/T (K -1 ) η (Paꞏs) ln (η, Paꞏs) log (η, Paꞏs)

	587.9	861.0	0.001161	2.88E+12 28.69	12.46
	598.5	871.6	0.001147	8.91E+11 27.52	11.95
	606.1	879.3	0.001137	2.57E+11 26.27	11.41
	617.1	890.3	0.001123	7.76E+10 25.08	10.89
	624.3	897.4	0.001114	4.07E+10 24.43	10.61
	628.9	902.1	0.001109	2.04E+10 23.74	10.31
	639.5	912.7	0.001096	7.24E+09 22.70	9.86
	650.1	923.3	0.001083	2.34E+09 21.58	9.37
	660.3	933.4	0.001071	1.00E+09 20.72	9.00
	665.3	938.5	0.001066	6.46E+08 20.29	8.81
	1050.0 1323.2	0.000756	3.32E+02	5.81±3.22	2.52±1.40
	1100.0 1373.2	0.000728	1.37E+02	4.92±3.10	2.14±1.35
	1150.0 1423.2	0.000703	2.42E+01	3.19±1.60	1.38±0.69
	1200.0 1473.2	0.000679	5.92E+00	1.78±0.043	0.77±0.02
	1250.0 1523.2	0.000657	3.49E+00	1.25±0.78	0.54±0.34
	1300.0 1573.2	0.000636	2.92E+00	1.07	0.47

  date measured slow-cool examples, where crystal fraction is 24-36 mass%. S2 Similar crystal fractions were obtained in ongoing compositionally-dependent glass studies, where major phases are oxyapatite, powellite, lanthanide borosilicate, and cerianite. In the simplified glass composition studied here, there was no cerium, so formation of cerianite [(Ce,Zr)O 2 ] was not expected. On closer examination, it was found that this is in fact cubic zirconia, likely doped with Nd 3+ for stabilization. Additionally, lanthanide borosilicate (Ln 3 BSi 2 O 10 )

Na-La oxyapatite

							Journal of the American Ceramic Society			Page 84 of 93
						Phase Name	SG		a (Å)	c (Å)
						NaLa 9 Si 6 O 26	P6 3 /m	9.690949 7.185874
								Na-La Oxyapatite 1550C HT.raw_1
		30,000												NaLa9Si6O26 100.00 %
		28,000
		26,000
		24,000
		22,000
		20,000
		18,000
	Counts	16,000 14,000 12,000 10,000 8,000 6,000 4,000 2,000 0				F o r
		-2,000 -4,000 -6,000	5	10	15	20	30 P 25 e	2Th Degrees 40 35	45	50	55	60	65	70
								e r
									R e
										v
										i e
										w
														S-14
							Journal of the American Ceramic Society

9.2 1 x slow cooled (CCC) sample 9.2.1 Microanalysis

  , C.M. Taylor Corp. Spot size was set to the minimum possible value of approximately 1.2 µm Measured compositions

						EPMA Scan Parameters
				On-
	Element/ analysis line	Analyzing Crystal	Peak Time Count	Accel. (KeV) Voltage	Probe (nA) Current	WDS Standards
				(s)
	Al Kα	TAP		10			20	20	Anorthite, USNM
	B Kα Ca Kα Nd Kα Mo Kα Na Kα	LDEB PETJ LiFH PET TAP F 200 20 20 10 10 o r	10 20 20 20 10	10 20 20 20 10	International Simple Glass Anorthite, USNM NdPO4 CaMoO4, Powellite International Simple Glass
	Si Kα Zr Kα	TAP PET Zircon #1Residual Glass 20 10 10 Wollastonite #2, C.M. Taylor Corp 10 20 20 P e e r Powellite Lathanum silicate mass fr. St. dev. mass fr. St. dev. mass fr. R e St. dev. B 2 O 3 0.117 0.015 0.0281 0.0027 0.0281 v 0.0027 Na 2 O 0.039 0.030 0.0111 0.0012 0.0111 0.0012 SiO 2 0.473 0.067 0.0025 0.0008 0.0025 i e 0.0008 MoO 3 0.009 0.003 0.7023 0.0088 0.7023 0.0088 w
		CaO	0.081	0.004	0.2311	0.0019	0.2311	0.0019
		Nd 2 O 3	0.126	0.015	0.0606	0.0026	0.0606	0.0026
		ZrO 2	0.059	0.001	0.0025	0.0013	0.0025	0.0013
		Al 2 O 3	0.090	0.005	0.0018	0.0005	0.0018	0.0005
		Total	0.993			1.0397	1.0397

S-18

and a 30-min dwell time. Measurements started at 1300°C with 100 data points taken over a 5-min time periodinterval for each measurement interval to collect shear stress measurements as a function of shear

Acknowledgements

The authors gratefully acknowledge research funding from the U.S. Department of Energy (US DOE) office of Nuclear Energy and the Nuclear Engineering University Program (NEUP), contract # DE-NE0008431. A portion of this research was conducted at the Oak Ridge National Laboratory (ORNL)'s High Flux Isotope Reactor (HFIR) sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, US DOE. The authors greatly acknowledge the help of all the ORNL HFIR staff, including the user support staff, instrument hall coordinators, and radiological protection technicians, in particular Ken Littrell at HFIR GP-SANS. A portion of this research was conducted at the Advanced Photon Source (APS) beamline 9 ID-C. Use of the APS, an Office of Science User Facility operated for the US DOE Office of Science by Argonne National Laboratory, was supported by the U.S. DOE under Contract No. DE-AC02-06CH11357. The authors greatly thank Jan Ilavsky for assistance at the APS and with SAXS analysis. Pacific Northwest National Laboratory is operated by Battelle Memorial Institute for the US DOE under contract DE-AC05-76RL01830. JVH thanks the EPSRC, the University of Warwick and the Birmingham Science City Program for partial funding of the solid-state NMR infrastructure at Warwick. The latter program accessed the Birmingham Science City Advanced Materials Project 1: Creating and Characterising Next Generation Advanced Materials, which derived support from Advantage West Midlands (AWM) and the European Regional Development Fund (ERDF). The authors also thank Carmen Rodriguez for help with running thermal analysis experiments, Charlie Bonham and Jesse Lang Acknowledgements The authors gratefully acknowledge research funding from the U.S. Department of Energy (US DOE) office of Nuclear Energy and the Nuclear Engineering University Program (NEUP), contract # DE-NE0008431. A portion of this research was conducted at the Oak Ridge National Laboratory (ORNL)'s High Flux Isotope Reactor (HFIR) sponsored by the Scientific User Facilities Division, Office of Basic Energy Sciences, US DOE. The authors greatly acknowledge the help of all the ORNL HFIR staff, including the user support staff, instrument hall coordinators, and radiological protection technicians, in particular Ken Littrell at HFIR GP-SANS. A portion of this research was conducted at the Advanced Photon Source (APS) beamline 9 ID-C. Use of the APS, an Office of Science User Facility operated for the US DOE Office of Science by Argonne National Laboratory, was supported by the U.S. DOE under Contract No. DE-AC02-06CH11357. The authors greatly thank Jan Ilavsky for assistance at the APS and with SAXS analysis. Pacific Northwest National Laboratory is operated by Battelle Memorial Institute for the US DOE under contract DE-AC05-76RL01830. JVH thanks the EPSRC, the University of Warwick and the Birmingham Science City Program for partial funding of the solid-state NMR infrastructure at Warwick. The latter program accessed the Birmingham Science City Advanced Materials Project 1:

Creating and Characterising Next Generation Advanced Materials, which derived support from Advantage West Midlands (AWM) and the European Regional Development Fund (ERDF). The authors also thank Carmen Rodriguez for help with running thermal analysis experiments, Charlie Bonham and Jesse Lang for making the soak and quench glasses, Hua Chen for viscosity sample preparation, and Scott Burroughs for electron microprobe help.

Journal of the American Ceramic Society Journal of the American Ceramic Society Journal of the American Ceramic Society Journal of the American Ceramic Society Journal of the American Ceramic Society Journal of the American Ceramic Society

Supplemental Information

The supplementary section contains additional information, tables, and figures about thermal analysis, viscosity measurements, small angle scattering, TEM, XRD, standards synthesis for 23 Na measurements, absorption measurements, 23 Na T 1 measurements, and EPMA.

Supplementary Information

The supplementary section contains additional information, tables, and figures about thermal analysis, viscosity measurements, small angle scattering, TEM, XRD, standards synthesis for 23 Na measurements, absorption measurements, and EPMA.

Journal of the American Ceramic Society

Journal of the American Ceramic Society