V Aubin
email: vaubin@ext.jussieu.fr

F Studer

D Caurant

D Gourier

N Nguyen

A Ducouret

N Baffier

T Advocat

The stability under irradiation of hollandite ceramics, specific radioactive cesium-host wasteforms

Investigations are currently performed on matrices for the specific immobilization of long-lived radionuclides such as fission products resulting from an enhanced reprocessing of spent fuel. Hollandite (nominally BaAl 2 Ti 6 O 16 ), one of the phases constituting Synroc, receives renewed interest as specific Cshost wasteform. The radioactive cesium isotopes decay involves the emission of β particles, γ rays and the transmutation of Cs to stable Ba ions. This study deals with the synthesis of hollandite ceramics by oxide route and single crystals by a flux method having the Ba x Cs y (Al,Fe) 2x+y Ti 8-2x-y O 16 composition type (1≤x≤1.28; 0≤y≤0.28). The influence of the hollandite chemical composition on the hollandite structure and microstructure is studied. To estimate the hollandite radiation resistance, external electron irradiation experiments, simulating the β particles emitted by radioactive cesium, were carried on single phase materials. The radiation effects were characterized by electron paramagnetic resonance (EPR) and Mössbauer spectroscopy.

INTRODUCTION

Over the past few years, progress on the enhanced reprocessing of nuclear spent fuel and notably on separating the long-lived fission products from the high level radioactive wastes (HLW) has encouraged the investigations on specific host matrices [1]. Particularly, the success of cesium extraction by calixarenes complexes, except for some traces of rubidium, suggests that new materials with higher performances than borosilicate glassy matrices currently used to immobilize unseparated HLW, could be considered. Cesium appears as one of the most hazardous radionuclide to immobilize because of the high thermal power of 134 Cs and 137 Cs isotopes, the long lifetime of 135 Cs isotope (T 1/2 =2,3 10 6 years), its volatility at high temperature during waste forms synthesis and its tendency to form water-soluble compounds. The cesium radioactive isotopes decay into stable barium by emitting β particles (with energy E ranging from 0.09 to 1.17 MeV, mainly 0.5 MeV) and γ rays (E≈0.6 MeV). Therefore, the choice of new Cshost matrices rely on their waste (Cs, Rb) loading, their Ba and Sr (decay product of Rb) incorporation ability, their thermal characteristics, their long-term containment properties (chemical durability, βγ radiation resistance) and the ease of their synthesis.

Among several Cs-specific matrices proposed (refractory aluminosilicate glasses [2], zeolites [3], pollucite ceramics [4]), hollandite of [Ba 2+

x Cs + y ][(Al 3+ ,Ti 3+ ) 2x+y Ti

O33-01

composition, the A cations and vacant sites may take up an ordered arrangement along the tunnels (superlattice ordering). Hollandite has several advantages. It is able to accommodate chemistry and valence changes due to the cesium transmutation because of its high capacity to incorporate Ba in its structure and due to the possibility of reduction of a fraction of Ti 4+ ions to ensure the charge balance [5]. Moreover, its chemical durability is two orders of magnitude higher than the currently used borosilicate glasses, which is required to select Cs-specific matrices [6]. Nevertheless, the synthesis of [Ba 2+

x Cs + y ][(Al 3+ ,Ti 3+ ) 2x+y Ti 4+ 8-2xy ]O 16 composition, named afterwards the Synroc composition, remains actually too complex to be industrialized [START_REF] Kesson | Radioact. Waste Manage[END_REF]. Indeed, its synthesis consists of a hot pressing stage under controlled redox conditions of a precursor obtained by an alkoxide route. The substitution of Ti 3+ by Fe 3+ in the Synroc composition enables to obtain single phase (Ba,Cs)-hollandite using a natural sintering under air of a precursors mixture prepared by an alkoxide route [8], which clearly simplifies the protocol. The next step is to use the oxide route to get these hollandite ceramics, which will be described in this study. On the other hand, the βγ radiation resistance of hollandite is not yet proved due to the lack of both true natural analogs and experiments on synthetic hollandite samples containing radioactive cesium. Moreover, literature reports only electron external irradiation experiments which overestimate the real effects of β irradiation. Indeed, only experiments under intense electron irradiation in electron microscope were carried out on single phase hollandite, which generated microtwinnings followed by amorphization [9]. Other external irradiations were performed on hollandite samples, simulating the α irradiation which could come from the α-emitting surrounding actinide-contaning phases (such as in Synroc assemblage) [10]. However, these radiation effects are not to be considered in the case of Csspecific immobilization and can not be extrapolated to predict the hollandite βγradiation resistance.

In this work, we report the synthesis of hollandite ceramics in the Ba x Cs y (Al,Fe) 2x+y Ti [8] and the insertion of cesium in hollandite structure (increase of channels size) [START_REF] Kesson | Radioact. Waste Manage[END_REF]. We study this composition range by varying the Al/Fe and Cs/Ba ratios (Table I). The reagent powders (almost 30g) were mixed, ground in an agate mortar, pelletized and calcined under air for 4 hours at 810°C. The calcined powder was then ground with an attrition mill with zirconia-based glass-ceramics or yttria zirconia (indicated with * near the composition sample in Table I) balls in water for 1 hour at 550 rpm (rounds per minute). This fine grinding stage is absolutely necessary to synthesize (Ba, Cs)-hollandite because it increases the precursor homogeneity and reactivity and so decreases both sintering temperature and duration. Once the precursor dried (at 100°C for 1-2 days), it was pelletized and sintered at 1200°C for 30 hours under air. The optimization of this synthesis was already reported in [11]. Samples were characterized by X-ray diffraction (XRD) with CoKα wavelength (λ= 1.78897 Å), by Scanning Electron Microscopy (SEM), Energy Dispersive X-ray Analysis (EDX) and by Electron Probe MicroAnalysis (EPMA). For enough densified samples, the ceramic density d was deduced from the Archimede force applied in toluene solvent and was then compared to the theoretical density d th deduced from the lattice parameters extracted from XRD patterns. Hollandite single crystals of the most simple composition Ba 1.16 Al 2.32 Ti 5.68 O 16 were also synthesized by a flux method. The choice of the crystal growth method results from the incongruent melting of hollandite (as verified for the Ba 1.16 Al 2.32 Ti 5.68 O 16 composition by both differential thermal analysis experiments and slow cooling from the melt). This new synthesis protocol is precisely described in [12] End-member Ba-hollandite (y=0) materials were irradiated by 1, 1.5 or 2.

RESULTS AND DISCUSSION

XRD and SEM Characterizations of Hollandite Ceramics

The hollandite ceramics with C=Al (samples 1 and 2) were first synthesized using the protocol described above. The Ba-Al-hollandite was obtained as single phase and was relatively dense (d/d th =95%, Fig. 2a.). Its crystal structure is I4/m with the cell parameters indicated in Table II. Then, we attempted to insert 0.1% Cs/formula unit, keeping constant the cations in octahedral sites. This Cs-content was chosen relatively low because literature indicates an upper limit of 0.25% Cs/formula unit for these type of hollandite ceramics prepared under air [13].

Nevertheless, this material was multiphase (Table II) and not densified (Fig. 2d). EPMA (Table I) indicates that only half of cesium is incorporated in hollandite. Therefore, a fraction of Cs volatilized or was concentrated in a secondary phase not detected by SEM due to the high porosity of the material. The second hypothesis appears more probable because after a longer sintering treatment (96h), a Cs-rich secondary phase containing Al, Ti, Si cations (of too small size to be probed by EPMA to get its exact composition) was observed. We could suppose that this phase is the well-known leachable CsAlTiO 4 phase already reported in literature [START_REF] Kesson | Radioact. Waste Manage[END_REF] with a partial substitution of Ti by Si, by comparison with other synthesized materials not reported here. The occurrence of this phase seems to be due to kinetic factors which could be reduced by the partial substitution of Al 3+ by bigger trivalent cations [START_REF] Kesson | Radioact. Waste Manage[END_REF]. Moreover, this phase could be favored by the presence of Si impurities which originate in part from the TiO 2 raw material and from the attrition balls constituted of zirconia-based glassceramics.

ATALANTE
To facilitate Cs incorporation in hollandite structure, Al 3+ was substituted partially or totally by Fe 3+ . The BaFe-hollandite (samples 3 and 5) are almost single phase except for a small white (backscattered electrons) phase enriched in Ba, P and Si only detected by SEM. This parasitical phase has been almost avoided afterwards by using very pure rutile as raw material. This impurities issue was already reported in [14]. These two ceramics were relatively dense (d/d th =92,7% and 96% for the samples 3 and 5, respectively). By comparing their microstructures (Fig. 2.b and c), Al ions do not appear as a sintering reagent because the amount of porosities decreases when the Al/Fe ratio decreases. This Al-Fe substitution leads also to an increase of the lattice parameters a, c and of the unit-cell volume V (Table II), which is explained by the difference of ionic radii of Al 3+ (r=0.535 Å) and Fe 3+ (r=0.645 Å) [15]. Thus, a better insertion of Cs is expected in hollandite containing iron. Two different Cs contents (0.28 and 0.24/unit-cell Cs for samples 4 and 6 respectively) were incorporated because we want to synthesize exactly the Cs 0.28 Ba 1 Al 1.46 Fe 0.82 Ti 5.72 O 16 composition reported in [8] to compare directly the alkoxide and oxide routes. Sample 4 is a single phase material with a composition very close to the expected one (4.45 wt% Cs 2 O versus 5 wt% initially introduced, Table I). Unfortunately, this material has a high porosity as observed by SEM (Fig. 2e). Comparison with the BaFe-hollandite (Fig. 2b) microstructure shows that cesium prevents also a good densification of the ceramic. Sample 6 retains the entire cesium initially introduced (Table I) and is dense. However, traces of Fe 2 TiO 5 and BaTi 4 O 9 could be suspected by XRD (only 1-2 peaks), which does not raise issue because this secondary phase does not concentrate cesium. Increasing the time sintering would improve the material homogeneity. The densification of this ceramic is better than the sample 4 one (Fig. 2e,f), which confirms that Al ions do not favor material densification. As shown by the lattice parameters of samples 3, 4, 5 and 6 (Table II), the incorporation of Cs induced an increase of a and V due to the large size of this cation, without changing c parameter. Therefore, single phase (Ba,Cs)-hollandite ceramics could be synthesized by oxide route using natural sintering under air with a total retention of the cesium initially introduced. The best composition currently synthesized seems to be Cs 0.24 Ba

Electron Irradiation Effects

Mössbauer Study

The sample 5 of composition Ba 1.16 Fe 2.32 Ti 5.68 O 16 was selected in this study because no 57 Feenrichment was necessary and this composition is relatively simple with only one trivalent specie (Fe 3+ ) and one tunnel cation (Ba 2+ ). Its 300K Mössbauer spectrum (Fig. 3a) exhibits one quadrupolar doublet before irradiation. However, in order to obtain the best fit, it was necessary to consider two Mössbauer components whose parameters are indicated in Table III. Their isomer shift IS values (0.49 mm/s) are typical of Fe 3+ ions. Only the quadrupole splitting (QS) differentiate these two sites. It agrees well with the iron octahedral environment with a more important distortion for site 1. Therefore, no Fe 2+ ions are detected before irradiation.

In the tetragonal hollandite structure (I4/m space group), all the octahedral sites are equivalent and Fe 3+ and Ti 4+ ions are assumed to be randomized in them in the scale of the X-ray diffraction [16].

On the other hand, if the structure is monoclinic (I2/m space group), two crystallographic sites occur but without preferential occupation [16]. However, few works are reported about local spectroscopies such as NMR, EPR and Mössbauer spectrocopies. Nevertheless, two Mössbauer studies on K 1.45 Fe 1.45 Ti 6.55 O 16 [17]and Cs 1.40 Fe 1.40 Ti 6.74 O 16 [18] hollandites report also one quadrupolar doublet at 300K whose parameters seem to be closer to the sample 5 site 2 ones. Moreover, literature reports a monoclinic structure for a close hollandite composition Ba 1.14 Fe 2.28 Ti 5.72 O 16 [14]. Therefore, we can wonder whether our material does not display a slight monoclinic distortion not clearly detected by XRD and thus these two doublets (Fig. 3a) would be related to the two different crystallographic sites. This hypothesis is rejected by comparison with Mössbauer spectra of the samples 3 and 4 (I4/m space group) which display also 2 quadrupolar doublets, not shown here [19]. The crystal structure of these samples was confirmed by structural determination on single crystals of the same composition [12]. Therefore, only one crystallographic site is assumed for sample 5 and then these two different doublets could be assigned to different local environments of Fe 3+ ions. First, this difference could be due to the possibility to have Fe 3+ ions or Ti 4+ ions as nearest neighbors. This explanation does not seem to be probable given the supposed random distribution of the Fe and Ti ions in the octahedral sites and the presence of only two quadrupolar doublets in samples 3 and 4 containing in addition Al ions [19]. Therefore, the occurrence of these two sites could only be explained by the occupancy of cation site in the tunnels. Indeed, this occupancy is not full and thus Fe 3+ ions could have vacancies or Ba cations in their close environment (Fig. 1). Moreover, it is known that B cations (Fe 3+ notably) change their position in their octahedral site according to the local tunnel content [20], which could explain the different quadrupolar splitting of the two sites. This hypothesis is confirmed by the changes of the site fraction R with the hollandite chemical composition [19]. The exact environment of each site is nevertheless not yet fully determined. Although 300K experimental Mössbauer spectrum of the electron irradiated sample 5 exhibits only one quadrupolar doublet (Fig. 3b), 3 iron components are necessary to reproduce it. The two first sites correspond to those observed before irradiation as shown by their associated parameters in Table III. Nevertheless, their relative fraction change in favor of the site 1 which becomes also more symmetric. This suggests that this irradiation modify the local arrangement of Ba cations around Fe 3+ ions by atomic displacements of Ba ions. Moreover, a third site was generated in small quantities. Its smaller isomer shift and higher quadrupole splitting values, with respect to the two first sites, are more typical of a fivefold coordinated pyramidal Fe 3+ ions. Therefore, this electron irradiation generates also some oxygen displacements. Finally, no change of the iron oxidation seems to occur: no Fe 2+ ions were detected within the Mössbauer sensitivity scale. Briefly, Mössbauer results show that this electron irradiation (E=1.5 MeV, F=5.8 10 18 cm - 2 ) seems to generate atomic displacements of at least Ba and O ions without reduction of a significant fraction of the Fe 3+ ions.

EPR study

The EPR study was carried out on ceramics and single crystals of the Ba 1.16 Al 2.32 Ti 5.68 O 16 composition to avoid any interference between the irradiation-induced defects and Fe 3+ ions (paramagnetic species). Fig. 4 exhibits the typical EPR spectra of pristine and β-irradiated ceramics as a function of fluence and energy of electrons. The EPR spectra of irradiated samples display three new signals T r , E 1 and E 2 whose intensities increase with increasing fluence without saturation up to 1.4 10 19 cm -2 for 1 MeV electron irradiations. Their relative proportions depend not only on fluence but also on electron energy. Indeed, at almost the same fluence (5. 10 18 cm -2 ), the E 2 signal dominates the spectrum of the 2.5 and 1.5 MeV irradiated samples whereas it is weak for the sample irradiated by 1 MeV electrons. This influence of the electron energy suggests an effect of energy displacement threshold and obviously a mechanism of ballistic collisions associated with this signal E 2 . In order to attribute and quantify these signals, experimental spectra, recorded at different temperatures, were simulated [21]. The g-tensor and the linewidth associated to each signal are indicated in Table IV. Thus, two types of signals could be differentiated:

• The T r signal , located near 337 mT corresponding to g-factor larger than the free electron g-factor (g e =2.0023). Its position imply that the associated defects are hole centers. This signal is simulated considering the contribution of three orthorhombic sites differing only from the g z value (Table IV) to account for the shoulders on the left-hand side of the signal. The computed g-factors are typical of oxygen species [START_REF] Marfunin | Spectroscopy, luminescence and radiation centers in Minerals[END_REF]. However, this signal cannot be assigned to a O -center but rather to a superoxide ion O 2 according to its g-shift sequence g // = g z > g ⊥ =1/2(g x +g y ) > g e . The associated spin concentration is in the range 4.2 10 15 -5.4 10 16 cm -3 when the fluence of 1 MeV electron is increased from 8.1 10 17 to 1.4 10 19 cm -2 .

Table IV. g-factors (g x , g y and g z components) and corresponding half-width at half-maximum (∆B mT) determined by the simulation of the three signals detected in ceramics after electron irradiation. All simulations were performed using gaussian shape lines. ) were also irradiated. First, to confirm that the same signals were generated in ceramics and single crystals, the EPR powder spectrum of several ground single crystals was compared to the ceramic one (Fig. 5). This spectrum (Fig. 5b) displays also the three signals T r , E 1 and E 2 with only small lineshape differences which could simply be explained by the lack of some crystallite orientations versus the external magnetic field in the ground crystal powder. Their relative intensities are also very O33-01 occupy a site with quasi -axial symmetry . Thus, its variation in the plane (110) could be fitted by Eq.1 where Θ is the angle between the external magnetic field and the principal axis of g-tensor.

) ( sin ) ( cos The principal values g ⊥ =1.8868, g // =1.9646 are thus obtained, with z following the c axis. These g values agree well with one of the two possible simulations of the signal E 2 observed in ceramics (noted signal E 2 (b) in Table IV), which allows us to eliminate the other one. Moreover, these results show that the principal axis of the gtensor follows the direction Ti-Ti along the [TiO 6 ] octahedra chains. Moreover, this axial site symmetry is quite surprising given the important distortion of the octaedral sites . We can therefore wonder if this center would not be located in interstitial sites rather than in normal sites. In this latter case, it imply irradiationinduced titanium displacements.

CONCLUSIONS

Thus, single phase hollandite ceramics in the Ba x Cs y (C=Al,Fe) 2x+y Ti 8-2x-y O 16 compositional range could be obtained by natural sintering under air of oxide powders only when iron is introduced. Nevertheless, the ceramic without Al seems to be the best one in terms of density and cesium retention. The behavior of several hollandite under external electron irradiation simulating the β-irradiation of radioactive cesium were also studied by EPR and Mössbauer spectroscopy. Both spectroscopies tend to show that atomics displacements of hollandite ions (O, Ba, B) could be generated, which modified the local environment of probed titanium and iron ions.

4+ 8 -Fig. 1 .

 81 Fig. 1. View of the tetragonal (I4/m) hollanditetype A 2 B 8 O 16 structure approximately along the c-axis: A cations (black atoms) and [TiO 6 ] octahedra are shown .

5

 5 MeV electrons generated by a Van de Graaff accelerator to simulate the effect of β particles emitted by radioactive cesium. The thickness of samples (0.5 or 1 mm) was chosen so that electrons go through the sample and radiations effects are quite homogeneous. Different fluences between 8.1 10 17 and 1.4 10 19 cm -2 were carried out. As electron irradiations generate mainly electron excitations of materials and few atomic displacements, the techniques used to characterize the radiation effects were sensitive local spectroscopies: EPR , Mössbauer spectroscopy, depending on samples composition. The 57 Fe transmission Mössbauer spectra were measured at room temperature using a 57 CoRh source on Fe-rich powders or thin massive samples. The isomer shift (I.S.) are corrected for the I.S. of metallic iron taken as a reference. The spectra were simulated thanks to a least square fitting routine assuming a Lorentzian line shape. The EPR spectra were recorded on a Bruker ESP 300e or an ELEXYS E500 spectrometer operating at X band (9.4 GHz) in the range of temperature 6-300K. The EPR signals were simulated using the software SIMFONIA. The spin concentration associated with EPR signals was determined by comparison with the organic radical standard DPPH. The main study was performed on Ba-Al hollandite (C=Al) because the occurence of Fe 3+ (paramagnetic 3d 5 ion) induced a strong and wide contribution on EPR spectra that could prevent the detection of irradiation-induced defects.

Fig. 2 .

 2 Fig. 2. Backscattered electron SEM micrographs of Ba-hollandite 1 (a), 3 (b), 5(c) and (Cs,Ba)hollandite 2 (d), 4 (e), 6 (f) showing notably the difference of porosity (darken areas) between samples (all the samples were prepared following the same method). TABLE II. Polyphased character of the hollandite ceramics. Lattice parameters a, c and unit-cell volume V (space group I4/m) are given for almost single phase samples. Lattice parameters Secondary phases (XRD and SEM) a (Å) c (Å) V (Å 3 ) 1 -9.968 2.923 290.5 2 Al 2 O 3 ,TiO 2 ,Ba 2 Ti 9 O 20 ---3 Impurities phase (traces) 10.003 2.944 294.5 4 -10.042 2.942 296.7 5 Impurities phase (traces) 10.103 2.971 303.2 6 Fe 2 TiO 5 (traces) 10.122 2.972 304.5

Fig. 3 .

 3 Fig. 3.Typical 300K Mössbauer spectra of sample 5 (Ba 1.16 Fe 2.32 Ti 5.68 O 16 ) before (a) and after irradiation (b) with 1.5 MeV electrons at a fluence F 5.8 10 18 cm -2 . The components used to simulate the experimental spectra are indicated by solid lines. The difference between the experimental and simulated spectra are shown below each spectrum.

Fig. 4 .

 4 Fig. 4. Typical EPR spectra of Ba 1.16 Al 2.32 Ti 5.68 O 16 hollandite before and after electron irradiation recorded at 70K with microwave power P =20 mW and microwave frequency ν =9.5 GHz. Energy of electrons and fluences are indicated near the spectra.

Fig. 5 .Fig. 6 .

 56 Fig. 5. Typical EPR spectra of Ba 1.16 Al 2.32 Ti 5.68 O 16 ceramics (a) and ground crystals (b), both irradiated by 1.5 MeV at a fluence of 5.8 10 18 cm -2 , recorded at 70K with P =20 mW and ν =9.4 GHz.

ATALANTE 2004 Nîmes (France) June 21-25, 2004 2 O33-01

  Ba 1.16 Al 2.32 Ti 5.68 O 16 Ba 1.18 Al 2.28 Ti 5.69 O 16 2 Cs 0.1 Ba 1.11 Al 2.32 Ti 5.68 O 16 Cs 0.05 Ba 1.17 Al 2.30 Ti 5.67 O 16 3 Ba 1.28 Al 1.64 Fe 0.92 Ti 5.44 O 16 Ba 1.29 Al 1.71 Fe 0.93 Ti 5.38 O 16 4 Cs 0.28 Ba 1 Al 1.46 Fe 0.82 Ti 5.72 O 16 Cs 0.25 Ba 1.05 Al 1.43 Fe 0.80 Ti 5.74 O 16 5 Ba 1.16 Fe 2.32 Ti 5.68 O 16 Ba 1.13 Fe 2.32 Ti 5.70 O 16 6 Cs 0.24 Ba 1.04 Fe 2.32 Ti 5.68 O 16 Cs 0.26 Ba 1.06 Fe 2.28 Ti 5.70 O 16

	ceramic powder at 40% wt. This mixture was
	melted at 1400°C and then slowly cooled to
	room temperature (3 to 12°C/h).
	TABLE I. theoretical and EPMA compositions of
	hollandite ceramics.
	Expected compositions EPMA compositions
	1
	: flux (47%
	mol. BaF 2 -53% mol B 2 O 3 ) added to hollandite

*

  1.04 Fe 2.32 Ti 5.68 O 16 in terms of microstructure (lowest porosity). Moreover, replacing entirely Al 3+ ions obviously reduces the amount of raw materials and avoids also the formation of the leachable CsAlTiO 4 phase. Thus, if Al ions do not influence the hollandite chemical durability, it appears useless to introduce these trivalent ions. The first measurements on the sample 6, currently in progress, tend to show that its chemical durability is as good as the hollandite ceramic reported in[8] and so that Al ions does not take part in the hollandite high chemical durability. Therefore, the Cs 0.24 Ba 1.04 Fe 2.32 Ti 5.68 O 16 hollandite composition is actually selected for this oxide synthesis.

TABLE III

 III

	sample	sites	I.S	Q.S.	Γ	R
			mm/s.	mm/s	mm/s	%
	Pristine	1	0.493 0.884	0.25 63
	sample 5	2	0.494	0.56	0.19 37
	Irradiated	1	0.490	0.76	0.23 77
	sample 5	2	0.492	0.55	0.12 19
		3	0.36	1.21	0.23	4

. Hyperfine parameters extracted from the least-squares fits of the 300K 57 Fe spectra of the pristine and electron irradiated sample 5 shown in Fig.

3

: Isomer shift (I.S.), quadrupole splitting (Q.S.), full width at half maximum (Γ) and fraction (R) of the different sites.

ACKNOWLEDGEMENTS

The authors are indebted to the DSM/DRECAM/LSI laboratory ( Palaiseau, France) and particularly to Stéphane Esnouf for the runs at the Van de Graaff accelerator.


P Loiseau
email: pascal-loiseau@enscp.jussieu.fr

D Caurant

K Dardenne

S Mangold

M Denecke

J Rothe

N Baffier

C Fillet

Characterisation of Nd-doped calcium aluminosilicate parent glasses designed for the preparation of zirconolite-based glass-ceramic waste forms

Zirconolite-based (nominally CaZrTi 2 O 7 ) glass-ceramics belonging to the SiO 2 -Al 2 O 3 -CaO-ZrO 2 -TiO 2 system are good waste forms for the specific immobilisation of actinides. The understanding of their crystallisation processes implies to investigate the structure of the glass. Thus, the environment around Ti, Zr (nucleating agents) and Nd (trivalent actinides surrogate) was characterised in parent glasses. Electron spin resonance (ESR) study of the small amount of Ti 3+ occurring in the glass enabled to identify two types of sites for titanium: the main one is of C 4v or D 4h symmetry. EXAFS showed that Zr occupied a quite well defined 6-7-fold coordinated site with second neighbours which could correspond to Ca/Ti and Zr. Nd environment was probed by optical spectroscopies (absorption, fluorescence), ESR and EXAFS. All these techniques demonstrated that the environment around Nd was very constrained by the glassy network. Notably, Nd occupies a highly distorted 8-9-fold coordinated site in the parent glass.

INTRODUCTION

Zirconolite (CaZrTi 2 O 7 ) is a naturally occurring phase which exhibits an excellent chemical durability and a good actinide containment capacity. Therefore, this crystalline phase is a waste form of choice for the specific immobilisation of minor actinides (mainly responsible for the long-term radiotoxicity of reprocessed high level nuclear waste) or of plutonium. An alternative to the ceramic route is to prepare zirconolite crystals by controlled crystallisation (nucleation+growth) of a parent glass. This glass-ceramic route is simpler to implement in radioactive plants and offers a greater chemical flexibility against waste composition fluctuations due to the presence of a residual glass surrounding zirconolite crystals. The basic parent glass composition which was studied is as follows (mol. %): 48.8 SiO 2 , 8.5 Al 2 O 3 , 25.3 CaO, 11.3 TiO 2 , 5.0 ZrO 2 , 1.1 Na 2 O. This composition was determined from [1] and various loadings of Nd 2 O 3 (trivalent actinide surrogate) were added up to 2.2 mol. % (10 wt. %). Previous studies showed that such glasses led to the only crystallisation of zirconolite in their bulk by controlled crystallisation (nucleation + crystal growth) [2]. However, the understanding of crystallisation processes implies to investigate the structure of the glass if further composition optimisations are to perform. This is why the environment around three elements was characterised in these parent glasses: Ti and Zr (which play a key role in nucleation processes and which are main components of zirconolite), and Nd.

EXPERIMENTAL

40g batch glasses were prepared following a standard method comprising two stages of melting and casting at 1550°C with intermediate grinding in order to increase glass homogeneity. Their compositions were checked by ICP-AES and they were perfectly amorphous. A similar glass transformation temperature around 760 °C was found for all the glasses (measured by differential thermal analysis), which means that the addition of Nd 2 O 3 up to 2.2 mol. % has no significant effect on the mean bond strength of the glassy network. Due to the high preparation temperature of the glasses, a small quantity of Ti 4+ ions are reduced into Ti 3+ state. Contrarily to Ti 4+ ions, Ti 3+ ions (3d 1 electron configuration) can be detected by Electron Spin Resonance (ESR). Ti 3+ ESR measurements were performed at 120 K using a Bruker ESP 300e spectrometer operating at Xband (ν ≈ 9.5 GHz). Room temperature EXAFS (Extended X-ray Absorption Fine Structure) measurements were performed at Zr-K edge (17998 eV) and Nd-L 3 edge (6208 eV) at ANKA synchrotron (Germany, Karlsruhe) in transmission mode through quenched glass samples containing 1.3 mol. % of Nd 2 O 3 (6 wt. %) dispersed in boron nitride pellets. EXAFS signals were extracted and simulated using the UWXAFS package [3]. Nd 3+ optical properties were characterised at 15 K by absorption from 400 to 900 nm and by emission from 860 to 960 nm ( 4 F 3/2 → 4 I 9/2 channel). At 15 K, only the lowest energy Kramers doublet of the 4 I 9/2 or of the 4 F 3/2 multiplet is populated for absorption or emission measurements respectively.

RESULTS AND DISCUSSION

ESR study of native Ti 3+ ions in parent glass

During glass meting, a Ti 4+ ⇔ Ti 3+ equilibrium was established (displaced towards reduction with increasing temperature) and then was frozen by quenching at room temperature. Therefore, a small quantity of Ti 3+ , detected by ESR (Fig. 1), formed. Their amount was quantified by ESR at 120 K with a DPPH standard. For all the glasses, it was shown that the proportion of Ti at the trivalent oxidation state was below 0.7 ppm (case of the undoped glass) and decreased with increasing Nd 2 O 3 concentration (no Ti 3+ ion detected for the 10 wt. % Nd 2 O 3 glass). Such an evolution is difficult to explain (Nd only exists as Nd 3+ ion) but could result from an increase of the glass basicity [4] with Nd concentration, if we admit that Nd acts as a network modifier. Ti 3+ ESR signal was simulated with the assumption of an axial symmetry for simplicity reasons. Without taking into account any distribution of g factor, it was necessary to consider two contribution (with gaussian line shape) to simulate the signal and the shoulder occurring around 344 mT. This gives evidence of two types of sites for Ti 3+ ions in the glass, referred as site 1 and site 2. Results of the simulation for the undoped glass are shown in Fig. 1 and Table I. Site 2 contributes only to 8.5 % of the overall signal. The knowledge of the g factors gives pieces of information about the symmetry of Ti 3+ environment in the glass [START_REF] Wertz | Electron Spin Resonance: Elementary theory and practical applications[END_REF]. As Ti 3+ ESR signal can be detected at room temperature, several simple geometries can be excluded: tetrahedron, octahedron and tetragonally elongated octahedron. A compressed tetrahedral symmetry can also be excluded because it would imply g ⊥ < g // . An elongated tetrahedral symmetry would impose (g // -g e )/(g ⊥ -g e ) = 4 (g e = 2.0023), which could only be verified by the site 2, although a four-fold coordination is quite unlikely for a cation as big as Ti 3+ (r VI (Ti 3+ ) = 0.67 Å> r VI (Ti 4+ ) = 0.605 Å [6]). However, ESR results can be fully explained by considering a compressed octahedral symmetry, highly distorted to have the order g // < g ⊥ (more distorted for site 2). This geometry could correspond either to a C 4v or a D 4h symmetry (5 to 6 coordination number). It is interesting to notice that the C 4v symmetry is the one of the square pyramidal O=TiO 4 frequently proposed for Ti 4+ environment in glasses [7].

EXAFS characterisation of Zr 4+ environment in parent glass

The short range order around Zr in the 6 wt. % Nd 2 O 3 parent glass was determined from the simulation of the EXAFS spectrum recorded at room temperature (Fig. 2). Zr occupies a well defined site (debye-waller factor σ 2 = 0.007 ± 0.001 Å -2 of 6.5 ± 0.7 coordination number with a consistent mean Zr-O distance d(Zr-O) = 2.15 ± 0.02 Å. This result shows that Zr 4+ ion is able to impose its own environment in the glass. This environment is relatively close to the one which was reported in literature for a depolymerised glass such as R7T7 glass, with a coordination number around 6 and a mean Zr-O distance ranging from 2.07 to 2.10 Å [8]. The fact that slightly higher coordination number and Zr-O distance were found in our glass could be due to the higher polymerisation of the glassy network.

For 100 moles of undoped glass, assuming that Al, Ti and Zr adopt respectively the geometries (AlO 4 ) -, (O=TiO 4 ) 2-and (ZrO 6 ) 2-, 24.7 moles of (CaO + Na 2 O) would be mobilised for their charge compensation, so that only 1.7 moles of (CaO + Na 2 O) would indeed be free to depolymerise the glassy network. In such a case, oxygens are less polarisable and a higher coordination number is required for Zr 4+ charge shielding. In zirconolite, Zr coordination number is 7 and d(Zr-O) = 2.20(1) Å. This short range order similarity with the parent glass could predispose to zirconolite nucleation. The EXAFS analysis of the next nearest neighbours around Zr is complicated (Fig. 2): at least two shells (likely three) with overlapping shapes are necessary to simulate the spectrum, and they partly interfere destructively. Nevertheless, the Fourier peak around 3.2 Å seems to be associated to Ca/Ti on the one hand and Zr on the other hand, both occurring at distances around 3.5 Å from Zr.

Spectroscopic investigation of Nd 3+ environment in parent glass

The fluorescence spectra corresponding to the 4 F 3/2 → 4 I 9/2 Nd transition were recorded at 15 K by exciting the Nd-doped glasses at the maximum of the 4 I 9/2 → 4 F 5/2 -2 H 9/2 Nd absorption band around 806.5 nm. Although the five expected emission lines cannot be resolved due to the inhomogeneous line broadening, these spectra allow to estimate the total splitting of the 4 I 9/2 ground multiplet which is about 500 cm -1 .

To compare, it was evaluated at 775 cm -1 for Nddoped zirconolite samples prepared by solid state reaction at 1460°C. The 4 I 9/2 splitting is a measure of the mean crystal field experienced by Nd 3+ ion in a host. The limit between compounds with low and high crystal field strength was set at 470 cm -1 by Auzel [9]. The greater crystal field of Nd in zirconolite as compared to parent glass indicates not only that the environments are different between the two hosts, but also that the mean Nd-O distance is shorter in zirconolite. Nd optical absorption spectra were recorded for the Nd-doped glasses at 15 K. They are all very similar, even though a slight broadening of absorption lines occurs with increasing Nd 2 O 3 concentration from 0.5 to 10 wt. % (Fig. 3). Three absorption bands from 4 I 9/2 state are particularly interesting: towards 2 P 1/2 (∼ 23200 cm -1 ), 4 F 3/2 (∼ 11400 cm -1 ) and 2 G 7/2 -4 G 5/2 (∼ 17200 cm -1 ). The first two bands ( 2 P 1/2 and 4 F 3/2 ) notably give information about the variety of Nd environment because of the low degeneracy of the excited state, whereas the latter band ( 2 G 7/2 -4 G 5/2 ) is described as hypersensitive and is very sensitive to Nd environment. For instance, the FWHM of the 4 I 9/2 → 2 P 1/2 band is approximately 150 cm -1 in Nd-doped parent glasses, much more than the one measured for zirconolite (30 cm -1 ), which reflects a higher distribution of environment in the glasses. Moreover, by comparison with literature [10], the shape of the 4 I 9/2 → 2 G 7/2 -4 G 5/2 is structure less (Fig. 3) and is characteristic either of highly polymerised glassy network or of glasses containing high field strength network modifiers. All these observations agree with an irregular Nd environment partly imposed by the glassy network topology, contrarily to Zr (see above). The environment of Nd was probed by EXAFS at L 3 edge (6208 eV). The short range order around Nd in the 6 wt. % Nd 2 O 3 glass was determined from the simulation of the EXAFS spectrum recorded at room temperature (Fig. 4).

Nd occupies a highly distorted site, the EXAFS simulation implies a cumulant analysis of the data (non-gaussian asymmetric Nd-O pair distribution [11]) with a large debye-waller factor σ 2 = 0.029 ± 0.001 Å -2 . Nd-O coordination number is 8.6 ± 0.8 and the mean Nd-O distance is about 2.53 ± 0.02 Å in the glass. In zirconolite, Nd coordination by O is 8 and, by EXAFS, Nd-O bond distance was found to be 2.44 ± 0.01 Å with a debye-waller factor σ 2 = 0.010 ± 0.001 Å -2 . So, the environment around neodymium in the glass shows striking differences with the one in the zirconolite crystalline phase and is much more distorted. Therefore, the incorporation of neodymium into zirconolite crystals during devitrification will imply strong reconstructive phenomena which are assumed to occur at high temperature, when the viscosity of the supercooled liquid is significantly lowered. In these conditions, it can be inferred that neodymium did not participate actively to the earliest steps of zirconolite nucleation processes. On the contrary, neodymium increases diffusionnal problems and can hinder zirconolite nucleation in the parent glass.

CONCLUSION

The environment around three cations, Ti 3+ , Zr 4+ and Nd 3+ was characterised by spectroscopic techniques in aluminosilicate glasses designed for the preparation of zirconolite-based glassceramic waste forms.

ESR study of Ti 3+ ions formed during glass melting at 1550°C revealed that their environment could be described by a C 4v or D 4h symmetry. This could correspond to a square pyramidal O=TiO 4 or to a six-fold coordinated octahedron axially compressed. The coordination numbers involved in such geometries are close to the ones encountered in zirconolite (6 and 5) which is aimed to crystallise by heat treatment of the glasses. Zr-K edge EXAFS spectra showed that Zr 4+ ions occupied well defined 6-7 coordinated sites, close to the ones of zirconolite. This result confirms the nucleating effect of Zr which was observed on zirconolite crystallisation in the glass. Contrarily to Zr 4+ ion which can be described as a cross-linking agent that reinforces the glassy network, Nd 3+ ion must rather be considered as a network modifier or a charge compensator which exhibits a lower ability to impose its own environment in a relatively polymerised glassy network. All the optical spectroscopies and EXAFS agree to describe Nd environment as being constrained by the glassy network with coordination number and Nd-O distance significantly greater than the ones of zirconolite. Therefore, Nd 3+ ion cannot act as a nucleating agent for zirconolite crystallisation.

Fig. 1 .

 1 Fig. 1. Experimental and simulated Ti 3+ ESR signal for the undoped glass.

Fig. 2 .

 2 Fit (solid lines) to the EXAFS (open circles) of the 6 wt. % Nd 2 O 3 parent glass (Zr-K edge, room temperature).

3 Fig. 3 .

 33 Fig. 3. Portion of Nd optical absorption spectra measured at 15 K in two parent glasses ( 4 I 9/2 → 2 G 7/2 -4 G 5/2 transition).

  )) (a. u.) R-α (Å) Fig. 4. Fit (solid lines) to the EXAFS (open circles) of the 6 wt. % Nd 2 O 3 parent glass (Nd-L 3 edge, room temperature).

TABLE I

 I

		. Results of the ESR simulation.
		g ⊥	FWHM	g ⁄⁄	FWHM
	Site 1	1.946 6.5 mT	1.88	12 mT
	Site 2	1.973	4 mT	1.89	4 mT