Investigation of the Role of Nucleating Agents in MgO−SiO2

DOI: 10.1021/cg101318p

Investigation of the Role of Nucleating Agents in MgO-SiO2-Al2O3-SiO2-TiO2 Glasses and Glass-Ceramics: A XANES Study at the Ti K- and L2,3-Edges

2011, Vol. 11 311–319

)

L. Cormier,*,† O. Dargaud,†,‡ N. Menguy,† G. S. Henderson,§ M. Guignard,† N. Trcera, and B. Watts^ †

)

Institut de Min
eralogie et de Physique des Milieux Condens
es, CNRS UMR 7590, Universit
e Pierre et Marie Curie - Paris 6, Universit
e Denis Diderot - Paris 7, IPGP, 4 place Jussieu, 75005 Paris Cedex, ‡ § France, Saint-Gobain Recherche, 39 Quai Lucien Lefranc, 93 303 Aubervilliers, France, Department of Geology, University of Toronto, 22 Russell St, Toronto, Ontario M5S 3B1 Canada, Synchrotron Soleil, BP 48, 91192 Gif-sur-Yvette Cedex, France, and ^Swiss Light Source, Paul Scherrer Institut, 5232 Villigen-PSI, Switzerland Received October 7, 2010; Revised Manuscript Received November 12, 2010

ABSTRACT: We report a X-ray absorption study at Ti K- and L2,3-edges to determine the role of TiO2 content as a nucleating agent in glass ceramics. It is found that the local Ti environment is not modified with TiO2 addition, indicating that medium range organization is responsible for its ability to promote internal versus bulk nucleation. We have identified a Ti coordination change between the nucleation front and the crystallized part of the glass ceramic. These changes correspond to conversion from 5-fold coordinated to 4-fold coordinated Ti in the remaining glassy part, resulting from compositional changes. This reveals that active sites for nucleation could be experimentally detected and that reorganization of the glass matrix during nucleation has a major influence on the ongoing nucleation processes.

Introduction Crystal nucleation in the disordered structure of glasses is the key operation in the production of many advanced materials.1 Understanding the nucleation pathways is also important in regard to the applicability of standard nucleation theories.2 The challenge and requirement for the study of nucleation are to obtain structural information at the atomic level, especially during the initial stages of the glass-to-crystal transformation. It is then important to separate the nucleation stage in which only a small fraction of the parent glass is affected from the crystallization of the glass, observed at higher temperatures. Nucleation occurs usually from the external surface of the glass, which is circumvented by adding various nucleating agents (for instance, TiO2, ZrO2). This will promote bulk nucleation and improve crystallization kinetics,1,3,4 giving materials having improved properties (low thermal expansion coefficient, low dielectric constant, and high chemical and mechanical stability).3,5-7 Though they are extensively used industrially, the fundamental structural role of these agents is still not well established: how do they induce structural changes in the glass and promote nuclei formation? The difficulty lies in the ability to gain information when nucleating agents start to be active, that is, in the first stages of the nucleation during the formation of nanometric crystals. This requires spatially resolved experiments, the most popular being various electron microscopy methods.8 However, the development of nano- or micro-X-ray absorption spectroscopy allows the determination of the nucleating agent environment with a resolution that has not been achieved in the past. X-ray absorption spectroscopy (XAS) has been long used to investigate the environment of typical nucleating agents

(Ti, Zr, Cr) in glasses and glass-ceramics.9-14 Extensive information about the coordination environment of Ti4þ in oxide glasses has been obtained by using XAS either at the Ti K-edge15,16 or at the L2,3-edges.17 Both edges are very sensitive to local order and can give information on the mixing of coordinated sites and on the site geometry. Those XAS techniques were used to determine the environment of titanium mainly for the TiO2-SiO2 binary glasses18-20 and Ti-bearing alkali or alkaline-earth silicate glasses.16,17,21-24 All these studies show that titanium atoms are present in the glass with multiple coordination numbers from four to six, with often the predominance of 5-fold coordinated sites. The 5-fold coordination has been determined using neutron diffraction using Ti isotopic substitution and has been unambiguously assigned to the symmetry of a square based pyramid.25,26 In this paper, X-ray absorption near-edge structure (XANES) spectroscopy is used to investigate the local structure around Ti4þ in magnesium aluminosilicate glasses and glass-ceramics, in order to investigate the effects of different additions of TiO2 and different thermal treatment leading to surface or bulk crystallization. Detailed information on the Ti environment in both the glass and the glass ceramics is obtained. Experimental Section

*To whom correspondence should be addressed. E-mail: cormier@ impmc.upmc.fr.

Sample Preparation. Glasses were prepared with the molar chemical composition 2MgO-2Al2O3-5SiO2 þ xTiO2 with x = 0, 2, 4, 6, 8, 10 mol %. The glasses were prepared by melting the dried starting powders (MgO, Al2O3, SiO2, and TiO2) at 1600 C for 1 h in a Pt crucible and quenched from high temperature by immersing the bottom of the crucible into water. The obtained glasses were ground and melted once again to ensure a good homogeneity. The compositions were determined using electron microprobe microanalyser (CAMECA SX50) at the Camparis Centre (Universit
e Pierre et Marie Curie, France). The chemical compositions of the studied glasses are reported in Table 1.

r 2010 American Chemical Society

Published on Web 12/08/2010

pubs.acs.org/crystal

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Table 1. Composition (mol %) of the Glasses Investigated in This Study glass

MgO

Al2O3

SiO2

TiO2

MAS-2Ti MAS-4Ti MAS-6Ti MAS-8Ti MAS-10Ti

21.64 21.44 21.00 20.66 19.64

21.83 21.42 20.84 20.74 20.22

54.52 53.19 52.17 51.26 50.07

2.01 3.95 5.99 7.34 10.07

The glass samples exhibit a slight brownish color. This may be due to a Ti-O-Fe charge transfer band resulting from Fe impurities or to the reduction of some Ti4þ ions into Ti3þ ions. Trivalent titanium ions are the only ions with a free electron, and their amount can be calculated by using electron paramagnetic resonance (EPR) (Bruker). Results indicate less than 300 ppm Ti3þ ions of total titanium atoms. Transmission electron microscope images (JEOL) confirmed the absence of nanometer-size heterogeneities (crystalline or amorphous). Several glass samples of approximately 5  5  1 mm3 were heat treated at a temperature up to 920 C at a heating rate of 5 K min-1 and then held at this temperature for 15 min. The samples were then removed from the furnace and allowed to cool down at room temperature. MAS-8Ti glass specimens of approximately 5  5  1 mm3 were heat treated at a temperature of 775 C, with a heating rate of 5 K/min, for different periods of time (1, 2, 4, 6, 8, and 12 h). One MAS-8Ti specimen received the same heat treatment at 775 C for 12 h and a further heat treatment at 950 C for 2 h (heat rate of 5 K/min). X-ray powder diffraction (XRD) patterns were recorded using a Xpert’Pro Panalytical diffractometer operating with a Co X-ray radiation. Differential scanning calorimetry (DSC) thermograms were recorded on a Setaram Multi-HTC 96, under N2 flux, in Pt crucible with a heat rate of 5 K min-1 from 450 C up to 1550 C. Energy calibration was made with respect to sapphire, and calculations were carried out as mentioned previously.27 Reference samples are those used in a previous study.17 They correspond to 4-fold (Ba2TiO4), 5-fold (Ba2TiS2O8, fresnoite), and 6-fold (TiO2, rutile) coordinated titanium. Eleven phases in the MgTi2O5-Al2TiO5 (MAT) solid solution were synthesized with the molar compositions (MgTi2O5)x-(Al2TiO5)1-x for which x is comprised between 0 and 1. Starting materials (MgO, Al2O3, and TiO2) were finely ground together in ethanol, and the resulting powders were pressed to make pellets. The pellets were then heated 24 h at 1500 C and quenched to room temperature. Ti L2,3-Edges X-ray Absorption Spectroscopy. The Ti L-edge spectra were recorded on the PolLux-STXM beamline of the Swiss Light Source (SLS, Switzerland).28 Scanning transmission X-ray microscopy (STXM) allows spatial resolution of ∼20 nm. The XANES spectra have been recorded in transmission mode over an energy range between 455 and 472 eV with steps of 0.1 eV and a counting time of 50 ms. The samples are prepared on a semitransparent silicon nitride membrane. The spectra have been corrected by a pre-edge background using a polynomial function and normalized to the peak with the maximum intensity.17 Ti K-edge X-ray Absorption Spectroscopy. Ti K-edge XANES spectra have been collected on the LUCIA beamline of the SOLEIL synchrotron (France) using a Si(111) double crystal monochromator.29 The XANES spectra have been recorded in the fluorescence mode with a silicon drift diode detector, over an energy range from 4850 to 5200 eV, with steps of 0.2 eV from 4960 to 4980 eV and a counting time of 2 s. At least two spectra have been averaged to increase the signal-to-noise ratio. The probed spot area was ∼5  5 μm. A Ti-foil was used to provide an absolute calibration of the monochromator. The bulk samples were put on a graphite tape stuck on copper slides. Experimental data have been reduced by subtraction of the background, and the spectra have been normalized to the high energy side of the spectra.30 Pre-edge parameters (absolute position and height relative to the edge jump) were extracted from the normalized spectra by fitting two Lorentzian profiles in the region 4962-4975 eV.

Results Results for DSC Data. Figure 1 shows the DSC curves for the glasses with different TiO2 content, and Table 2 compares

Figure 1. Differential scanning calorimetry curves of glasses with different TiO2 content showing the variation in the crystallization peaks for MAT (MgTi2O5-Al2TiO5) and β-quartz phases. Table 2. Characteristic Parameters Extracted from the DSC Curves in Figure 1,a sample

Tg ( 2 (C)

ΔCp ( 0.5 (J g-1 K-1)

Ton ( 2 (C)

TX ( 2 (C)

MAS-0Ti MAS-2Ti MAS-4Ti MAS-6Ti MAS-8Ti MAS-10Ti

813 803 784 772 773 761

0.6 0.6 0.6 0.68 0.50 0.51

920 896 877 878 872 866

947 912 896 896 895 889

a Tg is the glass transition temperature, ΔCp is the heat capacity jump at Tg. Ton and TX are the temperatures corresponding to the onset of the first crystallization phenomenon and to the maximum of the first crystallization peak, respectively.

characteristic parameters extracted from the DSC data. The attribution of the different events is in agreement with previous investigations.3,31,32 The curves exhibit similar shapes with an endothermic signal corresponding to the glass transition and exothermic features corresponding to crystallization phenomena. TiO2 plays a role in the thermodynamic properties of the system from its initial addition in the sample with 2 mol % TiO2 (MAS-2Ti), by lowering the glass transition temperature and the temperatures for the first crystallization event. The sample without TiO2 (MAS-0Ti) shows a first crystallization event starting at 920 C that can be ascribed to the formation of βquartz. With addition of TiO2, the first crystallization peak is shifted to lower temperatures, down to 866 C for 10 mol % TiO2. Ti L2,3-Edges XAS. L2,3 spectra obtained for the crystalline compounds (Figure 2) are similar to an earlier study, and we used the same labeling for the different features.17 The peak positions are given in Table 3. The spectra are dominated by the coupling of the 3d electron with the 2p core hole. Because of the 2p spin-orbit (LS) coupling of the 2p hole, the 2p edge is separated into the 2p3/2 (L3 edge) and 2p1/2 (L2 edge). The energy difference between both peaks is related to the crystal field parameter 10Dq. For each edge, two peaks (c-d and g-h) appear which are directly related to the t2g and eg orbital splittings in perfect Oh symmetry. The second peak (d and h) in both L3 and L2 edges is larger due to vibrational and dispersional broadening.33 The splitting of

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the second peak (d and e) in the L3 edge is not well explained: distortion of the local environment,34 long-range effects,35,36 or coupling of electronic and vibrational states have been proposed.36 All details of the L2 edge are blurred compared to the L3 edge because of the intrinsic broadening.34 It was shown that the separation and relative intensities of peaks g and h can be accounted for by the differences in the first coordination shell of the Ti atoms.35 Peaks a and b (small prepeaks on the low energy side of the L3 edge) are related to a forbidden transition in LS-coupling but which becomes allowed because of the multipolar 2p-3d interactions.34,35 The spectrum for the MAS-8Ti glass is compared in Figure 2 with the spectra for the crystalline references and Table 3 compares the characteristic peak positions. Peak c is at the same position than that for the fresnoite spectrum and shifted by þ0.2 eV compared to that of Ba2TiO4 spectrum. This peak was related to the basal-plane distortions of TiO5 square-pyramids, and it is shifted to low energies for distorted TiO5 pyramids.37 The peak f of the fresnoite is not observed in the glass spectrum, but the positions of peaks g and h are at similar positions for glasses and fresnoite. The L-edge XANES spectra suggest that Ti is predominantly present as [5]Ti. A similar conclusion was obtained by neutron diffraction, but, based on Reverse Monte Carlo modeling,

Figure 2. Ti L2,3-edges spectra for Al2TiO5, MgTi2O5, TiO2 (rutile), Ba2TiSi2O8 (fresnoite), Ba2TiO4, and glass MAS-8Ti. The upper curve is a linear combination of reference spectra (see text for an explanation). The Ti coordination for each crystalline phase is indicated.

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the presence of [4]Ti and [6]Ti was also assessed with proportion of 20.8% and 27.7%, respectively.38 Given the proportion obtained previously,38 an averaged spectrum (upper curve in Figure 2) is obtained using a linear combination of the spectra for Ba2TiO4, Ba2TiSi2O8, and Al2TiO5. Alhough this average spectrum does not perfectly fit the spectrum for the glass, it shows that mixing between various Ti coordination sites must be taken into account in the MAS glasses, though polyhedral distortions might also cause broadening of the features.39 In Figure 3, the spectra for the samples after heat treatment are compared with the crystalline references and Table 3 compares the characteristic peak positions. As the duration of the heat treatment increases, we observe an increase in intensity of peak c, a shift by þ0.2 eV and a broadening of the peak d, a shift by -0.3 eV and an increase in intensity of peak g, and a shift by þ0.25 eV of the peak h. The final spectrum (12 h-775 C þ 2 h-950 C) is close to that observed for the MgTi2O5 and Al2TiO5 crystals, in which Ti is 6-fold coordinated. All these changes suggest a coordination change from mainly [5]Ti to mainly [6]Ti during nucleation. The local changes around the Ti environment are consistent with X-ray diffraction patterns (Figure 4) where we observe the initial formation of a magnesium aluminotitanate crystal in the MgTi2O5-Al2TiO5 solid solution.3,32,40,41

Figure 3. Ti L2,3-edges spectra for MAS-8Ti glass with various heat treatment and comparison with Ba2TiSi2O8 (fresnoite, [5]Ti, lower curve) and Al2TiO5 and MgTi2O5 crystals ([6]Ti, upper curves).

Table 3. Peak Positions (( 0.1 eV) of the L-Edge Spectra for the Crystalline Phases and the MAS-8Ti Glass and Glass-Ceramicsa peak

a

Ba2TiO4 Ba2TiSi2O8 TiO2 (rutile) MgTi2O5 Al2TiO5 MAS-8Ti glass 1 h-775 C 2 h-775 C 4 h-775 C 6 h-775 C 8 h-775 C 12 h-775 C 12 h-775 C þ2 h-950 C

∼456.6

a

456.7 456.8 456.8

b 457.1 457.3 457.6 457.8

457.0 ∼457.0 ∼456.7 456.7 456.8 456.8

∼457.6 ∼457.6

c

d

458.2 458.4 458.2 458.4 458.6

460.2 460.1 ∼460.1 ∼460.2 460.6

458.5 458.55 458.5 458.4 458.4 458.4 458.4 458.4

Peaks a-h correspond to the features in Figure 2 (see text for an explanation).

460.2 460.3 460.3 460.3 460.3 460.4 460.4 460.4

e

∼460.9 ∼460.9

f 462.7

g

h

463.8 464.1 463.6 463.6 463.9

465.8 465.6 466 ∼466.5 466.0

464.1 464.0 464.1 463.9 463.9 463.9 463.85 463.8

466.6 466.6 466.6 466.6 466.7 465.9 465.9 465.85

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Figure 4. X-ray diffraction patterns obtained for the glasses with various heat treatments and for crystals in the MgTi2O5-Al2TiO5 solid solution. The Bragg peaks correspond to crystals in the MgTi2O5-Al2TiO5 solid solution. A background was used in the glass ceramic patterns to remove the glass contribution.

Figure 5. Normalized heights vs energy positions of the pre-edge features at the Ti K-edge, showing the different coordination domains, from Farges et al.24 MAS glasses with various TiO2 content are represented by squares. Various points below the surface for the MAS-4Ti glass ceramic are plain circles, and various points below the surface for the MAS-6Ti glass ceramic are plain diamonds.

We have identified that the molar composition of the nanocrystals is close to (MgTi2O5)40-(Al2TiO5)60.42 These phases are built from edge-shared MgO6, AlO6, and TiO6 octahedra.43,44 Bragg peaks appear only for the samples heat treated for 4 h or more at 775 C (Figure 4). The changes in the Ti environment during the nucleation are particularly visible in the peak c in Figure 3; though its position is unchanged, an important increase in intensity is observed

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Figure 6. Ti K-edge XANES pre-edge features for MAS glasses with various TiO2 content.

as the proportion of [6]Ti increases. The changes appear continuous with the heat treatment indicating that a larger fraction of Ti is involved in the first nanocrystal as the heat treatment increases or that the Ti environment in the remaining glass during the early nucleation becomes more similar to that in the MAT phases. Ti K-Edge XAS. A typical XANES spectrum at Ti K-edge shows an intense pre-edge at 4970.8 eV, followed by a shoulder at ∼4978 eV and two asymmetric maxima in the region of the absorption edge, respectively, at 4989 and 5000 eV. The spectrum is dominated by the 1s f p transition and the pre-edge is due to transitions from 1s energy levels of Ti to the Ti3d/O2p molecular orbitals. These latter transitions are particularly interesting as they are very sensitive to the local environment. The 1s f 3d electronic transition is forbidden in centrosymmetrical sites (e.g., regular octahedron) but partially allowed in noncentrosymmetrical sites where p-d orbital mixing occurs (e.g., tetrahedron). The pre-edge is thus very sensitive to the coordination polyhedron, its geometry, and its distortion.45 Using a large selection of Ti-bearing model compounds, Farges et al. has established a correlation between the coordination environment around Ti and the pre-edge parameters (position and height), which is summarized on the diagram in Figure 5.16,45 Three domains can be separated corresponding to [4]Ti, [5]Ti, and [6]Ti. The pre-edge has the greatest height for [4]Ti and is shifted toward lower energies by ∼2 eV when compared to [6]Ti. Pre-edge feature for [5]Ti is intermediate between [4]Ti and [6]Ti. Various mixtures of Ti coordinations can also be detected. The pre-edges for glasses with different amounts of TiO2 are shown in Figure 6. The pre-edge information has been extracted by fitting the energy region between 4962 and 4975 eV using two Lorentzian functions. The pre-edge parameters for the studied glasses are in the domains of various coordination sites (squares in Figure 5), mainly the [5]Ti and [6]Ti domains, with normalized intensities ranging from 40 to 43% at 4970.8 ( 0.2 eV. There is no significant variation of the pre-edge features with an increasing amount of TiO2. Bulk MAS samples were heat treated at 920 C for 15 min, and fractured sections were measured at different points starting from the surface to the bulk with steps of ∼100 μm, in order to monitor the Ti structural changes during nucleation from the surface to the bulk. The extracted pre-edges are

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Figure 7. Extracted normalized pre-edge features for glasses and glass ceramics obtained by a heating ramp of 5 K mn-1 up to 920 C for samples with 2 (A), 4 (B), 6 (C), 8 (D), and 10 (E) mol % TiO2. Each curve reports the glass pre-edge and those obtained each 100 μm step below the surface of the glass ceramic.

displayed in Figure 7 and the extracted parameters are reported in Table 4. The corresponding XRD patterns are shown in Figure 8. They can be interpreted similar to previous XRD investigations.3,31,32,46,47 At 2 mol % TiO2, the pre-edge features present no modification, while the corresponding XRD pattern displays Bragg peaks corresponding to a phase with the structure of β-quartz solid solution (also called μ-cordierite) and to the R-cordierite

phase (also called indialite or high cordierite). The crystallization is only on the surface. At 4 mol % TiO2, the surface crystallization is more important and this affects the Ti environment. The pre-edge features show a reduced intensity below the surface indicating the formation of [6]Ti. The last pre-edge feature, at 500 μm below the surface, is close to the pre-edge for the glass, indicating that, beyond the crystallization front, the Ti environment is slightly changed. At 6

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Table 4. Pre-Edge Parameters Obtained by Fitting with Pseudo-Voigt Function for the Different MAS Glasses and Glass-Ceramicsa sample MAS-2Ti glass 200 μm 300 μm 400 μm MAS-4Ti glass 200 μm 300 μm 400 μm 500 μm MAS-8Ti glass 200 μm 300 μm 400 μm 500 μm 600 μm 700 μm 800 μm 900 μm MAS-8Ti glass 200 μm 300 μm 400 μm MAS-10Ti glass 200 μm 300 μm

position (eV)

normalized height

fwhm (eV)

4970.7 4970.8 4970.7 4970.7

0.42 0.41 0.42 0.41

1.00 0.98 0.99 0.99

4970.7 4970.9 4971.0 4971.0 4970.8

0.40 0.33 0.23 0.28 0.38

1.03 1.08 1.28 1.10 1.04

4970.7 4970.6 4970.6 4970.6 4970.6 4970.6 4971.0 4971.0 4970.8

0.40 0.55 0.59 0.55 0.60 0.47 0.25 0.29 0.40

1.01 0.89 0.88 0.88 0.88 0.99 1.18 1.11 1.02

4970.6 4971.3 4971.2 4971.2

0.42 0.21 0.22 0.21

1.00 1.23 1.20 1.21

4970.7 4971.2 4971.2

0.43 0.21 0.20

0.99 1.32 1.32

a

Uncertainties for position, normalized height, and FWHM are (0.1 eV, ( 0.01 and (0.02 eV, respectively.

mol % TiO2, the crystallization is still from the surface. The evolution of the Ti K-edge XANES spectra from the surface to the bulk shows a strong nonlinear variation. Up to 500 μm below the surface, the pre-edge information corresponds to the center of the [5]Ti domain, and then it falls to the [6]Ti domain. At 900 μm below the surface, it corresponds to the initial glass, outside the crystallization zone. At 8 and 10 mol % TiO2, bulk crystallization operates with Ti playing its role of nucleating agent. The pre-edge feature for the glasses and glass-ceramics are distinctly different. Glass-ceramics are characteristic of [6]Ti (pre-edge at 4971.2 eV with intensity of 24-21%), with no change between the surface and the bulk. The pre-edge corresponding to [6]Ti is asymmetric at low energy values and the asymmetry increases as the TiO2 content increases. The shoulder at ∼4969.5 eV is close to the position expected for [4]Ti. However, such a shoulder also exists for crystalline references containing [6]Ti.48 Discussion Above 8 mol % TiO2, the L- and K-edges spectra results show agreement that the MAS-8Ti glass contains a majority of 5-fold coordinated species. This is in accordance with a previous neutron diffraction investigation on the same glass composition where the method of isotopic substitution has been used.38 The [5]Ti coordination is usually found as the predominant species in silicate glasses,16,17,24,49 and it is usually associated with a square based pyramid, characterized by two sets of Ti-O distances.25,26 The existence of various Ti coordination sites for the MAS-8Ti glass has been quantified in a Reverse Monte Carlo model with a ratio of 2:5:3 for [4] Ti/[5]Ti/[6]Ti.38 Similarly, it was previously shown that [6]Ti is relatively important when small alkaline-earth cations (Mg2þ,

Figure 8. X-ray diffraction patterns obtained for the MAS glass ceramics with various TiO2 content heat treated at 920 C (heat ramp 5 K min-1) for 15 min.

Ca2þ) are present, though usually at a high TiO2 content (>30 mol %).24 The L2,3-edges XANES spectrum is compatible with a mixing of various Ti coordination sites (see Figure 2, upper curve), and the pre-edge information extracted from the K-edge XANES spectrum lies between the domains for [5]Ti and [6]Ti. During the nucleation, the L2,3-edges spectra undergo major changes which are compatible with Ti coordination changes from mainly [5]Ti to mainly [6]Ti. Modifications of the medium range order around Ti were also observed using neutron diffraction with Ti isotopic substitution, indicating the formation of MAT nanocrystals.42 The major changes appear in the intensity of the peak c and in the position of peak g. This indicates that L2,3-edges spectra can be successfully used to follow the Ti changes in the course of nucleation. However, due to the size of the nanocrystals (∼10-20 nm),42 the spatial resolution of the STXM technique (20 nm) is insufficient to resolve the Ti distribution. The nucleation role of TiO2 corresponds to the promotion of nucleation in volume, reversing the natural trend of surface nucleation for the parent glass. This is more easily achieved with additions of TiO2 in excess of approximately 7 mol %,50 close to the disappearance of incubation time of nucleation at 9.2-10.9 mol % TiO2.51 Surprisingly, our results indicate that there is no variation in the Ti local environment upon increase of the TiO2 content in the parent glasses (Figure 6). For all compositions, Ti is mainly present in a 5-fold coordinated site. Therefore, it seems that the Ti coordination is not involved in the explanation of a threshold in TiO2 content. We have previously shown that higher coordinated Al ([5]Al and [6] Al) is formed with the addition of TiO2, with an important increase in the fraction of the higher coordinated species around 7 mol % TiO2.52 Moreover, there is a preferential association of these high coordinated Al species with Ti, which yields to structural fluctuations that can correspond to preexisting seeds for nucleation. However, at this stage, we still cannot explain the threshold around 7 mol % TiO2. The investigation of the structural evolution of the Ti environment between the surface and the bulk of glassceramic specimen can shed light on the role played by Ti in the first stages of nucleation. At low TiO2 content (2 mol %), the Ti environment is not changed which would indicate that

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Figure 9. (A) Glass-ceramic MAS-6Ti measured by XAS at the Ti K-edge for different points along the line represented by the arrow. (B) Schematic representation of the glass-ceramic MAS-6Ti and the pre-edge features along the line of measurement.

Ti atoms remain in the glassy part of the materials, since crystals forming in this MAS systems should promote the formation of 6-fold coordinated Ti sites. The lack of changes in the pre-edges indicates that, for low TiO2 content, Ti atoms are not active in the nucleation process. We need to add at least 4 mol % TiO2 to observe a Ti coordination change with the formation of [6]Ti (Figure 7B). The presence of a higher proportion of [6]Ti just below the surface compared to the bulk suggests that Ti atoms enter the crystalline phases. Ti atoms can form MAT nanocrystals (not detected in the XRD pattern in Figure 4) or can be incorporated as substituting atoms in ~ β-quartz and indialite. The behavior at 6 mol % TiO2 presents an interesting evolution. Between 200 and 500 μm below the surface, in a region that is macroscopically crystallized, Ti pre-edge is more intense than in the initial glass and at a position lying in the middle of the [5]Ti domain. Then Ti environment is changed between 700 and 800 μm, a region that corresponds to the limit between the crystallized region and the remaining noncrystallized glass. The pre-edge is less intense and its position is on the top of the [6]Ti domain. The schematic picture in Figure 9 summarizes the observations for this sample. The last point, at 900 μm below the surface, corresponds to a point in the remaining glass, and, indeed, the extracted pre-edge parameters for this point (Table 4) are comparable to those for the MAS-6Ti parent glass. This indicates that no Ti coordination change appears in the glass before significant nucleation takes place. The bimodal behavior for Ti below the surface (200-500 and 700-800 μm) is difficult to interpret. The points at 700 and 800 μm present a small pre-edge intensity corresponding to [6]Ti as expected for Ti in MAT nanocrystals or in substitution in β-quartz and indialite. However, its position is at a low energy for [6]Ti alone, and we suspect that

[5]

Ti is still present in the front of nucleation. The changes in the pre-edge near the surface correspond to an increase in intensity which can only be understood by the presence of [4] Ti. As the position is different to that expected for [4]Ti alone, and since an important proportion of Ti must remain in the crystalline phases in octahedral position, there is likely a mixture of [4]Ti þ [6]Ti up to 500 μm below the surface. This suggests that not all Ti atoms are incorporated in the crystallization part and that Ti atoms in the remaining glass undergo a coordination change toward [4]Ti, an apparently stable coordination for Ti in glass. This coordination change may be the result of compositional changes and reorganization of the glass ceramic during the crystal growth and the transformation βquartz to indialite. In a recent TEM investigation using electron energy-loss spectroscopy (a technique similar to XANES), such compositional gradients have been evidenced as an Al2O3enriched layer around ZrTiO4 crystals,53 which could influence the further growth of secondary silicate crystalline phases. At higher TiO2 contents (8 and 10 mol %), Ti can be considered as an effective nucleating agent. This can be viewed by the Ti coordination change toward [6]Ti which is present in the MAT nanocrystals. As silicate phases develop, it is however still unclear if all Ti atoms remain in the MAT phases or if a part can substitute into the crystalline silicate phases. The pre-edge features are not modified between the surface and the bulk, which indicates a homogeneous distribution of Ti in the glass-ceramics with a common average environment. The major changes in Ti environment during nucleation are the local coordination change toward [6]Ti in the crystalline phases. However, [4]Ti is likely to be present in the remaining glassy part. This is observed in the MAS-6Ti glass where the surface that undergoes an incomplete nucleation contains a significant amount of [4]Ti, responsible for the high intensity in

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the pre-edge feature. This implies that Ti coordination may have important role in nucleation processes, and in particular the [5]Ti geometry could be the active sites for nucleation, contrary to [4]Ti sites. The reason lies in the specific association of [5]Ti sites with [5]Al and [6]Al sites, which results to few atomic rearrangements to form the MAT crystals during the heat treatment. A major thermodynamic barrier for nucleation thus corresponds to local coordination changes. This is emphasized in the surface zone of the MAS-6Ti where [4] Ti present in the glassy part prevents the nucleating role of Ti. Another remarkable result is the presence of a specific interface between the crystallized region and the vitreous bulk in the MAS-6Ti glass ceramic sample (Figure 9). This indicates that the structure is modified as the crystallization proceeds: there is a different Ti environment between the initial nucleation (probed at 800 μm at the glass-nucleation interface) and an advanced stage of crystallization (probed at 200 μm below the surface in the well-crystallized part). This implies that modifications of the glassy matrix, both structural and compositional changes, modify the pathways of nucleation. This emphasizes that not only the crystallized part but also the glassy part must be investigated to fully understand and control the formation of glass ceramics. Conclusions Using X-ray absorption spectroscopy at both Ti K- and L2,3-edges, we have investigated the role of TiO2 as a nucleating agent to promote internal nucleation as opposed to surface nucleation in glass ceramics and the influence of TiO2 content. We have successfully determined the Ti environment in the glass and in glass ceramics, from the surface to the bulk. We have shown that local coordination is not modified with the TiO2 content supporting the model that structural fluctuations at medium range are an important parameter controlling nucleation processes. We have for the first time elucidated a modification of the Ti environment between the nucleation front and the crystallized part, which correspond to structural rearrangements and compositional changes in the remaining glassy part. This is direct experimental evidence that some Ti sites, such as 4-fold coordinated Ti, should not be efficient reactive sites for nucleation. The study reported here provides a new approach to understand nucleation mechanism with structural information obtained by spatially resolved experiments. Acknowledgment. Authors wish to acknowledge Pierre Lagarde (SOLEIL), Anne-Marie Flank (SOLEIL), and C
ecile Jousseaume (SGR) for help in data acquisition. DSC thermograms were performed at Saint-Gobain Recherche with the valuable help from Samuel Pierre and Sophie Papin. G.S.H. acknowledges funding from NSERC in the form of a discovery grant. The Canadian Light Source is supported by NSERC, NRC, CIHR, and the University of Saskatchewan. This work was supported by the French national research agency (ANR) under contract no. 06-JCJC-0010 ‘Nuclevitro’.

References (1) H€ oland, W.; Rheinberger, V.; Schweiger, M. Phil. Trans. R. Soc. London A 2003, 361, 575–589. (2) Fokin, V. M.; Zanotto, E. D.; Yuritsyn, N. S.; Schmelzer, J. W. P. J. Non-Cryst. Solids 2006, 352, 2681–2714. (3) Barry, T. I.; Cox, J. M.; Morrell, R. J. Mater. Sci. 1978, 13, 594–610.

Cormier et al. (4) Dargaud, O.; Cormier, L.; Menguy, N.; Galoisy, L.; Calas, G.; Papin, S.; Querel, G.; Olivi, L. J. Non-Cryst. Solids 2010, 356, 2928-2934. (5) Beall, G. H. J. Eur. Ceram. Soc. 2009, 29, 1211–1219. (6) Knickerbocker, S. H.; Kumar, A. H.; Herron, L. W. Amer. Ceram. Soc. Bull. 1993, 72, 90–95. (7) Wange, P.; H€ oche, T.; R€ ussel, C.; Schnapp, J. D. J. Non-Cryst. Solids 2002, 298, 137–145. (8) H€ oche, T. J. Mater. Sci. 2010, 45, 3683–3696. (9) Dumas, T.; Ramos, A.; Gandais, M.; Petiau, J. J. Mater. Sci. Lett. 1985, 4, 129–132. (10) Ramos, A.; Gandais, M.; Petiau, J. J. Phys. (Paris) 1985, C8-Suppl No. 12, tome 46, C8-491-C8-494. (11) Greaves, G. N.; Bras, W.; Oversluizen, M.; Clarck, S. M. Faraday Discuss. 2002, 122, 299–314. (12) Dargaud, O.; Calas, G.; Cormier, L.; Galoisy, L.; Jousseaume, C.; Querel, G.; Newville, M. J. Am. Ceram. Soc. 2010, 93, 342–344. (13) Meneghini, C.; Mobilio, S.; Lusvarghi, L.; Bondioli, F.; Ferrari, A. M.; Manfredini, T.; Siligardi, C. J. Appl. Crystallogr. 2004, 37, 890–900. (14) Dumas, T.; Petiau, J. J. Non-Cryst. Solids 1986, 81, 201–220. (15) Waychunas, G. A. Am. Mineral. 1987, 72, 89–101. (16) Farges, F.; Brown, G. E., Jr.; Navrotsky, A.; Gan, H.; Rehr, J. J. Geochim. Cosmochim. Acta 1996, 60, 3039–3053. (17) Henderson, G. S.; Liu, X.; Fleet, M. E. Phys. Chem. Miner. 2002, 29, 32–42. (18) Sandstrom, D. R.; Lytle, F. W.; Wei, P. S. P.; Greegor, R. B.; Wong, J.; Schultz, P. J. Non-Cryst. Solids 1980, 41, 201–207. (19) Greegor, R. B.; Lytle, F. W.; Sandstrom, D. R.; Wong, J.; Schultz, P. J. Non-Cryst. Solids 1983, 55, 27–43. (20) Mountjoy, G.; Pickup, D. M.; W., W. G.; Anderson, R.; Cole, J. M.; Newport, R. J.; Smith, M. E. Chem. Mater. 1999, 11, 1253– 1258. (21) Dingwell, D. B.; Paris, E.; Seifert, F.; Mottana, A.; Romano, C. Phys. Chem. Miner. 1994, 21, 501–509. (22) Ponader, C. W.; Boeck, H.; Dickinson, J. E. J. Non-Cryst. Solids 1996, 201, 81–94. (23) Farges, F. J. Non-Cryst. Solids 1999, 244, 25–33. (24) Farges, F. Am. Mineral. 1997, 82, 36–43. (25) Yarker, C. A.; Johnson, P. A. V.; Wright, A. C.; Wong, J.; Greegor, R. B.; Lytle, F. W.; Sinclair, R. N. J. Non-Cryst. Solids 1986, 79, 117–136. (26) Cormier, L.; Gaskell, P. H.; Calas, G.; Soper, A. K. Phys. Rev. B 1998, 58, 11322–11330. (27) Della Gatta, G.; Richardson, M. J.; Sarge, S. M.; Stolen, S. Pure Appl. Chem. 2006, 78, 1455–1476. (28) Raabe, J.; Tzvetkov, G.; Flechsig, U.; Boge, M.; Jaggi, A.; Sarafimov, B.; Vernooij, M. G. C.; Huthwelker, T.; Ade, H.; Kilcoyne, D.; Tyliszczak, T.; Fink, R. H.; Quitmann, C. Rev. Sci. Instr. 2008, 79, 113704. (29) Flank, A. M.; Cauchon, G.; Lagarde, P.; Bac, S.; Janousch, M.; Wetter, R.; Dubuisson, J. M.; Idir, M.; Langlois, F.; Moreno, T.; Vantelon, D. Nucl. Instr. Methods Phys. Res. B 2006, 246, 269–274. (30) Ravel, B.; Newville, M. J. Synchrotron Rad. 2005, 12, 537–541. (31) Weaver, D. T.; Van Aken, D. C.; Smith, J. D. Mater. Sci. Eng. A 2003, 339, 96–102. (32) Goel, A.; Shaaban, E. R.; Melo, F. C. L.; Ribeiro, M. J.; Ferreira, J. M. F. J. Non-Cryst. Solids 2007, 353, 2383–2391. (33) Degroot, F. M. F.; Fuggle, J. C.; Thole, B. T.; Sawatzky, G. A. Phys. Rev. B 1990, 41, 928–937. (34) Degroot, F. M. F.; Figueiredo, M. O.; Basto, M. J.; Abbate, M.; Petersen, H.; Fuggle, J. C. Phys. Chem. Miner. 1992, 19, 140–147. (35) Crocombette, J. P.; Jollet, F. J. Phys.: Condens. Matter 1994, 6, 10811–10821. (36) Brydson, R.; Sauer, H.; Engel, W.; Thomas, J. M.; Zeitler, E.; Kosugi, N.; Kuroda, H. J. Phys.: Condens. Matter 1989, 1, 797–812. (37) H€ oche, T.; Grodzicki, M.; Heyroth, F.; van Aken, P. A. Phys. Rev. B 2005, 72, 205111. (38) Guignard, M.; Cormier, L.; Montouillout, V.; Menguy, N.; Massiot, D.; Hannon, A. C. J. Phys.: Condens. Matter 2009, 21, 10. (39) H€ oche, T.; van Aken, P. A.; Grodzicki, M.; Heyroth, F.; Keding, R.; Uecker, R. Phil. Mag. 2004, 84, 3117–3132. (40) Golubkov, V. V.; Dymshits, O. S.; Zhilin, A. A.; Chuvaeva, T. I.; Shashkin, A. V. Glass Phys. Chem. 2003, 29, 254–266. (41) Pinckney, L. R.; Beal, G. H J. Non-Cryst. Solids 1997, 219, 219–227.

Article (42) Guignard, M.; Cormier, L.; Montouillout, V.; Menguy, N.; Massiot, D.; Hannon, A. C.; Beuneu, B. J. Phys.: Condens. Matter 2010, 22, 7. (43) Yang, H.; Hazen, R. M. J. Solid State Chem. 1998, 138, 238–244. (44) Norberg, S. T.; Ishizawa, N.; Hoffmann, S.; Yoshimura, M. Acta Crystallogr. E 2005, E61, i160–i162. (45) Farges, F.; Brown, G. E., Jr.; Rehr, J. J. Geochim. Cosmochim. Acta 1996, 60, 3023–3038. (46) Park, K.-H.; Shin, D.-W. J. Ceram. Proc. Res. 2002, 3, 153–158. (47) Zdaniewski, W. J. Mater. Sc. 1973, 8, 192–202. (48) Farges, F. Phys. Rev. B 1997, 56, 1809–1819.

Crystal Growth & Design, Vol. 11, No. 1, 2011

319

(49) Henderson, G. S.; Liu, X.; Fleet, M. E. Mineral. Mag. 2003, 67, 597–607. (50) Fokin, V. M.; Zanotto, E. D. J. Non-Cryst. Solids 1999, 246, 115–127. (51) Loshmanov, A. A.; Sigaev, V. N.; Khodakovskaya, R. Y.; Pavlushkin, N. M.; Yamzin, I. I. J. Appl. Crystallogr. 1974, 7, 207–210. (52) Guignard, M.; Cormier, L.; Montouillout, V.; Menguy, N.; Massiot, D. J. Non-Cryst. Solids 2010, 356, 1368–1373. (53) Bhattacharyya, S.; H€ oche, T.; Jinschek, J. R.; Avramov, I.; Wurth, R.; Muller, M.; Russel, C. Cryst. Growth Des. 2010, 10, 379–385.