4uef/o3?f UM-P-88/124
Interdependence of Phase Chemistry, Microstructure and Oxygen Fugacity in Titanate Nuclear Waste Ceramics
by
W.J. Bukyx*. DM. Lcvinst, R.St.C. Smart", K.L. Smith*, G.T. Stevens*,
K.G. Watson*, D. Weedon§, and T.J. White* # x
* Advanced Materials Progam, Australian Nuclear Science and Technology] Organization, Private Mail Bag No. 1, Menai, New South Wales, 2234, Australia.
"School of Chemical Technology, South Australian Institute of Technology, The Levels, P.O. Box 1, Ingle Farm, South Australia, 5098, Australia.
§Julius Kruttschnitt Mineral Research Centre, University of Queensland St. Lucia, Queensland, 4067, Australia.
"National Advanced Materials Analytical Centre, School of Physics, The University of Melbourne, Parkville, Victoria, 3052, Australia.
Supported by the Australian National Energy Research, Development and Demonstration Programme.
* Author to whom correspondence should be addressed. * Member, The American Ceramic Society. 2
Abstract
Titanate ceramic waste forms were prepared using several combinations of calcination atmosphere (N2, ^-3.5% FU, K2) and metallic buffer (Ni, Fe, Ti, Al) to examine the dependence of microstructure and durabilitv upon oxygen activity. It was found that the microstructurcs and phase assemblages were mostly insensitive to the fabrication method, although in detail some systematic changes were recognized. The correlation between aqueous durability and oxygen fugacity was not straightforward due to density variations in the hot-pressed ceramics. These fluctuations in density dominated the dissolution characteristics of the waste forms and sometimes obscured the more subtle changes associated with redox potential.
It is concluded that although the best durability is achieved at lower fugacities (i.e. Ti metal buffer and ^ calcination atmosphere), a satisfactory
product can be produced using any of the preparative routes examined
providing the material is near theoretical density. 3
I. Introduction
The regulation of oxygen activity during the fabrication of high level
nuclear waste (HLW) ceramics has been the subject of much discussion.1.2.3,4
It is desirable to maintain reducing conditions for two reasons. First, the
mobility of volatile elements during processing is lessened at low oxygen
fugacity, thereby minimizing the concentration of active constitutents in
secondary waste streams. Second, low oxygen fugacity prevents the
formation of higher valence species (e.g. U ,Mo ) which stabilize water
soluble phases. To achieve these perceived advantages it is important that
,oxygen activity be regulated throughout, the fabrication process. During • |
calcination, evaporation of the HLW nitrate solution tafres place via the
evolution of oxygen and/or nitrogen oxides, which are removed using neutral (N2 or Ar) or reducing (N2/3.5% F^) carrier gases. However, if
thermal rise time.; are excessive, oxygen excursions will result in the
stabilization of higher valence species of the intermctallics and actinides.
During consolidation of the calcine, oxygen activity is usually buffered by
dispersing finely divided metallics throughout the powder, and/or by hot-
pressing in a self buffering cannister (e.g. a nickel tube or graphite die).
An earlier study by Ryerson, explored in detail the relationship
between oxygen activity and phase equilibria of a titanosilicate ceramic
immobilizing Savannah River Plant high level defense waste simulant.5 The
present article, describes variation? in the microstructure and phase
chemistry of aluminotilanate ceramics containing commercial high level
waste when prepared using different calcination atmospheres and solid state
buffers. 4
II. Experimental Methods
(i) Processing
Precursor powder having the composition given in Table I was prepared using the Sandia technique developed by Dosch and Lynch.6 A 10 wt% loading of HLW simulant (Table II) was slurried with the precursor, dried and then calcined at 750° C under a variety of atmospheres for one hour to voiatilize nitrates. During calcination an oxygen probe was used to monitor nitrate decomposition. Metallic buffers (2 wt%) were mixed with the calcine powders prior to • hot-pressing in graphite dies. Details - of the » calcination atmospheres, solid state buffers and hot pressing conditions are summarized in Table III.
(ii) Analytical Techniques
The phase assemblages of the calcines and consolidated wasteforms were characterized by X-ray diffraction (XRD) and selected area electron electron diffraction (SAD). Large scale inhomogeneities (> 10 |xm) were detected by backscattered electron imaging (BEI) of polished sections. Analytical electron microscopy (AEM) was used to probe individual grains and determine their chemistry. The surface chemistry of calcine powders and fracture surfaces were examined by X-ray photoelectron spectroscopy (XPS) and secondary ion mass spectrometry (SIMS) respectively. Aqueous durability of the ceramics was established using MCC-1 and hydrothermal testing. Details of our laboratory procedures are given elsewhere.7 5
III. Description of Calcine Powders
Pyrochlore is the predominant crystalline phase in all calcine powders, regardless of the atmosphere used to flush away the products of nitrate decomposition. However, except where conditions were most strongly
reducing (viz. pure H2), pyrochlore coexisted with smaller amounts of rutile and anatase (Fig. 1). The calcines consisted of small spherical panicles, 50-
100 nm in diameter which clumped together as partially sintered masses of variable size (Fig. 2). EDS and XPS analysis of the clumps gave compositions
for the major elements in reasonable agreement (± 5%) with their expected proportions (Table I); the similarity of bulk composition and surface (analysis suggests there is little zonation of matrix species with the calcine particles. The color of the powders progressively altered from cream (100% N2) to gray
3+ (N2/3.5% H2) to dark blue (100% H2) as greater quantities of Ti ions, which behave as color centres, were stabilized.
XPS revealed increasing surface enrichment of cesium and molybdenum as calcination conditions became less reducing (Table IV).
Moreover, these elements are weakly surface sorbed as their levels were significantly depleted after soaking in doubly distilled deionized water
(DDDW) for 5 minutes at room temperature. This effect was most dramatic for
the powder prepared under a neutral atmosphere (100% N2) where 79% of the cesium and 18% of the molybdenum were removed by this treatment (Fig.
3). These XPS data are in qualitative agreement with chemical analyses of
leach liquors (Tabic V). It will be noted that the losses determined by XPS are systematically lower than atomic absorption specfoscopy/inductively coupled plasma (AAS/ICP) analyses, probably as a result of redeposition of cesium and molybdenum during drying of the calcine; th's residue would contribute to the XPS signals from the washed powder. 6
IV. Characterization of Waste Forms
(i) Phase Assemblage
A summary of the phase assemblages as determined by XRD and SAD are given in Table III. Hollandite. perovskitc and zirconolite were common to all
preparations. Those ceramics fabricated without a metallic buffer, or above
the Ti-Ti02 buffer (viz. with Ni or Fe particles) contained significant
quantities of rutile, whilst both the Al and Ti metal buffered waste forms
contained reduced rutile (Magneli phases) (Figs. 4 and 5). It should be noted
that although reflections in the XRD traces are assigned to Magneli phases,
their diffracted intensities are variable and not in agreement with published
values. The reasons for this are not yet clear, however, their presence has
been confirmed by SAD. Further, the peak shapes of the predominant phases
differ between preparations end overlap is common. Consequently, a
quantitative interpretation was not possible. Minor phases were identified
by SAD and include hibonite, loveringitc, alumina and intermetallic alloys.
(ii) Homogeneity
Backscattered electron images show that all waste forms were
homogeneous with minor phase segregation. As noted previously, the
detailed analysis of these images is difficult as the volume of the
backscattered electron source exceeds the grain size." However, some
systematic differences are noteworthy. When a Ti metal buffer was used a
microstructure typified by Fig. 6(a) resulted - its dominant features are small
(<_ 1 urn) white areas indicative of intermetallic alloys and large (~ 4 tim),
irregularly shaped black areas which are usually Magndli phases. Those
preparations in which the metallic buffers were excluded from the
fabrication process contained acicular (3-5 urn long) black crystallites of stoichiometric TiO^ (Fig. 6(b)). These occur because limited reduction of tetravalent titanium occurs (even with pure PU z* the calcination atmosphere). Those waste forms prepared with Fe and Ni oxygen getters also contain numerous acicular rutile crystallites, as these metals cannot reduce oxygen activity below the Ti-TiCK buffer (Fig. 7). However, it should be noted that the overriding redox control for samples containg Ni and Fe particles was provided by the C-CO-CO2 buffer established by the graphite hot-pressing die. The result of this is illustrated in Fig. 8 which compares the BSE images of Fe and Ti metal remnants. In both instances some reaction has occurred with the surrounding matrix. However, the white contrast (arising from efficient backscattering due to high average atomic number) of the Fe-relic shows that little, if any, reduction has occurred. This demonstrates that its buffering capacity was not utilized due to the dominant effects of the graphite die. The Ti relic, on the other hand, has been oxidized and shows
dark contrast (due to a relatively low average atomic number). Darker bands
within the latter relic (visible in the original BSE image) reflect the slightly different backscatter co-efficients of twin domains within the reduced rutile;
such domains have been recognized in an earlier high resolution
transmission electron microscope study of titanate waste forms.9 Smaller
black areas within the mass of reduced rutile are shrinkage pores which arise as Ti is converted to TiC>2. The Al-buffered wasteform also contains
numerous black, acicular crystals (Fig. 9), but these are not stoichiometric TiC>2 (since alumina buffers below the Ti-TiC>2 buffer). Instead, they reflect
an unusual morphology exhibited by aluminium-rich Magneli phases. SAD established that the calcium aluminium titanate (CaTu AloO «n) described by
Morgan and Koutsoutis,10 has not been stabilized. 8
(iii) Nanostructure and Crystal Chemistry
Ail the waste forms were composed of submicrometer oxide grains and smaller (-0 1 u,m) alloy particles; the characteristic morphologies and distinguishing featu'es of these phases are detailed elsewhere.7 The main difference between the nanostructures of the materials prepared under highly reducing conditions (e.g. using Al or Ti buffers) compared with those of less reduced material (using a Ni buffer) was the preponderance of metallic precipitates in the former (Fij». 10).
A summary of chemical analyses for the most abundant constituents of the phase assemblage are given in Table VI. As the level of most fission products was low it was not possible to derive partitioning data quantitatively, although the results are in agreement with earlier studies.^*7-11 In brief, rare earth elements were found to partition strongly into perovskite, and to a lesser extent, zirconolite. For these phases, no correlation with the applied redox was disccmable. Because the hollandite-type phase can accommodate trivalent metals (viz. Fe , Ti , Al3+)12,13 it was expected, and found, that its stoicr iometry changed systematically with oxygen fugacity. In general, hollandite became more aluminous as conditions during fabrication became less reducing. Some hollandite grains in the Ni-NiO buffered material contained anomalously high concentrations of iron. This contaminant may have been introduced during milling of the calcine.
The stoichiometry of (reduced) rutile phases could be used as an oxygen probe, since the extent of reduction reflected the redox potential prevailing during hot-pressing. Analyses obtained from rutile in the Ti-TiCK and Al-
A^Oj buffered waste forms can be rationalized by considering them as the compound Ti^Oo. This Magneli phase contains considerable Al in place of 9
Ti3+, and divalent ions (e.g. Ca ) which enter the phase via the coupled -X substitution Ca^j T^ Ti^j Tipj Ti^j Ti^j
il 10
(iv) Durability
The correlation of applied redox potential with durability was complicated due to fluctuations in density which mask the more subtle changes in dissolution characteristics.14 For the sample set in which the calcination atmosphere was varied and no metal buffer added, the density remained constant at 99% of the maximum density (MD)* , thereby allowing direct interpretation of the data. These data are summarized in Fig. 11(a), and shows that for the three elements chosen (cesium, strontium, barium), a small ibut systematic reduction in elemental, losses accompanied the imposition of successively more reducing conditions.
For the series prepared using a titanium metal buffer (Fig. 11(b)), there is an order of magnitude improvement in durability as oxygen fugacity is minimized. Significantly, in the case of the ^73.5% H2 calcined sample
(specimen no. 355) where the density falls to 97% MD, cesium losses increase whilst those of the non-volatile elements (strontium and barium) are at expected levels. This is due to entrapment of cesium vapour in micropores during hot-pressing in a form which is easily released during dissolution tests.15 On the basis of the above criteria, it is not unexpected that the most durable waste form (specimen no. 350) fortuitously yielded the maximum density observed, and was prepared under conditions of mininum oxygen fugacity (H2 calcination atmosphere and Ti metal buffer).
In the waste form series fabricated using different metal buffers the durability deteriorated as the oxygen potential decreased (Fig. 11(c)), in
We have chosen to use the term maximum density (MD) rather than theoretical density (TD) since uncertainty regarding the number of phases and their proportions make calculation of TD unduly speculative. Instead, the highest density achieved in this series of experiments (4.39 g.cm ) was choosen, on the basis of scanning electron microscopy, to represent the fully dense ceramic. 11 contradication to the above result*. However, in this instance the oxygen activity and density decreased in unison. Therefore, the poorer durability of these samples correlates directly with less complete compaction and nullifies the benefits of lower oxygen activity. The iron-buffered waste form was the least resistant wasteform, largely as a result of the significant corrosion of Fe relic particles; see the mass losses giv^n in Table III. As the metals are bound to the oxide mainx by soluble, nuclide bearing glassy phases, dislodgement of iron particles exposes larger areas of the ceramic to aqueous attack (Fig. 12).
Ion beam thinned specimens were studied by TEM using techniques described previously.16 Our results are identical to those obtained in other
studies,17-18 and are briefly described here as a matter of course. Micrographs of specimen nos. 409 and 410 treated for 3 days at 150°C are shown in Fig. 13. Perovskite has been subjected to the most complete dissolution, as evidenced by the precipitation of TiC^ brookite crystals.
Hollandite was slightly dissolved by this treatment, whilst zirconolite is highly resistant. In the case of the Fe buffered sample, the brookite crystals do not have a well defined morphology, due perhaps to the inclusion of iron which inhibits grain growth (Fig. 14). 12
V. Discussion
(i) Crystal Chemistry and Durability
The present work generally confirms and extends earlier investigations which showed that chemical durability of aluminotitanate waste forms increases with decreasing oxygen activity during fabrication.1 «4»5,19
Durability is maximized when every waste species resides in a limited number of crystallochemical environments. This is achieved by controlling oxygen fugacity during two k"y events in the fabrication process - calcination, and to a? lesser extent, hot-pressing. As pyrochlore bis the crystalline progenitor } | » j If if of all other phases,20 it is important to stabilize as much of this phase as possible during calcination. To achieve this, it is necessary to operate under neutral or reducing atmospheres. Calcination in strongly reducing H2 results in almost a 100% pyrochlore yield, while in the neutral atmosphere of N 2 a portion of the TiO^ (as anatase and rutile) remains unreacted. The quantity of these latter phases should be minimized, as they cannot incorporate waste species into their structure, but rather, provide a substrate onto which volatiles such as cesium and molybdenum absorb. It is believed that sorbed species become trapped as grain boundary phases and in micropores during consolidation. This is undesirable, as these structural features are known to be susceptible to hydrothermal attack.15 However, we cannot present direct evidence to substantiate the formation of chemically different grain boundaries and pore surfaces in waste forms calcined under different oxygen potentials. Static SIMS depth profiles obtained from intergranular fracture faces of the fully dense waste form prepared in h^, gave cesium peaks which where 2-3 times greater in amplitude, but considerably narrower, than that for ceramics prepared under less reducing conditions (e.g. N2 calcination atmosphere). A further difference, was the 13 unusually wide titanium, calcium, aluminium and strontium peaks in the strongly reduced waste form compared with N2 calcined material. Although
SIMS does indicate variation in grain boundary composition and the concentration of glassy films per unit volume, we have been unable from this evidence to quantify the changes.
Calcination is considered successful if sufficient titanium is reduced to the trivalent state to ensure that the wide variety of altervalent subsititutions, necessary for the incorporation of waste elements in the major phases can operate, thereby reducing the quantity of metastable compounds in. the hot-pressed product.19 As a safeguard against an unsuccessful calcination, a metallic buffer is added prior to hot-pressing to getter oxygen, and in the case of Ti- and Al-buffers, produce additional trivalent titanium. On the laboratory scale, metal buffers need not be used if hot-pressing is carried out in graphite as the die itself results in a moderately
reducing environment close to the Fe-FeO buffer. For large scale production, where hot-pressing may take place in stainless steel bellows or cannisters, the inclusion of a metal oxygen getter is essential.21 If calcination fails catastrophically, then the use of Ni (and possibly Fe) buffers should be
avoided, as these metals will not getter oxygen effectively enough to restore the desired phase assemblage and soluble compounds such as CS2M0O4 will be
stabilized.'-^
No substantial difference in the partitioning of rare earth elements
(REE) could be detected with changes in redox as inhomogeneity on a grain- to-grain basis masked systematic changes. Sampling a statistically significant number of grains could have overcome this problem,5 but this was beyond the scope of the present study. Importantly no new REE-rich phases were detected indicating that all waste forms had sufficient capacity to immobilize these fission products. It should be noted however, that the 14
REEs may not accurately simulate the behavior of transuranic elements, where the oxidation st-tes and partitioning would vary over the range of these oxygen fugacity experiments.22 Because the hollandite-type can incorporate a number of trivalent m~tal ions in its structure, the composition cf this phase varied from titanium-rich (in the Al and Ti buffered material) to aluminium-rich 'in Ni buffered material). It has been shown that the former stoichiometries permit higher concentrations of cesium to enter hollandite,12*2^ but at a 10 wt% waste loading this additional capacity was not required.
(ii) Density I |
Oversby^ reported that Sethi et al.25 prepared a series of titanate waste forms by sintering under controlled CO/CCU atmospheres. These workers found that not only was chemical durability redox dependent, but also the density, with the percentage TD improving from 89 to 99% in passing from pure CO^ to CO. This contrasts with this study where density, although somewhat variable, did not correlate with oxygen activity. Therefore, hot- pressing is superior to reactive sintering, not only because it will generally lead to higher TDs, but also because it minimizes product deterioration should the redox control unexpectedly fail.
A more important parameter controlling waste form quality during hot- pressr._, is homogeneity of the calcine and blending of the metal buffer. For example, in the series of Ti-buffered waste forms calcined using different atmospheres (Fig.11(b)), the density of the ^-3.5% H2 material (specimen no. 3SS) was anomalously low (97% TD) leading to greater cesium dissolution.
Optical micrographs (Fig. IS) clearly show this ceramic to be heterogeneous when compared with the more completely densified material calcined in rU or N~. 15
(iii) Solid State Buffers
Density fluctuations disguised correlations between durability and oxygen fugacity in the waste forms prepared with different metallic buffers
(Fig. 11(c)). Furthermore, the study of iron and nickel buffers was effectively negated by hot-pressing in a graphite die which itself established the oxygen activity during hot-pressing. However, it is clear that with the exception of iron, all the metal buffers should yield a satisfactory product if near theoretical density is achieved. Iron is not a recommended buffer for, if
it remains in its reduced state, rapid corrosion will expose highly soluble, cesium-bearing glassy films. f I
The absence of substantial aluminium partitioning into the waste- bearing phases of the Al-A^O^ buffered material, even though the Al was
completely oxidized, suggests that the crystal chemistry of the waste phases is
predetermined at calcination (i.e. when pyrochlore is formed). It appears
that little solid state diffusion of the buffer species occurs. This is
presumably the case for the other buffers also, though we were unable to
determine this for the waste forms containing (unoxidized) iron and nickel
metal. The lack of solid state diffusion suggests that none of the buffers
. xamined should deliteriously modify the chemistry of the waste form
assemblage. 16
VI. Conclusions
Providing slightly reducing conditions ar; employed during fabrication
(sufficient to prevent the formation of uranyl and molybdate compounds) a waste form with reproducible chemistry and durability can be prepared.
More important than redox control is the attainment of high density, which will override small fluctuations in oxygen activity. As the density achieved by hot-pressing is strongly dependent upon homogeneity, it is the precalcining and calcination conditions, rather than the metal buffer, which is of crucial importance. Therefore, compact density would be a key quality control device ink commercial production. it I! Acknowledgements: Argon-ion-beam-thinned sections were skillfully prepared by Roy Warren (ANSTO). Optical microscopy was carried out by
Alan Bellrose (ANSTO). 17
References
'A.E. Ringwood, Mineral Waste Form Development for U.S. Defense Wastes,"
Report No. UCRL-15347, Lawrence Livermore National Laboratory, 1981.
2C.M. Jantzen, J. Flintoff, P.E.D. Morgan, A.B. Harker, and D.R. Clarke, "Ceramic
Nuclear Waste Forms," Ber. Kernforschungsanlage Juelich [Conf.] Juel-Conf-
42 (Vol.2), Proc. Inst. Semin. Chem. Process Eng. High-Level Liquid Waste
Solidification, 693-706 (1981).
3V.M. Oversby, "Properties of Synroc C Nuclear Wast. Form: A State-of-the-
Art Review," Report t No. UCRL-53326, Lawrence Livcrmore National 1 » I Laboratory, 1982.
4R.G. Dosch, T.J. Headley, and P.Hiava, "Crystalline Titanate Ceramic Nuclear
Waste Forms: Processing and Microstructure," J. Am. Ceram. Soc., 6_7_ [5] 354-
361 (1984).
5F.J. Ryerson, "Phase Equilibria of Nuclear Waste Ceramics: the Effect of
Oxygen Fugacity," J. Am. Ceram. Soc., 6_Z [2] 75-82 (1984).
6R.G. Posch and A.W. Lynch, "Solution Chemistry in Synroc Preparation,"
Report No. SAND 80-2375, Sandia National Laboratory, Albuquerque, NM, 1980.
7W.J. Buykx, K. Hawkins, D.M. Levins, H. Mitamura, R.St.C. Smart, G.T. Stevens,
K.G. Watscii, D.V/^edon, and T.J. White, "Titanate Ceramics for the
Immobilization of Sodium-Bearing High-Level Nuclear Waste," J. Am. Ceram.
Soc, 21 [8] 678r688 (1988). 18
8P.E. Fielding and T.J. White, "Crystal Chemical Incorporation of High-Level
Waste Species in Aluminotitanate-Based Ceramics: Valence, Location,
Radiation Damage and Hydrcthermal Durability," J. Mater. Res., 2_ [3] 387-414
(1987).
9C.J. Ball, W.J. Buykx, F.J. Dickson, K. Hawkins, D.M. Levins, R.St.C. Smart, K.L.
Smith, G.T. Stevens, K.G. Watson, D. Weedon, and T.J. White, "Titanate Ceramics for the Stabilization of Partially Preprocessed Nuclear Fuel Elements," J. Am.
Ceram. Soc, in press (1988)
^P.E.D. Morgan and M.S. Koutsoutis, "Phkse Relations in the Ca-Ti-Al-O if XL System; Further Studies on Members of the Cm Family," J. Mat. Sci. Lett., 4,
321-323 (1985).
"T. Murakami, "Distribution of Radioactive Waste Elements Among
Constitutcnt Minerals of a Cermaic Waste Form," Mineral J., 11 [7] 344-351
(1983).
12 Svnr S.E. Kesson and T.J. White, "[BaxCsy] [(Ti,Al)2*+y Tig*2x-y] °16 oc-type
Hollandses I Phase Chemistry," Proc. R. Soc. Lond. A4Q5_, 73-101 (1986).
13R.W. Cheary, "An Analysis of the Structural Characteristics of Hollandite
Compounds, "Acta Crystallogr., B42. 229-236 (1986).
14J.L. Woolfrey, D.M. Levins, R.St.C. Smart, and M. Stephenson, "Effect of Hot-
Pressing Conditions on the Structure and Leachability of Synroc," Am. Ceram.
Soc. Bull., 6j6_ [12] 1830-1846 (1987). 19
^J.A. Cooper, D.R. Cousens, J.A. Hanna, R.A. Lewis, S. Myhra, R.L. Scgall,
R.St.C. Smart. P.S. Turner, and T.J. White, "Intergranular Films and Pore
Surfaces in Synroc C: Structure, Composition and Dissolution Characteristics,"
J. Am. Ceram. Soc., 6j> W 347-352 (1986).
16D.R. Clarke and J.F. Flintoff, "Preferential Dissolution Phenomena of
Nuclear Waste Materials," pp.29-36 in Scientific Basis for Nuclear Waste
Management VI. Edited by D.G. Brookins, Elsevier, New York, 1983.
17T. Kastrissios, M. Stephenson, P.S. Turner, and T.J. White, "Hydrothermal
Dissolution of Perovskite: Implications for Synroc Formulation," J. Am. if I Ceram. ^Soc., ZQ [7] C-144 - C-146 (1987). ' "
18S. Myhra, P. Delogu, R. Giorgi, and J.C. Rivi6re, "Scanning and High-
Resolution Auger Analysis of Zirconolitc/Perovskite Surfaces Following
Hydrothermal Treatment," J. Mat. Sci., 22, 1514-1520 (1988).
19S.E. Kesson, "The Immobilization of Cesium in Synroc Hollandite," Radioact.
Waste Manage. Nucl. Fuel Cycle, 4 [1] 53-72 (1983).
20F.J. Dickson, H. Mitamura, and T.J. White, "Radiophase Development in Hot-
Pressed Alkoxide Derived Titanate Ceramics for Nuclear Waste Stabilization," J.
Am. Ceram. Soc, 21, in press (1989).
21A.b. Ringwood, A. Major, E.J. Ramm, and J. Padgett, "Uniaxial Hot-Pressing in Bellows Containers," Nucl. Chem. Waste Manage., 4, 135-140 (1983).
22Hj. Matzke, B.W. Seatonberry, I.L.F. Ray, H. Thielc, H. Trisoglio, C.T. Walker, and T.J. White, "The Incorporation of Transuranic Elements in Titanate
Nuclear Waste Ceramics," unpublished report. 20
23R.W. Cheary and J. Kwaitkowska, "An X-ray Structural Analysis of Cesium
Substitution in the Barium Hollandite Phase of Synroc," J. Nucl. Mat., 125, 236-
243 (1984).
24A.E. Ringwood, "Disposal of High-Level Nuclear Wastes: A Geological
Perspective," Mineral. Mag., 49_ 159-176 (1985).
25V.K. Sethi, R.M. Arons, J.P. Guha, J.T. Dusek, and R.B. Poeppel, "Fabrication,
Characterization and Evaluation of Synroc Wasteform," Program Review,
High-Level Waste Technology, Atlanta, GA (1981). :| ll 21
Table I. Matrix Element Compositior of Calcine as Determined by XPS
Oxide Precursor Element Expected Calcine Experimental Composition Composition Calcine (wt%) (at%; Composition (at%)
Ti02 71 2 Ti 24.95 20.7
GO 11.1 Ca 5.55 3.8
Zr02 6.8 Zr 1.54 1.9*
BaO 5.5 Ba 1.01 0.7
A1203 5.4 'Al 2.97 12.8
63.99 60.2
100.01 100.1
* includes waste Zr02. t Overlap of Al-2p with Cs-4d combined with low sensitivity for Al make quantification for this element difficult. 22
Table II. Simulated High Level Waste
Component wt% Component wt% Component v/i%
Mo0 13.12 8.26 3.72 3 c&p °Wh
7.54 BaO 3.93 U 5.16 Ru02 3°8
Rh 0 1.24 SrO 2.68 P 1.65 2 3 2°5
PdO 3.72 Zr0 12.50 Fe 3.82 2 2°3
Ag 0 0.21 Y 1.55 Cr203 0.83 2 2°3
cao Ce02 12.19 NiO 0.31
TeQ i| | 1.86 Nd 0 15.50 2 2 3 i 100.00 Table III. Conditions Used for the Preparation of Titanate Ceramics. All Specimens were Hot-Pressed in Graphite Dies (2h, 1200°C, 16 MPa)
Specimen Precursor Calcination Range of Metal Density Phase MCC-1 Test
3 -2 -1 No. Batch No. Atmosphere f02 at 750°C Buffer (gem" ) Assemblage Mass Loss (mg.m .d )
22 24 350 43 H2 io- - io- Ti 4.39 H, P, Z. M, A 7±3
351 43 H2 - - 4.3-w" *H, P, Z, R 12 ±7
1 5 353 43 N2 10" - 10" Ti 4.35 *H, P, Z, M 16 ±7
354 43 N2 - - 4.37 *H, P, Z, R 17 ±9
1 20 355 43 N2/3.5% H2 10" - 10" Ti 4.27 *H, P, Z, M 1±0
356 43 N2/3.5% H2 - - 4.37 *H, P, Z, R 9±2
20 409 43 co/co2 -IO- Fe 4.28 H, P, Z, R, A, L, Aa 563 ± 38
24 410 43 N2/3.5% H2 -IO- Ti 4.24 H, P, Z, M, A, Hi 4±2
16 411 43 co/co2 -IO- Ni 4.33 H, P, Z, R, A 6+6
422 43 vacuum - Al 4.17 H. P, Z, M, A, Aa 26 ± 6
H - hollandite, P - perovskite, Z - zirconolite, M - Magneli phases, R - rutile, Hi - hibonite, Aa - alumina, L - loveringite, A - alloy * Characterized by XRD only 24 Tabic IV. Composition of Calcines Before and After Washing in DDDW for 5 Minutes
Composition (at%)
3.5% H /N 100% N 100% H2 2 2 2
Element/Ratio Initial Washed Initial Washed Initial Washed
Ti 20.35 20.46 19.63 21.54 19.14 21.49
O 59.30 61.71 55.90 62.86 55.78 62.22
Ba 0.70 0.80 1.10 1.63 1.10 1.65 a 1.16 0.80 1.87 0.58 1.98 0.47 3.72 3.54 3.64 3.73 3.41 3.54 &u i Hi 1.86 1.83 1.76 II 2.21 1.87 2.01 Mo 0.35 0.34 0.44 0.35 0.55 0.35
Al 12.56 10.51 15.66 7.10 16.17 8.26
Cs/Ti (x 102) 5.7 3.9 9.6 2.7 10.3 2.2
Mo/Ti (x 102) 1.7 1.7 2.2 1.6 2.9 1.6 25
Table V. Percentage of Cs and Mo Lost from Calcine
100% H2 3.5% H2/N2 100% N2
Technique Cs Mo Cs Mo Cs M:>
XPS 32 0 72 14 79 18
AAS/ICP 48 17 100 38 100 40
I II 26
Table VI. Chemical Analyses
Phase Oxygen Stoichioiretry Buffer 3+ 4+ 4+ Zirconolite A1-A1203 [Ca 75 Nd Q2 Zr 231 L1 [Ti 6Q Ti „ Ti Q1] £1 [Tij ?6 Al Q9 Si Q9 MoQ4] u 9g 07
3+ 4+ 3+ 4+ Ti-Ti02 [Ca 7g Nd Q3 Zr 19] j., [Zr 69 Ti 31] u [Tij 6? Al Q7 M0()4 Si 2Q] £, 9g 0?
Fe-FeO 3+ 4+ 4+ 3 + Ni-NiO [Ca ?9 Nd Q2 Zr 19] ^ [Zr 6Q Ti 33 Ti 0?] £1 [Tij ?9 MoQ 4 SiQ g Al 0?] z{ 9g 0? 4+ .3+ .4 + Perovskite A1-A1203 [Ca 79 Nd 07 Ce01 SrQ j] z gg [Ti Qg Ti 9g] £1 Q6 03 14+ .3+ .4+ . Ti-Ti02 [Ca g6 Nd^ Ce01] r 93 [Ti 07 Ti %] ri Q3 03 Fe-FeO 1' ,, ,• 'I lCa.80 Nd.07 Sr.0l ^.Ol'l^O rri.10 T,.91 ^.Ol111.02 °3 .3+ .4 + Ni-NiO [Ca ?6 NdQ6 SrQ1 Ce01] z g4 [Ti07 Ti 99] L1 Q6 03 .3+ .4 + Hollandite A1-A1203 [Ba 36 Cs02] j. 3g l(Al 3? Ti 3?) (Zr 1Q Ti? 13>] I? 9? 0J6 Ti-Ti0 Ba Cs r 2 t .31 .04> 1.35 ^.35 \3^ ^.06 ^.M" 17.97 °16 .3+ , .4 + Fe-FeO [Ba 34 Cs 05]r 39 [(Al 49 Ti 23) (Zr Q7 ^7M^i,7.99 °16 + Ni-NiO [(Ba 40 CsQ1 Ca Q3] ZM [(Al 5fi Fe*/,) (Zr 1Q Ti* 12)] ip Q9 0,6 [(Ba.40 Cs.01)] 1.41 «A!.75 ^oV (Zr.01 T\l2\)] £8.02 °16 4+ _ .3+ 3+ .4 + Rutile- A1-A1203 [Ti2 92 ZrQ6] I2 9g [Ti lg Alj og Fe Q2 CaQ35 Sr Q] Ti 3(-] I2 Q{) 09 4+ 4+ 3+ 4 + Magneli Ti-Ti02 [Ti2 91 Zr 05] n96 [Ti 52 Al 92 Ca 2g Ti 2g] £2 09 Phases Fe-FeO lTi.98 ^ °2 1 Ni-NiO ^98 Zr.02l °2 27 Table VI. Chemical Analysis (Con'd) Phase Oxygen Stoichiometry Buffer Alloy A1-A1203 Mo ?4 Fe ,9 Pd Q5 Ni^ Rhu Ti-Ti0 Ti Mo A1 Fe 2 .50 .24 .04 .06 ^.05 Pd.04Mo.64Ti.17Fc.09Pd.06Rh.03 Fe-FeO Mo.46 Fe.25 Ti.13 ^.10 Pd.05 ^.01 Ti.33 Mo.25 Fe.24 Pd09 ^.09 Ni.01 Ni-NiO Ni Mo .87 .13 Ni.28! Mo—.3.311 Ti".1.122 ^.1.13 3Fc .0.06 6Pd "".0.08 8 Notejthat the titanium content in perovskite is overestimated due to convolution of Ce-L and Ti-K X-ray I fluorescent energies 28 Figure Captions Fig. 1. XRD traces of calcine powders prepared under H2, ^-3.5% H2 and N2- (Py - pyrochlore, R - rutile, A - anatase) Fig. 2. Secondary electron image of calcine powder prepared in pure Nj. Fig. 3. XPS traces of Ba and Cs signals for calcines before and after washing in DDDW for 5 minutes. Fig. 4. XRD traces for hot-pressed ceramics fabricated from calcine powders prepared under different atmospheres anid |witkwith and without a Ti 1 k metal buffer. (H - hollandite, P - perovskitevskite,. Z - zirconolite,. RR - 'I rutile, X - unknown) Fig. 5. XRD traces for ceramics hot-pressed with different metallic buffers. Fig. 6. BSE images of waste forms fabricated from H2 calcined powder, (a) with Ti-metal buffer (specimen no. 350), and (b) without metal buffer (specimen no. 351). Fig. 7. Correlation between oxygen fugacity and specimen number. Fig. 8. BSE images of metal relic particles, (a) Fe relic (specimen no. 409), and (b) Ti-relic (specimen no. 410). Fig. 9. BSE image of Al-metal buffered waste form (specimen no. 411). Fig. 10. Bright field electron micrographs of (a) Ni-NiO buffer (specimen no. 411), and (b) Ti-TiCK buffered (specimen no. 410) waste forms. 29 Fig. 11. Correlations between density, elemental dissolution and redox potential during fabrication. Dependance en (a) calcination atmosphere when no metal is added during hot-pressing, (b) calcination atmosphere with the subsequent addition of titanium metal buffer, and (c) metal buffer species. Fig. 12. Dissolution of alloy particle in oxide matrix. Fig. 13. Bright field image of dissolving oxide matrix in (a) Ti metal buffered waste form (specimen no. 410), and (b) Fe metal buffered waste form (specimen no. 409). i . U I) Fig. 14. EDS of brookite precipitates in (a) Ti metal buffered waste form (specimen no. 410), and (b) Fe metal buffered waste form (specimen no. 409). Fig. IS. Optical micrographs waste forms prepa/ed using different calcination atmospheres and Ti metal buffers. (a) N2 atmosphere (specimen no. 353). (b) N2/3.5% H2 atmosphere (specimen no. 355). (c) H2 atmosphere (specimen no. 350). i§i!^§§ij^ 30 *2 II It py py VW> »'*»»>A*#~*HJ^+*S' ,^>«K ..J - r~^~ —T 1 r T 1 1 r- N2-i5l H2 ^^»»«^>^M —\ 1 T 1 1 T~ ki 1 .1 r- ^A^^LW^JJJU VA^W. S ' 3 ' IT "15 ' 7T 2» Co K« Figure 1 31 I Figure 2 32 as received washed _i i i i t_ Bo3tJ Cs3d ll ^UL_ w ~i 1 1 1 r 1 1 1 1 r 'c _i i i i i _i i I I L. CO I'^UL N -35%H w 2 2 c -i r -i 1 1 i i 8 _i u o —I 1 1 1 r— 850 ' 750 ' 650 850 750 650 Binding Energy (eV) Figure 3 33 "? Mj 15% H, R H » r ' •: P P buffer 1 1 R z I H 1 Z I Z| 2JL Rz H M III H UL 41 I) *^t 3f7 33 2'9 Nj a5*H2 Tl i Duffer r~iT,"°?" •!**»*• Figure 4 34 ll Al mate) butter Fa maw buttar M nwui buffer b 6 AJL JUL uL UX- iLL JU4t 37 13 ?fl 41 37 33 » mi41 37 33 » Uu41 37 L33 " ;r9 I j T.„0?n., Mapa* (*NS 2« CO K(» Figure 5 35 Figure b RS«!S3SWSH>;-; 36 ll ^sr>^^4 351,354,356, .409,411 350,353,355,410 422 Figure 7 J?-~X*P 37 I Figure 8 38 Figure 9 39 ll 1.0 ntn Figure 10 KH."~:>:p 40 oc. • V OBo N3 ^-35*82 rtj Cdcinolnii AmottJwra Figure 11 41 Figure 12 42 Fi Kiire 1 i 43 Ti Fe Ti Mo L J^'^'.ltb KeV Figure 14 44 •: ''' -r.Vy '^ • *^-»^7?A - Figure '.5