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[Yohei Onodera](https://orcid.org/0000-0002-3080-6991), Yasuyuki Takimoto, Hiroyuki Hijiya, Qing Li, Hiroo Tajiri, Toshiaki Ina, [Shinji Kohara](https://orcid.org/0000-0001-9596-2680)

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[Formation of a zirconium oxide crystal nucleus in the initial nucleation stage in aluminosilicate glass investigated by X-ray multiscale analysis](https://mdr.nims.go.jp/datasets/af15f0e8-4cb3-4cda-bb5e-e87ee4991712)

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Formation of a zirconium oxide crystal nucleus in the initial nucleation stage in aluminosilicate glass investigated by X-ray multiscale analysisOnodera et al. NPG Asia Materials           (2024) 16:22 https://doi.org/10.1038/s41427-024-00542-y NPG Asia MaterialsART ICLE Open Ac ce s sFormation of a zirconium oxide crystal nucleus inthe initial nucleation stage in aluminosilicate glassinvestigated by X-ray multiscale analysisYohei Onodera 1,2, Yasuyuki Takimoto3, Hiroyuki Hijiya4, Qing Li4, Hiroo Tajiri5, Toshiaki Ina6 and Shinji Kohara 1AbstractUnderstanding the nucleation mechanism in glass is crucial for the development of new glass-ceramic materials.Herein, we report the structure of a commercially important glass-ceramic ZrO2-doped lithium aluminosilicate systemduring its initial nucleation stage. We conducted an X-ray multiscale analysis, and this analysis was used to observe thestructure from the atomic to the nanometer scale by using diffraction, small-angle scattering, absorption, andanomalous scattering techniques. The inherent phase separation between the Zr-rich and Zr-poor regions in thepristine glass was enhanced by thermal treatment without changing the spatial geometry at the nanoscale. Element-specific pair distribution function analysis using anomalous X-ray scattering data showed the formation of a liquidZrO2-like local structural motif and edge sharing between the ZrOx polyhedra and (Si/Al)O4 tetrahedra during theinitial nucleation stage. Furthermore, the local structure of the Zr4+ ions resembled a cubic or tetragonal ZrO2crystalline phase and formed after 2 h of annealing the pristine glass. Therefore, the Zr-centric periodic structureformed in the early stage of nucleation was potentially the initial crystal nucleus for the Zr-doped lithiumaluminosilicate glass-ceramic.IntroductionGlass-ceramics are composed of precipitated crystalsand a glass matrix and are utilized in many industrialproducts, such as high-temperature furnace windows,cooktop panels, artificial teeth, and telescope mirrors1,2.These materials exhibit unique characteristics that are notobserved in conventional glasses. For instance, severalglass-ceramics exhibit high mechanical strength owing tothe presence of encrusted crystals, which block crackpropagation in a glass matrix. Glass-ceramics are syn-thesized by intentionally reheating quenched pristineglass above its glass transition temperature to obtainhomogeneous crystallization. The properties of glass-ceramics are affected by various factors, such as thecomposition, shape and size of the crystals and crystal-linity. Hence, these factors need to be controlled with highreproducibility to enable the practical use of glass-ceramics.The addition of nucleation agents to the host glass is aneffective method for producing glass-ceramics; nucleationagents aid and control the crystallization process3. Zir-conium oxide (ZrO2) is a commonly used nucleationagent that acts as the nucleation center for ZrO2 micro-crystals in various glass-ceramics1–3. Aluminosilicate(Al2O3–SiO2) systems with ZrO2 are the most importantcommercial glass-ceramics and are used in numerousoptical and photonic applications owing to their lowthermal expansion coefficients, high strengths, and highfracture toughness.Several studies on pristine aluminosilicate glasses andglass-ceramics with ZrO2 have been performed using acombination of X-ray diffraction (XRD), X-ray absorption© The Author(s) 2024OpenAccessThis article is licensedunder aCreativeCommonsAttribution 4.0 International License,whichpermits use, sharing, adaptation, distribution and reproductionin any medium or format, as long as you give appropriate credit to the original author(s) and the source, provide a link to the Creative Commons licence, and indicate ifchangesweremade. The images or other third partymaterial in this article are included in the article’s Creative Commons licence, unless indicated otherwise in a credit line to thematerial. Ifmaterial is not included in the article’s Creative Commons licence and your intended use is not permitted by statutory regulation or exceeds the permitted use, you will need to obtainpermission directly from the copyright holder. To view a copy of this licence, visit http://creativecommons.org/licenses/by/4.0/.Correspondence: Yohei Onodera (ONODERA.Yohei@nims.go.jp) orYasuyuki Takimoto (yasuyuki.takimoto@agc.com)1Center for Basic Research on Materials, National Institute for Materials Science,1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan2Institute for Integrated Radiation and Nuclear Science, Kyoto University,2-1010 Asashiro-nishi, Kumatori-cho, Sennan-gun, Osaka 590-0494, JapanFull list of author information is available at the end of the article1234567890():,;1234567890():,;1234567890():,;1234567890():,;http://orcid.org/0000-0002-3080-6991http://orcid.org/0000-0002-3080-6991http://orcid.org/0000-0002-3080-6991http://orcid.org/0000-0002-3080-6991http://orcid.org/0000-0002-3080-6991http://orcid.org/0000-0001-9596-2680http://orcid.org/0000-0001-9596-2680http://orcid.org/0000-0001-9596-2680http://orcid.org/0000-0001-9596-2680http://orcid.org/0000-0001-9596-2680http://creativecommons.org/licenses/by/4.0/mailto:ONODERA.Yohei@nims.go.jpmailto:yasuyuki.takimoto@agc.comspectroscopy (XAS), and transmission electron microscopy(TEM) to elucidate the local structure of Zr4+ and the for-mation mechanism of nanocrystalline ZrO24–8. In the case ofpristine glasses, Dargaud et al. reported that the Zr4+environment in MgO–Al2O3–SiO2–ZrO2 glass corre-sponded to a sevenfold coordinated site linked mainly withSiO4 tetrahedra by edge sharing, according to the resultsfrom extended X-ray absorption fine structure (EXAFS)measurements at 77 K4. Patzig et al. used a combination ofZr LII,III-edge X-ray absorption near edge structure (XANES),scanning TEM (STEM), and energy-dispersive X-ray spec-troscopy and observed sixfold coordination around Zr4+ inthe MgO–Al2O3–SiO2–ZrO2 glass5. Moreover, they showeda transformation of the Zr coordination number fromapproximately six in pristine glass to eight after the com-pletion of crystallization in MgO–Al2O3–SiO2–ZrO2 sys-tems5,6. Conversely, Cormier et al. investigated the role ofZrO2 in the nucleation of several aluminosilicate glasses byXAS at both Zr K- and LII,III-edges. They indicated that theZr4+ environment corresponded to a sixfold coordinated sitein Li- and Na-bearing glasses and that sevenfold coordinatedsites are formed in Mg-, Ca-, and Zn-containing glasses7.Based on combined XANES and TEM results from theLi2O–Al2O3–SiO2–ZrO2 glass-ceramics, Kleebusch et al.recently reported that the crystallization of nanoscale ZrO2occurred with a gradual change in the coordination state ofZr4+ from sixfold in pristine glass to eightfold in the crystal8.The Zr4+ environments in pristine glasses and suffi-ciently crystallized glass-ceramics have been studied usingvarious experimental techniques. However, in glass-cera-mics, the environment around the Zr4+ ions in an inter-mediate state between the glass and crystals, particularly inthe early stage of nucleation, is still not well understood.Structural modifications, for example, change in thecoordination number around Zr4+ and the formationprocess of the polyhedral connectivity of the ZrOx units atthe early stage of nucleation, are crucial for furtherunderstanding the role of Zr4+ in the nucleation process.In particular, determining the clustering process con-sidering the connectivity of the ZrOx units is indispensablefor understanding the origin of nucleation because theZrOx clustering potentially triggers nucleation and occurson an intermediate-range scale; this intermediate-rangescale exists on a length scale larger than the nearest-neighbor distance. However, the conventional approachusing XAFS is ineffective for investigating theintermediate-range ordering in oxide glasses owing totheir structural disorders. The EXAFS signals are imme-diately damped beyond the nearest-neighbor distance,particularly for glasses9. Therefore, XAFS is unsuitable forelucidating intermediate-range structural ordering inglasses and glass-ceramics in the early stage of nucleation.In contrast, pair distribution function (PDF) analysis isuseful for probing short- to intermediate-range orderingin glasses. The PDF is obtained by the Fourier transformof the structure factor S(Q), which is the normalizedscattering intensity obtained from quantum beam (X-rayand neutron) diffraction measurements. PDF analysisprovides the interatomic distances and average coordi-nation numbers10–12. In several previous investigations,the PDF analysis with structural modeling techniquesshowed that the coordination number of a cation in aglass was different from that in a crystal13–17. For instance,Mg–O species adopted four, five, and six-coordination insilicate glass that contains insufficient network formers13;ZnOx (average x < 4) polyhedra formed the network inZn-rich binary phosphate glass15; and K ions were trappedin highly coordinated K–O polyhedra, forming a corre-lated pair arrangement with Na–O polyhedra in silicateglass with the mixed alkali effect16. These cation coordi-nation environments are characteristic of glasses andrarely derived solely by EXAFS because the EXAFS ana-lysis of glassy materials is performed using a well-definedcrystalline structure as a reference. Hence, PDF analysis isadvantageous for investigating glass structures. However,extraction of the specific pair correlations, such as Zr-related correlations, from a PDF is challenging for mul-ticomponent glasses and glass-ceramics (which mostlyinclude four or more oxides to enable practical use)because the pair correlation peaks are greater than 10atomic pairs in the PDF. Therefore, an element-specifictechnique based on diffraction is required to probe Zr-specific structural information beyond short-rangeordering in practical glass-ceramics.In this study, we examine lithium aluminosilicate glassand its glass-ceramic, which is a commercially importantsystem due to its attractive properties, such as low/negativethermal expansion coefficients. A small amount of ZrO2(2.9mol%) was added as a nucleation agent to the glass tocrystallize the glass-ceramic. To investigate the structuralordering relative to the small amount of Zr in the glass andglass-ceramic at both short- and intermediate-range scales,we employed anomalous X-ray scattering (AXS) measure-ments. AXS utilizes anomalous variations in the atomicform factor of a specific element close to the X-rayabsorption edge and provides element-specific structuralinformation beyond the nearest-neighbor distance18–21.Therefore, the atomic distances and coordination numbersaround Zr were determined by element-specific PDF ana-lysis using the AXS data. The local structure of Zr was alsostudied using XAFS measurements. The XAFS spectra weremeasured using the Zr K-edge because K-edge XAFSmeasurements enabled the attainment of the EXAFSspectra up to a higher wavenumber (k) than those obtainedusing the L-edge. The XANES spectra were utilized todetermine the valence state of Zr, and the EXAFS spectrawere analyzed as a reference for element-specific PDFanalysis. The combination of AXS and EXAFS has beenOnodera et al. NPG Asia Materials           (2024) 16:22 Page 2 of 13    22 applied to studies of CaO–SiO2–ZrO2 glass-ceramics22 andlithium germanate glasses23 to investigate the local struc-tural motifs of Zr or Ge atoms. Small-angle X-ray scattering(SAXS) is a useful method for probing inhomogeneity inmaterials when examining the phase separation of nan-ometer or submicrometer order. Therefore, SAXS was usedto observe the behavior of the nanoscale structures duringnucleation. In this study, lithium aluminosilicate glass-ceramics in the initial nucleation stage were prepared bythermal treatment of pristine glass. The nanoscale struc-tures of the pristine glass and glass-ceramics were probedby combining SAXS with in-house XRD. The short-rangestructures of the pristine glass and glass-ceramics wereanalyzed using a combination of XAFS and AXS. Further-more, the AXS data were investigated in both Q and realspaces to determine the formation mechanism of thestructural ordering, which was consistent with the SAXSresults in the intermediate range in the early stage ofnucleation. Thus, we performed a state-of-the-art multi-scale structural analysis based on synchrotron X-rayexperiments on Zr-doped lithium aluminosilicate glassand glass-ceramic materials. This novel approach demon-strates the behavior of a small number of Zr cations fromthe atomic scale to the nanoscale in the initial nucleationstage in commercially available glass-ceramic materials.Materials and methodsMaterialsLithium aluminosilicate glass with small amounts of zir-conium oxide, sodium oxide, and phosphorus pentoxide(pristine glass) was prepared using a conventional melt-quenching method. The raw materials in the prescribedamounts were sufficiently mixed and melted at 1650 °C for12 h to achieve homogeneous melting. The material wasthen quenched and rapidly solidified to prevent partialcrystallization. The chemical composition of the pristineglass is listed in Table 1. Large portions of the obtained glasspieces were cut and polished to a thickness of 1.2mm forthe subsequent experiments. The residual glass was crushedfor thermal analysis. The pristine glass samples wereannealed at a crystallization onset temperature determinedby thermal analysis for several durations between 0 and 48 hto obtain glass-ceramic samples at different nucleationstages. The XRD measurements were performed using aRigaku SmartLab diffractometer with an X-ray source of CuKα radiation to confirm the formation of amorphous andcrystalline phases. The density of the pristine glass obtainedusing the Archimedean method was 2.55 g cm–3; this valuecorresponded to a number density of 0.0759 Å–3.Differential scanning calorimetry (DSC) measurementDSC was performed to elucidate the crystallization beha-vior of the pristine glass during heating. A crushed glasssample was sifted to obtain glass particles with sizes rangingfrom 106 to 180 μm for DSC measurements. DSC wasperformed at a constant heating rate of 10 °C min–1 using aDSC3300SA differential scanning calorimeter (Bruker, Inc.).Synchrotron X-ray measurementsThe SAXS measurements were performed using theBL8S3 beamline at the Aichi Synchrotron RadiationCenter. The energy of the incident X-rays was 13.5 keV.The camera length was 2174.1 mm, and a two-dimensional detector (PILATUS 100 K) was used. Theexposure time for each measurement was 480 s.The Zr K-edge (18.01 keV) XAFS measurements wereperformed using the BL5S1 beamline at the Aichi Syn-chrotron Radiation Center by employing a Si(111) double-crystal monochromator in transmission mode. Theobtained XAFS spectra were normalized and analyzedusing Athena and Artemis software24. The k3-weightedEXAFS spectra, k3χ(k), were Fourier transformed (FT)over a k-range of 3.0–12.0 Å–1 to obtain the radialstructure functions. The Zr–O distances (rZr–O) andZr–O coordination numbers (NZr–O) were determined bycurve fitting the EXAFS data using Artemis software24. Afixed Debye–Waller factor of 0.009 Å2 was used; thisvalue was obtained from the curve fitting to Zr K-edgeEXAFS data for yttria-stabilized zirconia (YSZ) crystals.The AXS measurements were performed using a dedi-cated AXS spectrometer21 built at the BL13XU beam-line25,26 at the SPring-8 synchrotron radiation facility. Thespectrometer consisted of a six-circle diffractometer, avacuum sample chamber, receiving slits, a beam stop, aLiF analyzer crystal, and a NaI (Tl) scintillation detector.The energy resolution of the LiF(200) crystal wasapproximately 12 eV at the full width at half maximum at12 keV, which enabled the discrimination of the con-tributions of fluorescence and Compton scattering. TheAXS measurements were performed at two incident X-rayenergies, 17.81 keV (Zr far edge, Efar) and 17.99 keV (Zrnear edge, Enear); these were 200 eV and 20 eV below theZr K-edge, respectively. The scattering intensities werecollected using reflection geometry for 200min for eachTable 1 Glass composition.Composition mol%SiO2 70.7Al2O3 14.3P2O5 0.7ZrO2 2.9Li2O 9.3Na2O 2.1Total 100.0Onodera et al. NPG Asia Materials           (2024) 16:22 Page 3 of 13    22 sample. The differential intensity (ΔZrI) between twoscattering intensities measured at Efar and Enear isexpressed as follows:αΔZrI Q; Efar;Enearð Þ ¼ ΔZr f Qð Þ2� �� f Qð Þh i2� �þΔZr f Qð Þh i2� �ΔZrS Qð Þ;ð1Þandf Qð Þ2� � ¼Xicif2i Qð Þ; ð2Þf Qð Þh i2 ¼Xicif i Qð Þ !2; ð3Þwhere α is a normalization constant, ΔZr[ ] indicates thedifference in values in the brackets at the energies of Efarand Enear of Zr, and ci and fi(Q) are the concentration andatomic form factor of the component atom i, respectively.The differential structure factor ΔZrS(Q) is given by alinear combination of the partial structure factors Sij(Q),as follows:ΔZrS Qð Þ ¼XNi¼1XNj¼1wij Q;Efar;Enearð ÞSij Qð Þ; ð4Þwhere the weighting factors are given by the following:wij Q;Efar; Enearð Þ ¼ cicjΔZr f i Qð Þf j Qð Þh iΔZr f Qð Þh i2� � : ð5ÞCompared with the total structure factor S(Q) obtainedfrom XRD, ΔZrS(Q) greatly enhances the contributionfrom Zr-related Sij(Q) and suppresses the contributionsfrom the other partials. By approximating fi(Q) by usingthe atomic number, the weighting factors in the ΔZrS(Q)can be calculated as follows:ΔZrS Qð Þ ¼ 0:033SZrZr Qð Þ þ 0:312SZrSi Qð Þþ 0:117SZrAl Qð Þ þ 0:018SZrLi Qð Þþ 0:015SZrNa Qð Þ þ 0:521SZrO Qð Þ;ð6Þwhere the weighting factors for the partials other thanthose mentioned above were almost zero; these partialstructure factors were reasonably eliminated inΔZrS(Q).Results and discussionPreparation of the glass-ceramic samples in the initialnucleation stageThe DSC curve of the pristine glass is shown in Fig. 1.Three broad exothermic maxima were observed atapproximately 800, 860, and 1045 °C. The in-house XRDpatterns of the samples annealed at different temperaturesare shown in Fig. S1. In the XRD pattern of the sampleannealed at 800 °C, several broad peaks were observed atapproximately 30°, 35°, 50°, and 60° and were assigned tocubic (c-) or tetragonal (t-) ZrO2. However, we could notdistinguish between these two phases because the XRDpeaks were extremely broad. Based on these results, thefirst exothermic peak in the DSC curve was assigned tothe nucleation of c- or t-ZrO2 in the pristine glassstructure. The second and third exothermic peaks wereassigned to the actual crystallization of the β-quartz solidsolution (s.s.)27 and the phase transition from β-quartz s.s.to β-spodumene s.s.27, respectively. A slight baseline shiftto the endothermic direction owing to the glass transitionwas observed at approximately 745 °C in the DSC curve.As shown in Fig. 1 and S1, the morphological changeduring annealing in pristine glass occurred in four steps:(i) a decrease in viscosity at the glass transition tem-perature, (ii) nucleation of the ZrO2 crystal, (iii) crystal-lization of β-quartz s.s., and (iv) phase transition fromβ-quartz s.s. to β-spodumene s.s.A series of glass-ceramic samples at different nucleationstages were synthesized to elucidate the structural changesaround Zr during the nucleation process. By elucidatingthe results of the thermal treatments based on the DSCcurve, an annealing temperature of 770 °C was found to besuitable; this temperature corresponded to the crystal-lization onset temperature. Based on this consideration, thepristine glass was annealed at 770 °C for various durations.The in-house XRD patterns of the samples prepared byannealing at 770 °C for various durations are shown inFig. 2a. A broad halo pattern was observed in the data forthe 0-h-annealed sample (pristine glass), indicating acompletely disordered structure. Subtle diffraction peakswere observed at approximately 30° and 50° in the XRDpatterns of the 2- and 3-h-annealed samples. Prominentdiffraction peaks assigned to c- or t-ZrO2 were found atapproximately 30°, 35°, 50°, and 60° in the XRD patterns ofthe samples annealed longer than 4 h (shown as closedcircles in Fig. 2a). The crystallite sizes of the samplesFig. 1 DSC curve at a heating rate of 10 °C min–1 for Zr-doped lithiumaluminosilicate glass.Onodera et al. NPG Asia Materials           (2024) 16:22 Page 4 of 13    22 annealed for longer than 4 h were calculated from theScherrer equation using the peak width observed at 30°.The crystallite size and peak height versus annealingduration are shown in Fig. 2b. The annealing durationdependence of the crystallite size was considerably weakerthan that of the peak height; thus, the thermal treatmentincreased the number of ZrO2 crystallites rather thantheir size. Similar behavior was reported for the nuclei ofZrO28 and ZrTiO428 in Li2O–Al2O3–SiO2-based glasses.As shown in Fig. 2b, the generation of ZrO2 crystalliteswas gradually saturated with successive annealing cycles.In the XRD pattern for the 48-h-annealed sample, anadditional peak was observed at 26° (shown by the closedtriangle in Fig. 2a) and was attributed to a peak ofβ-quartz s.s.27. The ZrO2 crystallites gradually pre-cipitated upon thermal treatment at 770 °C. Therefore,glass-ceramics suitable for investigating the structuralchanges during the initial nucleation stage were obtained.Analysis of nanoscale structureThe SAXS profiles of the samples prepared for variousannealing durations are shown in Fig. 3a, b. A relativelylow scattering intensity was observed for the 0-h-annealedsample, indicating that the pristine glass already had aninternal inhomogeneous structure. The mean distance, d,between scatterers was calculated using d= 2π/Q, whereQ is the peak position in the SAXS profile. The calculatedd values and SAXS peak heights are shown in Fig. 3c. TheSAXS peak height gradually increased with the annealingduration. These results indicated that the electron densitycontrast between the scatterers increased with increasingannealing duration. Conversely, the distance betweenscatterers was approximately 12 nm and did not show anynotable change with increasing annealing duration,indicating that the spatial distribution of scatterers wasunaffected by the thermal treatment.As mentioned in the discussion of the XRD results, thesize of the ZrO2 crystallites was almost unchanged afterFig. 2 XRD data for the Zr-doped lithium aluminosilicate glass and glass-ceramics. a In-house XRD patterns for the Zr-doped lithiumaluminosilicate glass and glass-ceramics obtained by annealing at 770 °C for various annealing durations (shown on the right side of the figure).Closed circles and triangles: peak positions of cubic/tetragonal ZrO2 and β-quartz solid solutions, respectively. b Annealing duration dependence ofthe peak height (closed circles) and crystallite size (open triangles) extracted from the diffraction peak observed at 30°.Fig. 3 SAXS profiles and structural parameters for Zr-dopedlithium aluminosilicate glass and glass-ceramics. a SAXS profilesof the samples annealed for 0 h (black), 1 h (green), 2 h (red), and 4 h(blue). b SAXS profiles of the samples annealed for 4 h (blue), 6 h(cyan), 8 h (red), 10 h (green), 16 h (orange), 24 h (magenta), and 48 h(brown). c. Annealing duration dependence of the peak height (closedcircles) and 2π/Q (open triangles) extracted from the SAXS profile.Onodera et al. NPG Asia Materials           (2024) 16:22 Page 5 of 13    22 thermal treatment, whereas the number of ZrO2 crystal-lites increased with annealing duration. If the SAXSprofiles originated from the correlation between ZrO2crystallites, the SAXS peak position should shift to ahigher Q with the annealing duration because an increasein the number of ZrO2 crystallites in a given volume couldcause the particles to approach the others. Furthermore,explaining the generation of the peak in the SAXS profileof the pristine glass that does not include ZrO2 crystallitesis difficult. Therefore, correlations between the ZrO2crystallite particles are unlikely to account for the SAXSpeaks in this system. Instead, the Zr-related phaseseparation was presumed to act as a scatterer for SAXS inthis system. Based on the SAXS profile with a scatteringpeak for the 0-h-annealed sample, the internal phaseseparation into Zr-rich and Zr-poor phases in pristineglass was inherent. The thermal treatment at 770 °Cinduced the aggregation of ZrOx and subsequent partialprecipitation of a ZrO2 crystallite. Furthermore, with anincrease in the annealing duration, the separation wasenhanced without changing the distance between thescatterers. The internal phase separation in glasses wasreported for B2O3–PbO–Al2O329 and Al2O3–SiO230 glasssystems. In these systems, the phase separations werecharacterized as binodal or spinodal31 textures by Porodanalysis of the SAXS patterns. In contrast, the SAXSprofiles of pristine glass and glass-ceramics in this studyshowed a clear scattering peak and did not obey Porod’slaw, causing difficulty to characterize the phase separationbehavior in the initial stage of nucleation based on thePorod plot. Based on the XRD patterns, the generation ofa particularly large number of ZrO2 crystallites wasinduced by thermal treatment. However, the SAXS dataindicated that the separation between the Zr-rich and Zr-poor phases was inherent in pristine glass and increasedwith increasing annealing duration. The results from thecombined in-house XRD and SAXS measurements indi-cated that the structural change in the Zr-containinglithium aluminosilicate glass-ceramic during the initialstage of nucleation originated from the separationbetween the Zr-rich and Zr-poor phases and that ZrO2crystallites precipitated in the Zr-rich phase. Notably,these structural changes occurred without any change inthe spatial nanoscale geometry.Nanoscale heterogeneities and the formation of nano-crystals have been reported in TEM-based studies of severalZrO2-containing lithium aluminosilicate systems. Klee-busch et al. used a combination of XANES and TEM toreport the formation of Zr-rich phase separation droplets,which evolved into crystalline ZrO2 in Li2O–Al2O3–SiO2glass-ceramics8. Höche et al. reported the precipitation ofZrTiO4 nanocrystals in phase-separated droplets accom-panied by the formation of a circumjacent shell consistingof an Al-rich region in Li2O–Al2O3–SiO2 glass-ceramicsusing TEM32. Raghuwanshi et al. quantitatively evaluatedthe core–shell structure using anomalous SAXS33. Based onthese aforementioned studies, the growth of nanoprecipi-tates was constrained by a shell-like region8,32,33. The XRDresults indicated that the size of the ZrO2 crystallites wasalmost unchanged upon thermal treatment; thus, the ZrO2crystallites that precipitated in the Zr-rich phase potentiallyformed a core–shell-like structure, and in the shell region,the growth of nanoprecipitates was prevented.Analysis of the short-range Zr-centric structures by XAFSTo obtain clear insight into the structural changes in theshort- and intermediate-range scales in the initialnucleation stage in the ZrO2-containing glass-ceramic, weconducted structural analyses using a combination ofXAFS and AXS measurements. These measurementswere performed at the Zr K-edge to investigate Zr-specificstructural information.The Zr K-edge XANES spectra of the samples preparedfor various annealing durations are shown in Fig. 4a. Theedge jump position did not change between these spectra,indicating that the valence state of the Zr ion was tetra-valent and was unchanged with the annealing duration.However, the shape of the XANES spectra graduallychanged with annealing duration, indicating that thestructure around the Zr4+ ions gradually changed.Moreover, the isosbestic points (indicated by arrows inFig. 4a) were observed in all XANES spectra; thus, theXANES spectra were formed by the superposition of theXANES spectra of the two states. These results wereconsistent with the aforementioned SAXS results, inwhich a structural change related to the separationbetween the Zr-rich and Zr-poor phases was proposed.The XANES spectrum of the 0-h-annealed sample (pris-tine glass) likely consisted of Zr-rich and Zr-poor phaseswithout the precipitation of ZrO2 crystallites. In contrast,a larger portion of the XANES spectrum of the 48-h-annealed sample should consist of the Zr-rich phase inwhich ZrO2 crystallites were precipitated by thermaltreatment. The XANES spectra of the 0- and 48-h-annealed samples were compared with those of thestandard crystalline samples in Fig. 4b. In this study, weused YSZ, with a stable c-ZrO2 at room temperature, andmonoclinic (m-) ZrO2 as the standard crystalline mate-rials. A comparison of these XANES spectra indicatedthat the structure of the sample changed from a glass-likedisordered structure to a crystal-like ordered structurebecause the XANES spectra of standard crystallinematerials were more similar to those of the 48-h-annealedsample than to those of the 0-h-annealed sample. More-over, in the XANES spectrum of the 48-h-annealedsample, the multiple scattering feature34 observed in theXANES spectra of t-ZrO2 and YSZ35 was observed as asmall rounded peak at approximately 18030 eV (indicatedOnodera et al. NPG Asia Materials           (2024) 16:22 Page 6 of 13    22 by the arrow in Fig. 4b). Based on these results for the 48-h-annealed sample, either a Zr-related local structuralorder formed in the c-ZrO2 (YSZ) or t-ZrO2 was intro-duced into its structure.The FT k3-weighted EXAFS spectra of the samples areshown in Fig. 5a. The first peak in the FT spectra observedat approximately 1.6 Å was assigned to the Zr–O nearest-neighbor correlation. The distance observed in the FT k3-weighted EXAFS spectra shown in Fig. 5a, b wasapproximately 0.2–0.5 Å shorter than the actual distancebecause of the phase shift9. The Zr–O distance slightlyshifted to a longer R with the annealing duration. Thesecond nearest-neighbor peak appeared at approximately3.0–3.2 Å for the samples annealed for 6 h or more; thepeak became stronger with increasing annealing duration.The Zr–O distances and coordination numbers in the 0-,2-, 4-, 6-, 10-, 24-, and 48-h-annealed samples weredetermined by curve fitting to their EXAFS spectra; thesefindings are summarized in Table 2. The Zr–O coordi-nation number gradually increased over 4 h of annealingand reached approximately 7 after 48 h. An increase in theZr–O coordination number associated with nucleationwas reported in previous XAFS studies of lithium alumi-nosilicate systems5,6,8. Similarly, an increase in the Ti–Ocoordination number during nucleation was reported inlithium aluminosilicate glass that contained TiO2 as anucleation agent in XANES studies36. Thus, an increase inthe coordination number around a cation in the nuclea-tion agent was a general trend during the nucleation ofthe aluminosilicate glasses. The FT k3-weighted EXAFSspectra of the 0- and 48-h-annealed samples are shownalong with those of the standard YSZ and m-ZrO2 sam-ples in Fig. 5b. The first and second peak positions werealmost the same for the 48-h-annealed and YSZ samplesFig. 4 XANES spectra of Zr-doped lithium aluminosilicate glass/glass-ceramics and ZrO2 crystals. a XANES spectra of samples annealed for 0 h(black), 2 h (red), 4 h (blue), 6 h (cyan), 10 h (green), 24 h (magenta), and 48 h (brown). The arrows mark the isosbestic points. b XANES spectra ofsamples annealed for 0 h (black) and 48 h (brown) along with those of YSZ (blue) and monoclinic zirconia (m-ZrO2) (green). The arrow indicatesmultiple scattering, which is a characteristic feature of tetragonal ZrO234 and cubic ZrO2 in YSZ35.Fig. 5 Fourier transform of k3-weighted EXAFS functions for Zr-doped lithium aluminosilicate glass/glass-ceramics and ZrO2 crystals.a Fourier transform of the k3-weighted EXAFS functions of the samples annealed for 0 h (black), 2 h (red), and 4 h (blue), 6 h (cyan), 10 h (green), 24 h(magenta), and 48 h (brown). b Fourier transform of the k3-weighted EXAFS functions of the samples annealed for 0 h (black) and 48 h (brown) alongwith those for YSZ (blue) and monoclinic zirconia (m-ZrO2) (green).Onodera et al. NPG Asia Materials           (2024) 16:22 Page 7 of 13    22 but differed greatly between the 48-h-annealed andm-ZrO2 samples; these results indicated that the ZrO2microcrystalline phase generated by thermal treatmentwas not monoclinic. The slight difference between the FTk3-weighted EXAFS spectra of the 48-h-annealed and YSZsamples indicated the presence of residual Zr4+ ions inthe glassy phase in the 48-h-annealed sample or thestructural distortion around Zr4+ ions in the 48-h-annealed and/or YSZ sample. The structural changes inthe post-nucleation stage were elucidated by Zr–Ocoordination number analyses using EXAFS data. How-ever, the second nearest-neighbor correlation, whichprovides crucial information in the early stage ofnucleation, was barely observed in the FT k3-weightedEXAFS spectra of the pristine glass and 2–4-h-annealedsamples. Thus, the local structure around the Zr4+ ionsfor the pristine glass and the 2–4-h-annealed glass-ceramic samples was considerably disordered for theobservation of the second nearest-neighbor correlationsperformed with the EXAFS measurements. Thus, inprinciple, observing atomic correlations beyond the firstcoordination shell in our samples during the early stage ofnucleation using EXAFS was difficult. Thus, we analyzedthe AXS data to determine a Zr-specific structure beyondthe nearest-neighbor correlations.Analysis of the Zr-related structures at short- andintermediate-range scales by AXSFigure 6a–c show the X-ray scattering intensities Icoh(Q)measured at Efar and Enear after annealing the samples for0, 2, and 4 h. The Icoh(Q) measured at Efar was slightlygreater than that measured at Enear. The first sharp dif-fraction peak (FSDP) was observed at approximately1.6 Å–1, which was close to the position of the FSDP ofSiO2 glass37, in all X-ray scattering patterns shown inFig. 6a–c. Because more than 70% of the components ofthe sample are SiO2, the FSDP was mainly attributed to anetwork structure stemming from the SiO2 glass, whichconsisted of SiO4 tetrahedra with shared oxygen atoms atthe corners. The differences in the positions and shapes ofthe FSDPs among the 0-, 2-, and 4-h-annealed sampleswere minimal, indicating that the SiO4 tetrahedralTable 2 Coordination number and distance of Zr–O obtained by EXAFS analysis. The Debye–Waller factors were fixed to0.009 Å2 in all curve fittings.Samples 0 h 2 h 4 h 6 h 10 h 24 h 48 hrZr–O (Å) 2.03 ± 0.04 2.05 ± 0.03 2.06 ± 0.04 2.07 ± 0.04 2.07 ± 0.05 2.07 ± 0.04 2.06 ± 0.03NZr–O 5.3 ± 1.7 5.2 ± 1.7 5.5 ± 1.7 5.9 ± 1.7 6.2 ± 1.6 6.5 ± 1.5 6.9 ± 1.932106543210 0 h 2 h 4 h60040020006543210 Far Near1005006543210 0 h 2 h 4 h60040020006543210 Far Near60040020006543210 Far Nearba cdQ ( 1)I coh(Q)0 hQ ( 1)I coh(Q)2 hQ ( 1)I coh(Q)4 hQ ( 1)I coh(Q)eQ ( 1)I coh(Q)Fig. 6 X-ray scattering intensities and differential intensities for Zr-doped lithium aluminosilicate glass/glass-ceramics. a X-ray scatteringintensities measured at Efar (red) and Enear (blue) for the 0-h-annealed sample. b X-ray scattering intensities measured at Efar (red) and Enear (blue) forthe 2-h-annealed sample. c X-ray scattering intensities measured at Efar (red) and Enear (blue) for the 4-h-annealed sample. d Differential intensitiesbetween two X-ray scattering intensities measured at Efar and Enear for the samples annealed for 0 h (black), 2 h (red), and 4 h (blue). e Enlarged viewof the high-Q region of the differential intensities of the samples annealed for 0 h (black), 2 h (red), and 4 h (blue).Onodera et al. NPG Asia Materials           (2024) 16:22 Page 8 of 13    22 network did not change significantly upon thermaltreatment. The Icoh(Q) values measured at both Efar andEnear increased at Q < 0.5 Å–1 with increasing annealingduration. The increasing scattering intensity atQ < 0.5 Å–1 was potentially due to the inhomogeneity atthe nanoscale caused by thermal treatment, as mentionedin the discussion of the SAXS results. In addition, thedifferential intensity ΔIcoh(Q) between the two scatteringintensities measured at Efar and Enear, as shown in Fig. 6d,increased at Q < 0.5 Å–1 with increasing annealing dura-tion; thus, the phase separation enhancement by thermaltreatment was related to Zr4+. The behavior of ΔIcoh(Q) inthe lower-Q region supported the SAXS results. Fur-thermore, ΔIcoh(Q) significantly changed in the higher-Qregion (Q > 0.5 Å–1), as shown in Fig. 6e, indicating thatthe Zr-related structure at short- and intermediate-rangescales underwent changes upon thermal treatment.Figure 7a shows the differential structure factorsΔZrS(Q) for the 0-, 2-, and 4-h-annealed samples. The firstpeak of ΔZrS(Q) for the 0-h-annealed sample wasobserved at approximately 2.1 Å–1; this was different fromthe position of the FSDP observed in the X-ray scatteringintensities. Therefore, the scale of the periodicity relatedto Zr4+ was different from that related to the SiO4 net-work. The total structure factor, S(Q), for liquid ZrO238 isshown in Fig. 7c. The positions of the first peak of ΔZrS(Q)for the 0-h-annealed sample and S(Q) for liquid ZrO2were significantly close, indicating that the periodicityrelated to Zr4+ in pristine glass was similar to that inliquid ZrO2. Compared with that of pristine glass, the firstpeak positions of ΔZrS(Q) for the 2- and 4-h-annealedsamples did not shift but became sharper with an increasein the annealing duration. In addition, the minute Braggpeaks (shown as open circles in Fig. 7a) appeared inΔZrS(Q) for the 2- and 4-h-annealed samples. The peakpositions were close to the diffraction peaks of c- ort-ZrO2, indicating that the initial crystal nucleus wasformed during the annealing of pristine glass in a shorttime. Furthermore, compared to in-house XRD mea-surements, AXS measurements enable more sensitivedetection of crystallization behavior involving a specificelement.Differential total correlation functions, ΔZrT(r), for the0-, 2-, and 4-h-annealed samples and the total correlationfunction, T(r), for liquid ZrO2 obtained from the Fouriertransforms of ΔZrS(Q) and S(Q) are shown in Fig. 7b, d,respectively. The Zr–O correlation peak was located at2.1 Å for all ΔZrT(r) and T(r) values. For liquid ZrO2, theasymmetric Zr–O correlation peak with a tail ofapproximately 2.8 Å indicated the formation of distortedZrOx polyhedra38. The second peak, with a shoulder onthe low-r side, was observed at 3.6 Å for all ΔZrT(r) values.In the case of liquid ZrO2, the second peak observed at3.6 Å in T(r) was assigned to the Zr–Zr correlation, cor-responding to the distance between the centers of ZrOxpolyhedra with a large fraction of edge sharing of oxygenin addition to corner sharing38. In the crystalline materi-als, a Zr–Zr distance of approximately 3.6 ÅFig. 7 Differential and total structure factors and differential and total correlation functions for Zr-doped lithium aluminosilicate glass/glass-ceramics and liquid ZrO2. a Differential structure factor ΔZrS(Q) and b differential total correlation function ΔZrT(r) of the samples annealed for0 h (black), 2 h (red), and 4 h (blue). c X-ray total structure factor S(Q) and (d) X-ray total correlation functions T(r) for liquid ZrO2 at 2800 °C38. The S(Q)and T(r) values calculated using the pair function method are plotted as green curves.Onodera et al. NPG Asia Materials           (2024) 16:22 Page 9 of 13    22 corresponding to the edge sharing of ZrO8 polyhedra wasobserved for t-ZrO2 and crystalline ZrSiO4 (zircon)39.Thus, the second peak in ΔZrT(r) was assigned mainly tothe Zr–Zr correlation corresponding to edge- and corner-sharing of ZrOx polyhedra. The Zr–Si, Zr–Al, andsecond-neighbor Zr–O correlations contributed to thesecond peak in ΔZrT(r) in addition to the Zr–Zr correla-tions according to Eq. (6). The shoulder peak on the low-rside of the second peak was likely composed of the tail ofthe asymmetric Zr–O peak and other Zr-related correla-tion peaks; this peak was observed at approximately2.8–2.9 Å for all ΔZrT(r) values. In Eq. (6), only six partialscontribute to ΔZrS(Q), and the weighting factors for Zr–Siand Zr–Al correlations are relatively large, with theexception of the Zr–O correlation in ΔZrS(Q); therefore,we could estimate that the shoulder peak mainly consistedof the Zr–Si and/or the Zr–Al correlation superposed onthe tail of the Zr–O peak. Both the second and shoulderpeaks of ΔZrT(r) evolved with increasing annealing dura-tion, indicating that the development of Zr-relatedstructural ordering beyond the nearest-neighbor dis-tance was induced by thermal treatment.The Zr–O coordination numbers of the 0-, 2-, and 4-h-annealed samples and liquid ZrO2 were determined usingthe pair function method proposed by Mozzi and War-ren40. The pair function is a useful method for analyzingreal-space functions, from which the structural para-meters such as the interatomic distance and coordinationnumber can be estimated. From the concept of the pairfunction, the calculated total correlation function Tcalc(r)was obtained using the following equation for the distancerZr–O and coordination number NZr–O of the Zr–O pair:T calcZr-O rð Þ ¼ 2πZ QmaxQmin2cZrNZr-Of Zr Qð Þf O Qð Þf Qð Þh i2 exp� 12l2Zr-OQ2� �sin πQ=Qmaxð ÞπQ=Qmaxsin QrZr-Oð ÞrZr-OsinQr dQ:ð7ÞThe term lZr-O is a convergence factor that representsthe static and thermal disorders of the Zr–O correlation.In contrast, the calculated differential total correlationfunction ΔTcalc(r) was obtained using the followingequation:ΔT calcZr-O rð Þ ¼ 2πRQmaxQmin2cZrNZr-OΔZr f Zr Qð Þf O Qð Þ½ �ΔZr f Qð Þh i2� � exp�12l2Zr-OQ2� �sin πQ=Qmaxð ÞπQ=Qmaxsin QrZr-Oð ÞrZr-O sinQr dQ:ð8ÞThe Zr–O coordination numbers can be estimated byfitting the calculated ΔT(r) and T(r) for the Zr–Ocorrelation to the experimental ΔZrT(r) and T(r). Weadopted two Zr–O pair functions to reproduce theasymmetric Zr–O peak in ΔZrT(r) and T(r). In this pro-cedure, the variations in the rZr–O and NZr–O in the cal-culated ΔT(r) and T(r) were ±0.02 Å and ±0.3,respectively. The ΔZrS(Q), S(Q), ΔZrT(r), and T(r), alongwith the results from the pair function method (greencurve), are shown in Fig. 7a–d, respectively. As shown inFig. 7a, c, good agreement between the pair function andthe experimental data was obtained at Q > 3.8 Å–1. Thestructural parameters obtained using the pair-functionmethod are summarized in Table 3. In the case of liquidZrO2, both the distance and average coordination numberof Zr–O were in agreement with those obtained by thedensity functional–molecular dynamics simulationreported by Kohara et al.38. The Zr–O coordinationnumber for the annealed samples increased from 5.1 to6.0 with increasing annealing duration. This increase inthe Zr–O coordination number for the 0-, 2-, and 4-h-annealed samples was not observed from the EXAFSanalysis, as shown in Table 2, because the presence of anasymmetric atomic distribution reduced the amplitude ofthe EXAFS signal; this reduction in amplitude is fre-quently accompanied by reductions in the bond distanceand coordination number. In contrast, the analysis of AXSdata using the pair function method enabled the use oftwo Zr–O pair functions to reproduce the asymmetricZr–O peak in ΔZrT(r) and T(r). These total correlationfunctions quantitatively described the atomic pair dis-tribution in glassy materials, enabling the direct analysisof the asymmetric atomic correlation peaks. A successiveincrease in the Zr–O coordination number from 5.1 to 6.0was achieved during nucleation in a glass structurebecause the rigid glass network sustained a metastablestructural unit of the Zr4+ ion in the glass matrix. TheZr–O coordination number of 6.0 for the 4-h-annealedsample was similar to that of liquid ZrO2, indicating that aliquid-like local structural motif around a Zr4+ ion wasformed during the initial nucleation stage.The Zr–Si and/or Zr–Al correlations observed as theshoulder peak in ΔZrT(r) corresponded to the distancebetween the centers of the ZrOx polyhedra and the (Si/Al)O4 tetrahedra. The Zr–(Si/Al) distance of 2.8–2.9 Å wasconsiderably short for the Si or Al at the center of a (Si/Al)O4 tetrahedron sharing an oxygen at the corner with aZrOx polyhedron. Therefore, the short Zr–(Si/Al) distanceindicated the formation of edge-sharing polyhedral con-nections between ZrOx and (Si/Al)O4. A short Zr–Sidistance of 2.98 Å corresponding to edge sharing betweenZrO8 and SiO4 was observed in zircon39. Furthermore,using EXAFS analyses, Cormier et al. determined that theZr–(Si/Al) distance at approximately 3.1 Å correspondedto edge sharing between ZrO6 or ZrO7 and (Si/Al)O4 inZr-containing aluminosilicate glasses7. The presence ofOnodera et al. NPG Asia Materials           (2024) 16:22 Page 10 of 13    22 the shoulder peak in ΔZrT(r) for the 0-h-annealed sampleindicated the inherent edge-sharing polyhedral connec-tion in pristine glass; this was not a typical structuralfeature of glass-forming materials. The shoulder peak inΔZrT(r) increased with annealing duration, indicatingthat thermal treatment promoted the formation of edge-sharing structures. Thus, we determined that a liquid-likelocal structural motif around a Zr4+ ion and theedge-sharing structures between ZrOx polyhedra and(Si/Al)O4 tetrahedra were formed by thermal treatment ofthe pristine glass. These Zr-related structural changeslikely played a key role in the initial stages of nucleation.To obtain clear insights into the initial stage of nucleationfrom the atomic scale to the nanoscale, we constructed aschematic representation of the structural change from thepristine glass to glass-ceramic after 4 h of annealing, asshown in Fig. 8a–d. Figure 8a, b show a comparison of therepresentations of the nanoscale structures based on theXRD and SAXS results for the pristine glass and the 4-h-annealed sample, respectively. As shown by the SAXSresults, pristine glass had an inherent phase separationbetween the Zr-rich and Zr-poor regions (Fig. 8a). Theelectron density in the Zr-rich region was greater than thatin the Zr-poor region. Thermal treatment of the pristineglass for 4 h caused aggregation of ZrOx and an increase incontrast in the electron density between the Zr-rich and Zr-poor phases (Fig. 8b). The precipitation of ZrO2 crystallitesin the Zr-rich regions occurred only during the 2–4-hTable 3 Structural parameters derived by the pair function method.Samples 0 h 2 h 4 h Liquid ZrO2Zr–O(1) Zr–O(2) Zr–O(1) Zr–O(2) Zr–O(1) Zr–O(2) Zr–O(1) Zr–O(2)rZr–O (Å) 2.10 ± 0.02 2.67 ± 0.02 2.05 ± 0.02 2.63 ± 0.02 2.10 ± 0.02 2.63 ± 0.02 2.09 ± 0.02 2.49 ± 0.02lZr–O (Å2) 0.01 0.01 0.01 0.01 0.01 0.01 0.09 0.09NZr–O 3.8 ± 0.3 1.3 ± 0.3 3.8 ± 0.3 1.3 ± 0.3 3.9 ± 0.3 2.1 ± 0.3 4.0 ± 0.3 1.7 ± 0.3Total 5.1 5.3 6.0 5.7NZr–OFig. 8 Schematic representation of the initial nucleation stage in zirconium-doped aluminosilicate glass and glass-ceramics. Nanoscalestructures of the 0-h-annealed sample (pristine glass) (a) and 4-h-annealed sample (b). Representations of the intermediate-range structure in the Zr-rich phase for pristine glass (c) and the 4-h-annealed sample (d). Green: zirconium, blue: silicon, cyan: aluminum, magenta: phosphorus, orange:lithium, purple: sodium, and red: oxygen. The yellow polyhedra indicate edge-sharing polyhedral connections.Onodera et al. NPG Asia Materials           (2024) 16:22 Page 11 of 13    22 annealing of pristine glass (Fig. 7a). Moreover, the crystallitesize of approximately 30 Å did not increase with theannealing duration, whereas the number of ZrO2 crystallitesincreased. Similar phase separation behavior during thenucleation process, in which the growth of nanoprecipitateswas limited by the size of the phase-separated domains, wasreported for various glass-ceramics8,28,32,33. In these cases,nanoscale heterogeneities, consisting of precipitated nano-crystals and a glass matrix, played an important role inyielding glass-ceramics that exhibited excellent mechanicaland thermal properties while maintaining their transpar-ency1,2. The addition of small amounts of nucleation agentsto the host glass induced phase separation at the nanos-cale3,32,41. DeCeanne et al. proposed the following definitionof a nucleation agent: “a nucleating (nucleation) agent is aminority component of the glass composition that leads toincreased internal nucleation rates or precipitation andcontrol of desired crystal phases, either by lowering thethermodynamic or the kinetic barrier for nucleation, orsome combination thereof”3. Thus, phase separationsinduced by nucleation agents could affect the thermo-dynamics or kinetics of nucleation by forming functionalphases (e.g., the Zr-rich regions shown in Fig. 8b) that couldact as precursors to nanocrystals, leading to enhancednucleation. Figure 8c, d show schematic representations ofthe intermediate-range structure in the Zr-rich region forpristine glass and the 4-h-annealed sample, respectively. Inthe pristine glass, SiO4, AlO4, and only a few PO4 polyhedraformed a tetrahedral network, whereas Zr4+, Li+, and Na+cations were distributed around the network (Fig. 8c).Several Zr4+ ions formed an edge-sharing structure with the(Si/Al)O4 tetrahedra in the structure of pristine glass. Afterannealing for 4 h, the structure became more ordered andwas centered on a Zr cluster, as shown in Fig. 8d. Comparedto the structure of pristine glass, Zr4+ ions congregated andformed an ordered arrangement, whereby the Zr–O coor-dination number increased and additional edge-sharingstructures formed between ZrOx and surrounding (Si/Al)O4tetrahedra. This specific configuration involving Zr4+ ionscould be the initial crystal nucleus of the Zr-doped lithiumaluminosilicate glass-ceramic. The Zr-related configurationformed a periodic c- or t-ZrO2-like structure (ZrO2 nano-crystal) during annealing for 2–4 h, generating the Braggpeaks in the differential structure factor ΔZrS(Q) (Fig. 7a).This finding was consistent with the SAXS results becausethe difference in the electron density between the con-gregated Zr4+ ions (Zr-rich phase) and the other ions (Zr-poor phase) increased with the formation of an orderedarrangement of Zr4+ ions in the Zr-rich phase. The ZrO2nanocrystal surrounded by the (Si/Al)O4 tetrahedral net-work shown in Fig. 8d corresponded to the core–shellstructure reported in previous studies8,32,33,36, as mentionedin the discussion of the SAXS results. The shell structure,consisting of a (Si/Al)O4 tetrahedral network, could act as anucleation barrier that prohibited the growth of ZrO2crystals. This behavior supported the theoretical modelpredicted by the generalized Gibbs approach42; here, thecomposition and structure of a nucleus significantly chan-ged, while the size remained nearly constant in the earlystage of nucleation. Hence, our proposed model for theinitial crystal nucleus of ZrO2 (Fig. 8d) was consistent withboth nanoscale structural studies8,32,33,36 and the theoreticalapproach42. Notably, the short Zr–(Si/Al) correlations,described as Zr–O–(Si/Al) bonds in Fig. 8c, d, playedimportant roles in the formation of the initial crystalnucleus. Zr–(Si/Al) correlations beyond the nearest-neighbor distance have been rarely observed using con-ventional approaches; however, we succeeded in theirinvestigation by using Zr-specific PDF analysis with AXSdata. These findings may provide new insights into theintermediate-range structure of nucleation agents in glassesand promote a better understanding of the nucleationmechanism in the initial stages of glass-ceramic materials.In this study, we performed a state-of-the-art multiscalestructural analysis by combining XRD, SAXS, XAFS, andAXS measurements. Based on a comprehensive analysis ofthe atomic-to-nanoscale structure, we succeeded in thedetermination of the Zr-related intermediate-range struc-ture that was the initial crystal nucleus for the nucleationprocess in lithium aluminosilicate glass. Multiscale struc-tural analysis combined with the element-specific mea-surements enabled the attainment of crucial information toelucidate the relationship between the structures andmaterial properties in multicomponent disordered systems.AcknowledgementsThis work was partially supported by JSPS Grant-in-Aid for Transformative ResearchAreas (A) “Hyper-Ordered Structures Science” (grant numbers 20H05878 (to S.K.),20H05881 (to Y.O., H.T. and S.K.)), and for Scientific Research (C) (grant number19K05648 (to Y.O.)). Synchrotron radiation experiments were performed at BL13XUof SPring-8 with the approval of the Japan Synchrotron Radiation ResearchInstitute (JASRI) (Proposal No. 2019A1722) and at BL5S1 and BL8S3 of the AichiSynchrotron Radiation Center (Approval Nos. 2019D1001 and 2019D1002).Author details1Center for Basic Research on Materials, National Institute for Materials Science,1-2-1 Sengen, Tsukuba, Ibaraki 305-0047, Japan. 2Institute for IntegratedRadiation and Nuclear Science, Kyoto University, 2-1010 Asashiro-nishi,Kumatori-cho, Sennan-gun, Osaka 590-0494, Japan. 3Innovative TechnologyLaboratories, AGC, Inc., 1-1 Suehiro-cho, Tsurumi-ku, Yokohama, Kanagawa230-0045, Japan. 4Materials Integration Laboratories, AGC, Inc., 1-1 Suehiro-cho,Tsurumi-ku, Yokohama, Kanagawa 230-0045, Japan. 5Scattering and ImagingDivision, Japan Synchrotron Radiation Research Institute (JASRI, SPring-8), 1-1-1Kouto, Sayo-gun, Hyogo 679-5198, Japan. 6Spectroscopy Division, JapanSynchrotron Radiation Research Institute (JASRI, SPring-8), 1-1-1 Kouto, Sayo-cho, Sayo-gun, Hyogo 679-5198, JapanAuthor contributionsY.O. and Y.T. designed the study. The samples were prepared by H.H. and Q.L.Differential scanning calorimetry was performed by H.H. Small-angle X-rayscattering was conducted by Y.T. X-ray absorption was performed by Y.T. andT.I. Anomalous X-ray scattering was conducted by Y.O., Y.T., H.T. and S.K. Theobtained data were analyzed by Y.O., Y.T., H.H., T.I. and S.K. The manuscript waswritten by Y.O., Y.T. and S.K. with input from all authors.Onodera et al. NPG Asia Materials           (2024) 16:22 Page 12 of 13    22 Conflict of interestThe authors declare no competing interests.Publisher’s noteSpringer Nature remains neutral with regard to jurisdictional claims inpublished maps and institutional affiliations.Supplementary information The online version contains supplementarymaterial available at https://doi.org/10.1038/s41427-024-00542-y.Received: 9 August 2023 Revised: 31 January 2024 Accepted: 18 March2024References1. Beall, G. H. & Pinckney, L. R. Nanophase glass-ceramics. J. Am. Ceram. Soc. 82,5–16 (1999).2. 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NPG Asia Materials           (2024) 16:22 Page 13 of 13    22 https://doi.org/10.1038/s41427-024-00542-y Formation of a zirconium oxide crystal nucleus in the initial nucleation stage in aluminosilicate glass investigated by X-ray multiscale analysis Introduction Materials and methods Materials Differential scanning calorimetry (DSC) measurement Synchrotron X-ray measurements Results and discussion Preparation of the glass-ceramic samples in the initial nucleation�stage Analysis of nanoscale structure Analysis of the short-range Zr-centric structures by�XAFS Analysis of the Zr-related structures at short- and intermediate-range scales�by AXS Acknowledgements Acknowledgements