Pergamon
0149-1970(94)0001&6
RADIATION EFFECTS IN NUCLEAR
Progressin NuclearEnergy,Vol. 29, No. 2, pp 63-121, 1995
Elsewr Science Ltd
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WASTE FORMS FOR
HIGH-LEVEL RADIOACTIVE WASTE
R. C. Ewing,* W. J. Webert and F. W. Clinard, Jrf
*University of New Mexico, Albuquerque, NM 87131, U.S.A.
tPacific Northwest Laboratory, P.O. Box 999, Richland, WA 99352, U.S.A
$Los Alamos National Laboratory, Los Alamos, NM 87545, U.S.A.
(Received 12 July 1994)
Abstract-High-level nuclear waste in the United States comprises large volumes (tens of millions of cubic meters),
high total activities (billions of Curies) and highly diverse and complex compositions. The principal sources of
nuclear waste are: (i) spent nuclear fuel from commercial and research nuclear reactors; (ii) liquid waste produced
during the reprocessing of commercial spent nuclear fuel; (iii) waste generated by the nuclear weapons and naval
propulsion programs. The latter category now includes over 100 metric tons of plutonium and many hundreds of
tons of highly enriched uranium from the dismantling of nuclear weapons.
Most of these wastes will require chemical treatment, processing and solidification into waste forms for
permanent disposal. The long-term effects of radiation on waste form solids is a critical concern in the performance
assessment of the long-term containment strategy. In the case of spent nuclear fuel, the radiation dose due to the inreactor
neutron irradiation is already substantial, and additional damage accumulation during disposal is not
anticipated to be significant; thus, this is not a subject addressed in this review paper. In contrast, the post-disposal
radiation damage to waste form glasses and crystalline ceramics is significant. The cumulative cc-decay doses which
are projected for nuclear waste glasses reach values of 1016 a-decays g- ’in 100 yr. Similarly, crystalline waste forms,
such as Synroc will reach values of > 10’s cc-decay events g - ’in 1000 yr for a 20 wt% waste loading. These doses are
well within the range for which important changes in the physical and chemical properties may occur, e.g. the
transition from the crystalline-to-aperiodic state in ceramics.
This paper provides a comprehensive review of radiation effects (due to y-, p- and a-decay events, as well as from
actinide doping experiments and particle irradiations) on nuclear waste form glasses and crystalline ceramics,
particularly Synroc phases, zircon, apatite, monazite and titanite. The paper also includes recommendations for
future research needs.
1. INTRODUCTION
High-level nuclear waste (HLW) in the United States
has three major sources: (i) spent nuclear fuel from
commercial and research nuclear reactors; (ii) liquid
waste produced at West Valley, New York, in the
1960s during the reprocessing of commercial spent
nuclear fuel; (iii) waste generated by the nuclear
weapons and naval propulsion programs. Much of the
defense waste is stored in tanks at Hanford and
Savannah River sites in Washington and South
Carolina, respectively. The volumes are substantial. At
the Hanford site alone the total inventory of low- and
high-level waste includes 1l,OOO,OOOm3 of fluids and
6900 metric tons of nuclear materials, which includes
4100 metric tons of uranium and 15 metric tons of Cs
and Sr in capsules. The total activity of the waste at the
Hanford site is estimated to be 446,000,OOO Curies
(Illman, 1993). For all of the principal DOE sites
(Hanford, Savannah River, Idaho Chemical Processing
Plant and the West Valley Demonstration Project)
the combined total volume of high-level waste through
1988) is 385,000 m3, and the total activity is
1,200,000,000 Curies (Office of Technology Assessment,
1991). Although the defense wastes account for
95% of the volume of the current total United States
HLW inventory, they account for only 9% of the total
radioactivity. The majority of the radioactivity is
contained in commercial spent nuclear fuel (Weber,
1991a). Except in the case of direct disposal of spent
fuel, which is the current U.S. policy, these wastes must
be treated and solidified prior to permanent disposal.
Finally, under the first and second Strategic Arms
Reduction Treaties (START I and II) and unilateral
pledges made by the United States and the former
Soviet Union and Russia, many thousands of nuclear
weapons will be dismantled during the next decade.
This will generate over 100 metric tons of plutonium
and many hundreds of tons of highly enriched uranium
that will require disposal (National Academy of
Sciences Report, 1994).
63
64 R. C. Ewing er al.
To the extent that these radioactive wastes must be
treated and solidified prior to final disposition, the effect
of radiation on waste form solids is a critical concern.
One of the most important applications of the study of
radiation effects in ceramics (vitreous and crystalline)
during the past decade has been the effect of radiation
on nuclear waste forms, e.g. borosilicate glasses, single
and polyphase crystalline ceramics, and the UO, in
spent nuclear fuel (Weber and Roberts, 1983; Matzke,
1988a,b; Ewing et al., 1988; Weber, 1991a). In the case
of spent fuel, the radiation dose due to the in-reactor
neutron irradiation is already substantial, and additional
damage accumulation during disposal is not
anticipated to be significant. In contrast, the postdisposal
radiation damage to waste form glasses and
crystalline ceramics is significant. The cumulative
m-decay doses per gram which are projected for the
nuclear waste glasses to be used for the West Valley
Demonstration Project (WVDP) and the Savannah
River Plant (SRP) reach values of 10” a-decay events
g -r in 1000 yr (Weber, 1991a). Similarly, crystalline
waste forms, e.g. the titanate assemblage, Synroc,
would reach values of > 10” a-decay events g-’ in
1000 yr for a 20 wt% waste loading. For this Synroc
composition, the total accumulated dose at the end of
10,000 yr is nearly equivalent to 0.5 displacements per
atom. This dose is well within the range for which
important changes in physical and chemical properties,
most notably the transition from the crystalline-toaperiodic
(amorphous or metamict) state can occur.
Although present performance assessments of
radionuclide containment rely primarily on the geologic
repository, the physical and chemical durability of
the waste form can contribute greatly to successful
isolation of nuclear waste (Ewing, 1992a). Thus, the
effect of radiation on physical properties and chemical
durability are of importance. The changes in chemical
and physical properties occur over relatively long
periods of storage, 104-lo6 yr, and at temperatures
which range from 100 to 450°C depending on waste
loading, age of the waste, depth of burial, and the
repository-specific geothermal gradient. Thus, one
challenge is to simulate effectively radiation damage at
relatively low fluxes over long periods of time. In this
regard, studies of naturally occurring U- and Thbearing
phases that are structurally and chemically
analogous to waste form phases have provided
important data and reference points for long-term
effects (Ewing et al., 1988).
The essential issues addressed and summarized in
this review are the following:
(1) What are the changes in physical properties (e.g.
volume, fracture toughness and hardness), chemical
properties (e.g. thermodynamic stability and
leach rate) and stored energy as a function of
radiation dose, temperature and time?
(2) How can the changes in physical and chemical
properties to relevant radiation doses be simulated
in accelerated experiments (for a waste form age of
104-lo6 yr)?
We begin with a brief history of the development of
nuclear waste forms and a summary of their types.
1.1. History
The early history of the development of nuclear
waste forms is summarized by Lutze (1988). These
efforts, mainly in the 1950s focused on the incorporation
of radioactive waste into glasses of widely varying
compositions. The first radioactive wastes were incorporated
into a nepheline syenite (a silica-deficient
igneous rock) glass at Chalk River, Canada. The
processing temperature of the nepheline syenite glass
(1350°C) was much higher than that of the now widely
adopted borosilicate glasses (110&l 150°C). Major
programs during the 1960s in the United States,
Canada, Germany, France, Italy, Japan, the United
Kingdom and the Soviet Union were directed at
developing glass compositions that had lower melting
points in order to reduce loss of radionuclides by
volatilization, incorporated waste without phase separation,
resulted in glasses which were durable, and for
which vitrification was a feasible large scale industrial
process. Borosilicate glass is now the waste glass
composition of choice for most nations. Industrial use
of the vitrification process has been in progress since
1978.
During the period from 1977 to 1982, there was a
tremendous diversity in the types of nuclear waste
forms which were under development (Lutze and
Ewing, 1988). In the United States, much of this work
ended with the decision to use borosilicate glass as the
waste form for U.S. defense waste at Savannah River
(Hench et al., 1984). Synroc, a titanate ceramic waste
form, was selected as the alternative U.S. waste form,
but further development in the United States ended
with the absence of funded programs. A major
research and development program for the development
of Synroc continued in Australia, culminating in
the construction of a ‘cold’ Synroc pilot-scale processing
plant. Basic research on the properties of Synroc
were continued at the Australian National University
by the late Professor Ted Ringwood and his colleagues
(see Ringwood, 1985 and Ringwood et al., 1988) and
at the Australian Nuclear Science and Technology
Organization, with collaborative work at the Japanese
Atomic Energy Research Institute and AERE Harwell
Radiation effects in nuclear waste forms 65
in the United Kingdom. Synroc is perhaps the most
thoroughly studied ceramic alternative to borosilicate
glass.
Investigations into the properties and performance
of other ceramic waste forms have revived during the
past 10 yr for application to special waste stream
compositions. At Lawrence Livermore National
Laboratory, a Mixed Waste Management Facility is
being developed in order to demonstrate an alternative
to incineration (Oversby et al., 1994). The waste form
under study is a derivative of a Synroc composition
originally developed for the immobilization of reprocessed
residues at Savannah River, and typical phases
include zirconolite, perovskite, spine], nepheline and
rutile. Although the radioactivity is low, this does
illustrate the ubiquity of a rather limited set of phases,
many of which are discussed in this paper. At the Idaho
National Engineering Laboratory, an iron-enriched
basalt waste form is under development, and the
addition of ZrO, and TiO, has produced zirconolite
as an actinide host in a silicate ceramic, i.e. a basalt
(Reimann and Kong, 1994). At Argonne National
Laboratory glassy slags have been developed for
mixed wastes with high metal contents (Feng et al.,
1994). The glassy slag may have crystalline components
which are as high as 80 vol.% in a vitreous
matrix. Spent fuel-a metal clad, ceramic oxide-has
become the major high level waste form under
consideration in the U.S., as well as in Canada,
Sweden and Switzerland, since these countries have
decided against reprocessing [see summary of recent
work in Ewing (1992b)].
1.2. Principles of nuclide isolation in solid
waste forms
Borosilicate glass is, at present, the waste form of
choice for most countries that have defense waste or
have reprocessed high-level waste (e.g. France, Great
Britain and the United States) and for most waste
compositions. The selection of borosilicate glass is
based mainly on an anticipated ease of processing
(glass frit and the waste are mixed, melted at relatively
low temperatures, and poured into canisters), the fact
that the technology is well demonstrated for actual
(radioactive) waste, and finally the assumption that
the glass as an aperiodic solid will easily accommodate
wide variations in waste stream compositions which
are extremely complex (29-30 component systems), as
well as not being susceptible to radiation-induced
transformations (it is already aperiodic). Ceramic
waste forms, with the exception of Synroc, have never
been developed to a degree that the processing
technology can be evaluated or demonstrated with
actual waste. The basis for comparison and evaluation
has been bench scale samples and usually samples that
are not radioactive. Despite the lack of data on waste
forms that have incorporated actual waste, the knowledge
of radiation effects in ceramic waste forms is
much further advanced than that of radiation effects in
glass waste forms.
In contrast to glass forms in which the radionuclides
are in principle homogeneously distributed throughout
the waste solid, ceramic waste forms may incorporate
radionuclides in two ways:
(1) Radionuclides may occupy specific atomic
positions in the periodic structures of constituent
crystalline phases, that is as a dilute solid solution. The
coordination polyhedra in each phase impose specific
size, charge and bonding constraints on the nuclides
that can be incorporated into the structure. This
means that ideal waste form phases usually have
relatively complex structure types with a number of
different coordination polyhedra of various sizes and
shapes and with multiple substitutional schemes to
allow for charge balance with radionuclide substitutions.
Extensive nuclide substitution can result in
cation and anion vacancies, interstitial defects, and
changes in structure type. The formation of polytypes
and twinning on a fine scale is common. The point
defects can themselves become sites for the radionuclides.
Except in unusual situations (e.g. monazite,
CePO,), the complexity of the waste composition
results in the formation of a polyphase assemblage
(e.g. Synroc consists of phases such as zirconolite,
CaZrTi,O,; perovskite, CaTiO,; and ‘hollandite’,
BaAl,Ti,O,,), with unequal partitioning of radionuclides
between the phases. In Synroc, actinides
partition preferentially into zirconolite and perovskite.
The polyphase assemblages are sensitive to waste
stream compositions, and minor phases form, including
glass segregated along grain boundaries. Ideally,
all waste stream elements, radioactive and nonradioactive,
are important components in the phases
formed. In some rare cases, a single phase (e.g.
monazite or sodium zirconium phosphate) can incorporate
nearly all of the radionuclides into a single
structure.
(2) Radioactive phases, perhaps resulting from
simply drying the waste sludges, can be encapsulated
in non-radioactive phases. The most common
approach has been to encapsulate individual grains of
radioactive phases in TiO, or Al,O,, mainly because
of their extremely low solubilities. This approach
requires major modifications to the waste stream
composition and special processing considerations to
keep temperatures low enough to avoid volatilization
of radionuclides. A similar approach may be taken
66 R. C. Ewing et al.
with low temperature assemblages (e.g. mixing with
concrete), but in this case there is the possibility of
reaction between the encapsulating phase and the
radioactive phases. Both the above types of waste
forms are specifically formulated for the incorporation
or encapsulation of radionuclides.
Spent fuel-a metal-clad, UO, ceramic-results
from the use of reactor fuel that is designed without
consideration of its waste form properties; but in order
to avoid reprocessing, spent fuel has become an
important potential waste form. The properties of
spent fuel as a waste form are determined primarily by
the irradiation history of the UO, in the reactor and
the disposal conditions (e.g. nearly insoluble under
reducing conditions and easily corroded under oxidizing
conditions). Radionuclides are distributed through
the fuel matrix as interstitial defects, as exsolved/precipitated
phases, along grain boundaries, or in voids
and cracks of the fuel-cladding gap. There is an
extensive literature on neutron damage effects in UO,
(Nakae et al., 1979) and a-particle damage (Weber,
1981), as well as radiation damage in naturally
occurring UO, +X (Janeczek and Ewing, 1991). A
summary of the work on spent fuel is beyond the scope
of this review; thus, radiation effects in UO, as a
nuclear waste form will not be discussed.
1.3. Types of nuclear wasteforms
The following sections give brief summaries of the
characteristics of the different waste forms. A careful
description of the characteristics of a waste form is
essential in anticipating potential radiation effects. As
an example, the waste loading ultimately determines
the radiation dose and the thermal history of the waste
form (i.e. the decay of fission products is responsible
for the thermal pulse in the early history of the waste
form). The final damage state will be a function of the
balance between accumulated damage and the kinetics
of annealing of the individual phases. For polyphase
waste forms, which include both glass and ceramics,
the partitioning of actinides and fission products into
separate phases can have dramatic differential effects,
such as the selective amorphization of phases, while
others retain their crystallinity. Anisotropic expansion
of phasesthat experience different radiation doses of
different types of radiation may lead to microfractures
and disaggregation of the waste form. Despite the
complexity of crystalline, polyphase ceramics, certain
structure types are common to a wide variety of waste
forms and waste stream compositions. This summary
focuses on the present knowledge of radiation effects in
the most common phases, e.g. pyrochlore, zirconolite,
hollandite, perovskite, zircon, apatite and monazite.
Extended discussions of the properties of the waste
forms and additional details on radiation effects can be
found in Lutze and Ewing (1988).
1.3.1. Glass. Glass waste forms can be of a wide
variety of compositions, including silicate glasses,
borosilicate glasses and phosphate glasses [summarized
by Lutze (1988)]. In principle, radionuclides are
evenly dispersed throughout the glass, although noble
metal precipitates and, in some cases, crystalline oxide
precipitates enriched in radionuclides may be present.
Waste loadings (defense and commercial) are typically
in the range of 10-30 wt%.
1.3.2. Synroc. Synroc is a titanium based, polyphase
ceramic developed at the Australian National
University and the Australian Nuclear Science and
Technology Organization. The primary phases are:
zirconolite (CaZrTi,O,), hollandite-like phase
(Ba,,,(Al, Ti),O,,), perovskite (CaTiO,), and titanium
oxides (e.g. TiO,). Other phases, such as
pyrochlore (A, _,BzO,(O, OH, F), _,pH,O; A=Ca,
REE, U, Th, Pu, etc.; B =Ti, Nb, Ta) may be
abundant depending on the waste stream composition.
Minor phases also include titanates and aluminates,
as well as noble metals. Processing conditions
are sufficiently reducing to form Ru-rich and PdTerich
phases. The principal components can each
accommodate a range of radionuclides. Hollandite
can retain fission products such as Cs, Rb, Ba;
zirconolite, U, Zr, Np, Pu; perovskite, Sr and
transuranics such as Np and Pu, The variable
compositions of the waste are accepted over relatively
wide ranges by simple and complex atomic substitutions
within the structure of individual phases and by
changing the relative proportions of the principal
phases. Certain elements (e.g. Na,O greater than
2 wt% together with SiO, and Al,O,) can lead to the
formation of more soluble silicate phases (e.g. nepheline,
NaAlSiO,). Still, waste loadings up to 20 wt%
can be made without deleterious effects on the
properties of the waste form.
1.3.3. Tailored ceramics. Tailored ceramics were
developed mainly at the Rockwell International
Science Center (Harker, 1988). Another silicate and
oxide based ceramic, supercalcine, was developed at
the Pennsylvania State University (McCarthy, 1977;
McCarthy et al., 1979). These crystalline, ceramic
waste forms are also polyphase assemblages, not
different in principle from Synroc, but perhaps
Radiation effects in nuclear waste forms 61
broader in conceptual design with a wider variety of
‘high-integrity mineral-like crystalline phases’
(McCarthy et al., 1979). The strategy again is to
modify the waste stream composition so as to form
specific, stable, crystalline phase assemblages. A
broader range of phases (oxides, silicates and phosphates)
are considered for radionuclide immobilization.
Actinides are incorporated into fluorite (CaF,),
zirconolite, pyrochlore, perovskite (CaTiO,), monazite
(CePO,), apatite (Ca,(PO,),F), and zircon
(ZrSiO,); Sr into magnetoplumbite (Pb(Fe,
Mn),,O,,), perovskite and hollandite; Cs into nepheline,
perovskite, magnetoplumbite, hollandite; Tc into
reduced metal alloys. Residual glass with fission
products can occur along grain boundaries. Nonradioactive
phases may be prominent, such as spine1
(MgAl,O,), corundum (Al,O,) or rutile (TiO,). As an
example, a ceramic waste form tailored for the
Savannah River defense waste consisted of magnetoplumbite,
spinel, nepheline, uraninite (UO, +X, and
higher oxides) and corundum, and a ceramic waste
form tailored for high-level Barnwell waste consisted
of pyrochlore, perovskite, monazite, fluorite and a Ru
alloy. In general, the silicate based assemblages are less
stable than the titanates in the presence of water.
Waste loadings are comparable to that in Synroc but
the higher density of Synroc (4.35 g cc-‘) versus a
typical tailored ceramic (4.05 g cc- ‘) means that a
slightly greater amount of waste can be incorporated
in Synroc for a weight-equivalent waste loading.
1.3.4. TiO,-ceramic. The TiO,-ceramic matrix
waste form was developed at the Kernforschungszentrum
Karlsruhe (Adehelm et al., 1988) and encapsulates
the waste in a rutile (TiO,) matrix. Waste
loadings of simulated waste oxides of up to 12 wt%
were attained in the laboratory scale experiments. The
final product has a low leach rate as compared to
borosilicate glass, because of the low solubility of
TiO,
1.35. Glass ceramics. Glass ceramics have been
developed at the Hahn-Meitner Institut in Berlin and
at the Whiteshell Nuclear Research Establishment in
Canada (Hayward, 1988). The Canadian glass ceramic
consists of discrete crystals of sphene (the proper
mineral name is ‘titanite’), CaTiSiO,, within a matrix
of aluminosilicate glass. The sphene structure is
capable of accepting a wide range of nuclides. Actinide
elements, as well as fission products such as Sr, can
substitute for the Ca. Larger cations, such as Cs, do
not fit into the sphene structure and remain in the glass
phase. The aluminosilicate glass matrix that remains
after sphene crystallization is rather durable and can
be produced in the glass-ceramic at lower temperatures
than those required to produce the glass alone.
Leach rates under comparable conditions are lower
than for borosilicate glass. In the Canadian repository
(granitic), sphene, instead ofperovskite, is the thermodynamically
stable phase. Optimal waste loadings are
approximately 15 wt%. A number of other glass
ceramic compositions have been studied in less detail.
These include: (1) celsian glass-ceramics which yield,
as individual phases, celsian (BaAl,Si,O,), perovskite
(CaTiO,), diopside (CaMgSi,O,), or eucryptite
(LiAlSi,O,); (2) fresnoite glass-ceramics which form
fresnoite (Ba,TiSi,O,), priderite ((K,Ba) (Ti,
Fe),O,,), pyrochlore (A, _,Bz06(0, OH, F), -”
pH,O) and scheelite (CaWO,); (3) basalt glassceramics
which form augite ((Ca, Mg, Fe),Si,O,) and
the spine1 trevorite (NiFe,O,); (4) iron-enriched
basalt glass-ceramics which form spine1 (enriched in
Fe,O,), plagioclase (CaAl,Si,O,), augite ((Ca, Mg,
Fe),Si,O,), fluorapatite (Ca,(PO,),F) and zircon
(ZrSiO,). Other glass-ceramic compositions have
been studied mainly in the hope of improving the leach
resistance of the borosilicate glasses. Perhaps the most
important feature of the glass-ceramics is that the
processing technology is not very different from the
demonstrated technology for the borosilicate glass.
The only difference is the final heat treatment required,
and depending on the specific glass-ceramic used (e.g.
the sphene glass-ceramic), there is a resulting increase
in chemical durability.
1.3.6. Monazite. Monazite developed at Oak Ridge
National Laboratory is unique in that it essentially
consists of a single phase, monazite (CePO,) (Boatner
and Sales, 1988). This composition can be synthesized
for the full range of lanthanide orthophosphates.
Depending on the lanthanide composition and the
temperature, at least three structure types are possible.
A low temperature hexagonal phase forms for compositions
in the first half of the lanthanide series (from La
to Dy). The hexagonal phase is metastable and will not
form once the structure is transformed to the monoclinic
monazite structure (Beall et al., 1981; Mullica et
al., 1985). The heavier lanthanides (Er to Y) have a
high temperature tetragonal structure isostructural
with zircon (ZrSiO,) (Mullica et al., 1990). Actinides
are incorporated into the lanthanide site. Important
divalent fission products, Sr and Ba, are also accommodated
in the monazite structure in order to provide
charge balance for tetravalent actinides. As with many
of the other ceramic waste forms, there is an important
68 R. C. Ewing et al.
improvement in chemical durability (factor of 20 in the
leach rate for selected elements) as compared to that of
borosilicate glass under comparable conditions. In
contrast to other waste form phases, the lanthanide
orthophosphates have a negative temperature coefficient
of solubility; thus, one may expect the chemical
durability in the presence of solutions to increase with
increasing temperature. These compounds also have a
high thermal stability with melting points in excess of
2000°C. Typical waste loadings for simulated U.S.
defense waste are 20 wt% (monazite waste form
density is in the range of 4.s5.0 g cc-‘).
1.3.7. Cementitious waste forms. Cementitious
waste forms are mainly considered for low-level waste
(Jiang et al., 1993; Quillin et al., 1994). Although
concrete has been used for low-level waste, FUETAP
(Formed Under Elevated Temperature and Pressure)
concrete was developed at Oak Ridge National
Laboratory and the Pennsylvania State University
(McDaniel and Delzer, 1988) and is unique in that it
uses the inherent heat of the high-level waste (as well as
external heat sources) to accelerate curing to drive off
up to 98% of the unbound water and form a hard,
dense product of improved (compared to normal
concrete) physical properties. The identification and
characterization of the specific phases formed and the
determination of the fate of radionuclides in FUETAP
have not been completed. Typical waste loadings are
of the order of 15-25 wt% (density is approximately
2 g cm-3). Leach rates are comparable to those of
borosilicate glass. In addition to radiation damage of
specific phases in cementitious materials, there is the
additional possibility of radiolysis of cement-pore
water and potential pressurization of waste canisters
due to an accumulation of the radiolytically produced
gases. Solid-state radiolysis of hydrated phases may
also affect their long-term stability.
1.3.8. Spent fuel. Spent fuel (Johnson and Shoesmith,
1988) was first considered as a possible waste
form in the early 1970s due to delays and cancellations
of fuel reprocessing programs. The continued,
depressed prices for uranium and lower projected
nuclear power generating capacity have made the
disposal of spent fuel a likely possibility. Typical fuel
consists of UO, pellets. Stacks of the pellets are clad in
Zircaloy-2 or -4. Individual pellets are pure UO,, with
95% of theoretical maximum density and with a
uniform grain size that is typically 24 pm. The
irradiated fuel is chemically and structurally different
from the unirradiated fuel. Actinides and lanthanides,
as well as some fission products, form solid solutions
with the UO,. In some cases new compounds form. Cs
and I react and form CsI, and some Cs is thought to
form cesium uranates or cesium molybdates. The
elevated temperatures (up to 1700°C) during irradiation
promote grain growth with the resulting segregation
of fission products that are incompatible with the
UO, structure. Cracking and void formation are
common, and volatile species (Xe, Kr, Cs and I) can
accumulate. Noble metal inclusions (up to 10 pm) are
also found. Despite the apparent complexity of spent
fuel as a waste form, its direct disposal in a repository
eliminates the need for reprocessing. The chemical
durability is difficult to assess because it is dependent
on the irradiation history and on the oxidation
potential of the repository. Under reducing conditions
UO, is very insoluble, but under slightly oxidizing
conditions matrix corrosion proceeds by oxidative
dissolution. Redox conditions may be influenced by uand
y-radiolysis of the solution in contact with the
spent fuel. Under the ‘extended-dry’ repository concept
(e.g. at Yucca Mountain), oxidation to U,O,,
accompanied by a volume increase and cladding
rupture, is a concern. If air enters the waste packages at
elevated temperatures, the subsequent entry of water
could result in relatively high leach rates for the spent
fuel.
1.3.9. ‘Novel’ ceramic waste forms. ‘Novel’ ceramic
waste forms include consideration of some additional,
unusual phases. Most of this work has been completed
at the Pennsylvania State University (Ewing, 1988).
One such phase is sodium zirconium phosphate,
NaZr,(PO,), (NZP). The structure consists of an
open three-dimensional framework with three distinct
atomic sties, which can accommodate a wide range of
radionuclides. At high simulated waste loadings
(20 wt%), an additional phase, monazite, forms. NZP
can also be mixed with concrete formulations to act as
a host for Cs, resulting in a CsZP. Its main advantages
he in its potential for low-temperature processing and
as a specific additive to concrete formulations. A
similar waste form ceramic is the low-temperature
hydroxylated ceramic. This consists of hydrated sheet
silicates (clays) and framework silicates (zeolites) that
in most cases form at low temperatures. Historically,
clay phases were among the first to be suggested for the
immobilization of nuclear waste (Hatch, 1953). The
principal advantage of both of these low-temperature
waste forms is that they are the common, thermodynamically
stable phases at the earth’s near surface,
the environment of most repositories.
Radiation effects in nuclear waste forms 69
1.3.10. Multi-barrier waste forms. Multi-barrier
waste forms are unique combinations of the previously
discussed waste forms and some special processing
technologies. Multi-barrier waste forms consist of a
series of glass or ceramic barriers. This is the ‘Russian
doll’ concept proposed by Roy (1979). Phases are
either encapsulated in other more durable phases or
the particles are coated with another phase. The
coatings are applied to enhance chemical durability,
mechanical strength or thermal stability. The coatings
may be conventional ceramics (Al,O,, TiO, or SiO,),
carbon products (PyC, Cr,C, or Sic), glass (borosilicate
or aluminosilicate) or metals (Ni, Si or Fe). A
common difficulty is the proper attachment ofcoatings
to substrate particles without cracking or flaking of the
layer. Differential expansion ofphases due to radiation
damage may exacerbate the cracking of the thin films.
1.4. Radiation sources and doses
The chief sources of radiation in high-level nuclear
waste forms are the b-decay of fission products (e.g.
’37Cs and 90Sr) and the a-decay of actinide elements
(e.g. U, Np, Pu, Am and Cm). There are minor
contributions to the damage process from spontaneous
fission events of the actinides, and E-neutron
reactions are additional sources of fission fragments
and neutrons, but the production rates are low.
Therefore, these latter events are not considered in this
paper. P-decay is the primary source of radiation
during the first 500 yr of storage, as it originates from
the shorter-lived fission products (e.g. the half-life of
’37Cs is 30.2 yr and the half-life of 90Sr is 28.1 yr). The
p-decay of fission products is responsible for heat
generation and the elevated temperatures early in the
history of waste form storage. The r-decay damage
becomes dominant as a radiation damage source after
approximately 1000 yr (e.g. the half life of 239Pu is
2.411 x lo4 yr).
The /I’- and E-decay events can cause radiation
damage through three processes in the waste form: (1)
elastic collisions between nuclear particles (e.g.
/?-particles, cc-particles and a-recoil nuclei) and the
atoms in the host matrix which cause atomic displacements
and the creation of isolated Frenkel defect pairs
(combined interstitials and vacancies) or intense
collision cascades; (2) ionization effects associated
with the y-rays, P-particles and a-particles; (3) the
transmutation of radioactive parent nuclei into different
elements. Of these three effects, the most important
are processes that cause atomic displacements, as these
are responsible for the atomic-scale rearrangement of
the structure and hence lead to the greatest changes in
physical and chemical properties. Each a-decay event
results in approximately 1500 atomic displacements,
while a p-decay event results in only 0.15-0.10
displacements. There is no compelling evidence that
ionization events result in important structural
changes in most oxide ceramics (Weber et al., 1981,
1982,1984; Weber and Roberts, 1983); however, more
recent work by DeNatale and Howitt (1985,1987) and
summarized by Weber (1991a) has shown that complex
borosilicate glasses can decompose by an ionization-driven
radiolytic process that produces bubbles
which contain molecular oxygen. Although systematic
studies are lacking, the bubble formation, glass
decomposition, and even changes in chemical durability
may correlate with the ionizing component of
energy deposition. These phenomena will be discussed
in Section 4.
The cumulative doses for the waste forms can be
substantial. Table 1 summarizes the anticipated cumulative
doses for a, p and y radiations for waste glasses
that will be generated at the Savannah River Plant in
South Carolina. If one converts these values to atomic
displacements per gram or absorbed dose, the glass
waste form with a 25 wt% waste loading will reach
values in excess of 10” atomic displacements g-’
(=0.0025 displacements per atom, dpa) or 7 x 10”
rads in only 1000 yr (Figs 1 and 2). Also included in
Figs 1 and 2 are the cumulative doses generated within
non-U.S. commercial waste glasses, which are several
orders of magnitude higher than those for defense
waste. Figure 3 shows the accumulated a-decay event
Table 1. Cumulative decay events and doses projected for SRP nuclear waste glasses (Weber, 1991a)
Time
(years)
1
10
IO2
lo3
lo4
lo5
106
Alpha decays
(per gram)
2 x 10’4
2 x 10’5
2 x 10’6
1 x 10”
3 x 10”
5 x 10”
1 x 10’8
Beta decays
(per gram)
5 x 10’6
5 x 10”
2 x 10’8
5 x 10’8
5 x 10’8
5 x 10’8
5 x 10’8
Alpha dose
(rad)
2 x 10’
2 x 108
2x 109
9x 109
3 x 10’0
4 x 10’0
8 x 10’0
Beta/gamma dose
(rad)
6x 10’
6 x lo9
3 x 10’0
6 x IO”’
8 x 10’0
1 x 10”
1 x 10”
70 R. C. Ewing et al.
I_
I_
ALPHA DECAY
ALPHA DECAY
BETA DECAY
KOMMERCIAL HLW)
)__--_---_--
CC.
0
/
// BETA DECAY
/ (DEFENSE HLW)
/-
____---_--
c-
0
/
/
/
10 16 103 104
Waste Storage Time, years
105 106
Fig. 1. Cumulative atomic displacements and decay events per gram in ceramic or glass nuclear waste forms. Basis:
Commercial HLW-3.3% enriched UO, fuel, burn-up of 33,000 MWD/MTHM, cooled 5 yr before reprocessing and
solidification, 25 wt% waste loading; Defense HLW-Savannah River Plant Waste, 25 wt% waste loading (see Table 1). After
Weber et al. (1982).
dose in Synroc for 20 and 10 wt% waste loadings.
Note, that for a 20 wt% waste loading, the accumulated
dose after lo4 yr is 4 x 10” a-decay events g-’
(=0.32 dpa). This is well within the range of doses in
which one may expect amorphization of actinidebearing
phases. Phases such as zircon can become
aperiodic (or metamict) at doses as low as 0.3 dpa.
2. RADIATION DAMAGE PROCESSES
When solids are subjected to irradiation, any or all
of three responses are possible: heating, localized
displacement of the constituent ions, and disordering
or other global rearrangements of those ions. Further,
the irradiating particles themselves can in some cases
have a significant effect on the evolution of the damage
microstructure, through their deposition in the structure.
In this section we discuss how damage effects vary
with type of radiation (which in nuclear waste
principally encompasses a, cl-recoil, p and y radiation),
and how damage mechanisms can vary with the type of
waste form.
2.1. Heating
When energy from any type of radiation is absorbed
in a solid, an increase in temperature will result. The
magnitude of the temperature increase depends on the
rate of energy absorption, physical properties of the
solid (e.g. its specific heat), and the extent of thermal
conductance to the surroundings. In the case of
nuclear waste, significant heating is possible: it can be
shown that commercial nuclear waste forms containing
typical waste loadings will, when placed in a
Radiation effects in nuclear waste forms
1oSl I I I I I I
1 10 102 103 104 105 106
WASTE STORAGE TIME, YEARS
71
Fig. 2. Ionizing radiation dose in commercial and defense high level forms (see basis for calculations in Fig. 1). After Weber and
Roberts (1983).
Synroc Age, years
Fig. 3. The r-decay dose for Synroc (10 and 20 wt.% loading of HLW). The a-decay damage is expressed as cc-decay events g- ’
and displacements per atom (dpa) and correlated to an equivalent Synroc ‘age’. After Ringwood et al. (1988).
repository, generate enough self-heat from decay of 2.2. CL-and a-recoil irradiation
fission products to result in an initial storage tempera- Self-damage in nuclear forms results primarily from
ture as high as 600°C. Even 100 yr after emplacement, decay of a-emitting actinide isotopes. The a-particle
the temperature may still be as high as 300°C. (helium ion) has an energy of 4.5-5.5 MeV; whereas,
Temperature increases of this magnitude can pro- the recoil nucleus has an energy of 7&100 keV; thus
foundly affect the response of nuclear waste forms to the a-particle carries 98% of the energy of the decay
self-irradiation damage. event. In estimating the consequences of irradiation of
12 R. C. Ewing et al.
a solid by light and heavy particles of varying energies,
it is important to estimate whether a given particle will
deposit its energy primarily by displacive (elastic) or
by ionization (inelastic) processes. This is determined
by the relative velocity of the bombarding particle and
that of the orbital electrons of the target ion. If the
particle velocity is below that of the orbital electrons,
the likelihood of electronic excitation is small and most
of the energy will be transferred to the nucleus of the
ion. However, if the particle velocity is higher than that
of the orbital electrons excitation will dominate. As a
rough rule, inelastic processes are important if the
energy of the bombarding particle, expressed in keV, is
greater than its atomic weight. Thus an cc-particle of
atomic weight 4 and energy 5 MeV will predominantly
deposit its energy by ionization, while an a-recoil ion
of mass 240 and energy 100 keV will lose most of its
energy in elastic collisions. Reeve and Woolfrey (1980)
have estimated that the cc-particle accounts for only
6% of the displacive energy from an a-decay event.
This energy is deposited over a range of l&20 pm, with
most of the damage occurring near the end of the path
of the helium ion.
In the case of irradiation with a light but energetic
ionizing particle, such as a 5 MeV a-particle, several
sequential atomic collisions usually occur before the
particle comes to rest. Each of these events displaces a
small number of ions, so that the overall effect is a
damage event spread widely throughout the solid in
small clusters. As a consequence, recombination of
vacancies and interstitials can easily occur and this
process is enhanced by the prevalence of nearby
undamaged material.
In most solids the only structural consequence of
electronic excitation is heating. However, in those
materials subject to damage by radiolysis, breaking of
atomic bonds and displacement of atoms can result.
For example, crystalline SiO, can be transformed to
the disordered or metamict state by electrons whose
energy is too low to directly displace atoms (Hobbs
and Pascucci, 1980). With respect to nuclear waste, it
has been observed that borosilicate waste glasses can
decompose by absorption of ionizing energy with the
resultant production of molecular oxygen (DeNatale
and Howitt, 1987).
The more massive but lower energy (approximately
100 keV) a-recoil particle has been shown by Reeve
and Woolfrey (1980) to account for 94% of the total
displacement energy resulting from an a-decay event
(Van Konynenburg and Guinan (1983) calculated a
value of 89% for Synroc phases). Such a particle
travels only a short distance in a solid, of the order of
10 nm. Thus the displacement event consists of a large
number of disordered atoms in a compact microvolume.
In this case most disordered atoms are
surrounded by other disordered atoms, leading to an
increased probability of damage retention.
When an a-particle stops in a solid, it becomes a
helium atom; if present in sufficient numbers and at
high enough temperatures to allow mobility, these gas
atoms can result in the formation of gas bubbles
(Clinard et al., 1970). In some cases these aggregates
affect mechanical properties, especially if located at
grain boundaries.
2.3. P-irradiation
In nuclear waste, decay of fission products produces
high-energy electrons @-particles). Because of their
low mass, these electrons usually induce only single
displacement events (Frenkel pairs); most of the
energy of these particles is dissipated by ionization
processes, through Coulombic interactions. However,
for those waste forms in which radiolytic damage is
significant, ionization effects from B-particles could
represent a major source of damage. Additionally,
b-decay will be the principal source of heating in waste
forms during the first 500 yr of waste storage, and
therefore plays an important role in determining the
temperature (and thus damage response) of the waste
form.
Calculations have been completed by Weber et al.
(1982) to determine the cumulative number of nonradiolytic
atomic displacements in commercial and
defense high-level waste that will result from the two
most effective sources of displacement damage, namely
a-decay and b-decay. The results, shown in Figs 1 and
2, demonstrate that a-decay is dominant and that
irradiation fluxes (and therefore heating effects) are
greatest early in the storage cycle.
2.4. y-irradiation
Ionizing radiation such as y-rays can, through
ionization processes, cause displacement damage or
even amorphization if the solid is sensitive to radiolysis.
Another source of damage related to y-irradiation
is the energetic electrons that result from interaction of
high-energy electromagnetic radiation with solids.
However, the inefficiency of production of these
electrons and their inefficiency in producing displacement
damage means that this is not a significant source
of radiation-induced changes in nuclear waste materials.
Total ionizing dose as a function of the age of the
waste form is illustrated in Fig. 2.
Radiation effects in nuclear waste forms 13
3. RADIATION SIMULATION TECHNIQUES
3.1. Actinide-doping of wasteforms and component
phases
Alpha-decay events associated with the decay of
actinides and their daughter products are responsible
for the atomic-scale structural damage that will occur
in radioactive waste forms. The cc-decay event consists
of two, separate, simultaneous processes: (1) An
a-particle (5 MeV) with a range of 10,00&20,000 nm
dissipates most of its energy by ionization; however, at
low velocities near the end of its track, it displaces
lo&200 atoms creating Frenkel defect pairs. (2) The
a-recoil atom (0.1 MeV) with a range of 10 nm
produces 100&2000 atomic displacements creating
‘tracks’ of disordered material. These two damaged
areas are separated by thousands of unit cell distances
and, as is illustrated in later sections of this review, will
have different effects on the atomic structure of the
solid, particularly for crystalline materials (see particularly
the work by Weber (1981) on UO,). The
long-term effects of a-decay events can be simulated by
incorporating highly-active, short-lived actinides,
such as 238Pu (half-life of 87.7 yr) or 244Cm (half-life of
18.1 yr) into the waste form phase in concentrations
great enough that the total a-decay dose reaches values
of lO’S-10’9 a-decay events g-r in reasonable
amounts of time (this may still require several years).
This is the accepted test procedure drafted by the
International Standards Organization and is the basis
for the MCC-6 method for the preparation and
characterization of actinide-doped waste forms
(Weber and Turcotte, 1982). This test effectively
simulates the simultaneous CL-and a-recoil effects to
doses that one expects over long periods of time
(103-lo6 yr for waste forms and 10s yr for natural
phases), but at much accelerated dose rates. However,
as illustrated by the study of zircon (Section 5.2.5),
actinide-doping experiments, with a high dose rate, do
accurately simulate radiation effects at very much
lower dose rates over geologic time. This is also
supported by the comparisons presented by Van
Konynenburg and Guinan (1983).
3.2. Actinides in minerals as natural analogues
Naturally occurring phases ( = minerals) contain
238u
235U and 232Th and their daughter products.
Concentrations may be as high as 30 wt% UO, in
pyrochlores, which, depending on the age of the
sample, may reach doses of 10” a-decay events g-l.
Other phases, such as zircon, contain trace amounts of
uranium (up to 5000 ppm), but with ages that are up to
lo9 yr, the doses may reach 1Or9 a-decay events g-‘.
In many cases, these doses are sufficient to cause a
radiation-induced periodic-to-aperiodic transition,
that is the metamict state. Metamict minerals are one
class of amorphous materials that have long been
recognized to have resulted from radiation damage
(Ewing, 1994), and their properties have been summarized
by Pabst (1952), Ewing (1975) and Ewing et
al. (1987). Minerals that are isostructural with crystalline
phases in ceramic waste forms can therefore serve
as natural analogues in the investigation of radiation
damage effects (Ewing and Haaker, 1980; Ewing et al.,
1988). Minerals have the advantage of representing
long-term radiation effects, over hundreds of millions
of years, which are the result of very low dose rates.
The dose rates of the actinide-doping experiments
described above and the dose rates experienced by
minerals span a range of nine orders of magnitude.
Unfortunately, there are no natural glasses which
contain enough uranium and thorium that they can
serve as natural analogues for radiation effects in waste
form glasses. There is, however, one instance in which
a Th-doped glass, 17 yr in age, has been studied (Eyal
and Ewing, 1993).
3.3. Neutron irradiations
Neutron irradiations can be utilized in three ways
(Weber and Roberts, 1983). Fast neutrons dissipate
their energy by elastic collisions resulting in numerous
atomic displacements, and the samples become
moderately radioactive. This provides a fair simulation
of a-recoil damage, but there is no simulation of
the effect of helium build-up due to an absence of
a-particles. Additionally, the only phases in polyphase
waste forms that will experience a-recoil damage
are those which contain actinides. All phases will
experience a-particle, p-particle and y-irradiation.
Neutron irradiations damage all of the phases in the
waste form, and therefore, will not simulate the phase
selective a-decay damage experienced in actual waste
forms. Neutron irradiations have been used to simulate
damage in Synroc phases (Reeve and Woolfrey,
1980; Reeve et al., 1981). Using another approach
(Malow and Andersen, 1979), boron-containing
glasses (e.g. the borosilicate glass) or boron-doped
waste forms can be irradiated in a thermal neutron
flux, and because of the high ‘OB(n, a)‘Li crosssection,
an a-particle (E= 1.78 MeV) is generated.
With this technique, high rates of He formation are
possible, but this does not simulate the a-recoil
damage of the a-decay event. Finally, thermal neutron
capture by selected nuclides (e.g. 235U) can lead to
nuclear fission, which results in the formation of
extensive zones of atomic displacements, i.e. fission
14 R. C. Ewing et al.
tracks (Antonini et al., 1979). This is not a good
simulation of a-decay damage (Weber, 1981), and the
number of fission events in nuclear waste forms is so
low as to make the process an unimportant damage
mechanism.
3.4. Charged-particle irradiations
Charged-particle irradiations using electrons
(Hobbs and Pascucci, 1980) protons, cc-particles
(Weber, 1981,1982,1985) or heavy ions (Weber et al.,
1994) have been used to study radiation damage
effects. Significant doses can be reached in short
periods of time (e.g. minutes), but analysis of the
results is difficult because the damaged areas are thin
surface layers of restricted lateral extent. The high
surface area can act as a sink for migrating defects, and
the dose required for amorphization may be greater
than that for bulk irradiations. Implanted layers can
also change the characteristics of the target material
(e.g. an increase in the apparent leach rate) as a result
of compositional changes that are enhanced in thin
layers. Still, the x-particle irradiation is a good
simulation of a-particle damage, and heavy-ion
implantation (Pb’+) provides a good simulation of
a-recoil damage (Headley et al., 1982). In fact, the two
techniques used together can be used to unravel the
relative contributions of a-particle and cl-recoil
damage to the changes in the properties of the solid.
All of these techniques have been used to study cc-decay
damage effects in nuclear waste glasses, nuclear waste
ceramics and component phases.
Most recently, Weber et al. (1994) have completed a
detailed study of the radiation-induced crystalline-toamorphous
transition in zircon in which they made a
comparison between damage caused by a-decay events
(238Pu-doped synthetic zircon and U-bearing natural
zircons) and ion-beam irradiated zircon
(0.8 MeV Ne’, 1.5 MeV Ar’, 1.5 MeV Kr+,
0.7 MeV Kr+, 1.5 MeV Xe’). The results demonstrated
that the amorphization dose under the high
dose rates of the ion irradiation (10-4~10-3 dpa s-i)
for heavy-ion irradiations is nearly identical to that of
Pu-doped zircon (3 x lo-’ dpa s- ‘), which suggests
that the amorphization process is largely independent
of the damage source or the dose rate for heavy
particles (cc-recoil particles or heavy ions) at constant
temperature. This suggests that heavy ion beam
irradiations can be used to simulate cc-decay event
damage in individual phases in nuclear waste forms.
Irradiations with light ions (e.g. a-particles) can be
used to simulate a-particle damage in phases that do
not contain nuclides which decay by cc-emission.
3.5. y-irradiations
Gamma-irradiations are easily completed on waste
forms using 6oCo or 137Cs sources. The advantage of
this type of irradiation is that y-rays are so penetrating
that waste forms can be irradiated in sealed containers.
Dose rates on the order of 2.5 x lo6 R h- ’are easily
achieved. However, structural damage from
y-irradiation is minimal except in those waste forms
that are subject to damage by radiolysis (Hobbs and
Pascussi, 1980). In terms of waste form corrosion, the
most important effect of y-radiation is the radiolytic
decomposition of species such as nitrogen and carbon
dioxide which are present in air or dissolved in the
water, leading to the formation of nitric and carboxylic
acids. The oxidation dissolution of UO, can result
from y-radiolysis of water in contact with spent fuel
(Sunder et al., 1992).
4. RADIATION EFFECTS IN GLASS WASTE FORMS
The effects of radiation in nuclear waste glasses are
complex, and the fundamental understanding of the
radiation damage processes and results is limited. The
high-radiation environment provided by the a- and
p-decay of radionuclides in nuclear waste can affect
radionuclide release to the biosphere through radiation-induced
physical and chemical changes in the
nuclear waste glass at both the atomic and macroscopic
levels. Several comprehensive reviews (Burns et
al., 1982a,b; Weber and Roberts, 1983; Day et al.,
1985; Weber, 1988, 1991a; Wronkiewicz, 1993) of
radiation effects in nuclear waste glasses provide
excellent technical assessments, and many unresolved
scientific and engineering issues regarding radiation
effects were summarized by Matzke (1988a,b) following
a 1987 workshop.
4.1. Volume changes
Radiation effects from a-decay in actinide-doped
simulated nuclear waste glasses can result in either
expansion or compaction of the glass structure. The
macroscopic volume changes, AV/V,, are usually
determined by density measurements and generally
follow an exponential dependence on dose, D, of the
form:
AV/V,=A[l-exp(-BD)], (1)
where A is the volume change at saturation and B is the
amount of glass damaged per unit dose. The volume
changes normally saturate within the range of f 1.2%
at a dose equivalent to about 1018 cc-decays g-’
(10” rad), as illustrated in Fig. 4 for several actinidedoped
nuclear waste glasses studied at the Pacific
0.8
-0.8
-1.2
'0
Radiation effects in nuclear waste forms 15
ALPHA IONIZATION, 10’ ’ rad
1 2
I I
1 2 3
CUMULATIVE DOSE, 1018 ALPHA DECAYS/g
_..
Fig. 4. Volume changes versus dose in ‘*Ym-doped simulated Pacific Northwest Laboratory waste glasses. After Weber and
Roberts (1983).
Northwest Laboratory in the U.S. (Weber and
Roberts, 1983; Weber, 1988). Similar behavior has
been observed in a collaborative study on actinidedoped
simulated waste glasses by the Centre de
Marcoule, the Hahn-Meitner Institut, and the Atomic
Energy Reserch Establishment (Day et al., 1985) as
shown in Fig. 5. Several proposed glass compositions
for Savannah River Plant waste have also been studied
(Bibler and Kelley, 1978; Bibler, 1982; Weber, 1988),
and the volume expansions (Fig. 6) are larger than for
other waste glasses studied to date. In a more recent
study of an actinide-doped waste glass by Inagaki et al.
(1992), the volume expansion determined by density
measurements at a dose of approximately 5.5 x 10”
a-decays g- ’ ranged from 0.4 to 0.6%, which they
found to be in excellent agreement with the volume
change (0.51%) determined from the measured size
and concentration of irradiation-induced bubbles.
These results show a strong correlation between
bubble formation and macroscopic volume change in
nuclear waste glass and also support the hypothesis
suggested by Weber (1988, 1991a) that the volume
expansions observed in earlier studies (Figs 4-6) may
partly be due to the formation of bubbles. Banba et al.
(1994) have also observed a volume expansion of
0.75% after 1.5 x 10” a-decay events g-i in an
actinide-doped waste glass from the Tokai Reprocessing
Plant. The bubbles may be of trapped He from the
a-decay event or oxygen from radiolytic decomposition
of the glass. The increases in volume of neutronirradiated
nuclear waste glasses (Antonini et al., 1980;
Sato et al., 1988) are also within the bounds observed
for the actinide-doped glasses. Volume changes have
not been measured in any of the ion-irradiation
studies. In the simulated waste glasses that have been
studied, there is a general tendency for swelling in
glasses with low alkali/silicon ratios and for compaction
in glasses with high alkali/silicon ratios (see
shaded area of Fig. 7).
Gamma-irradiation studies in the U.S. on two
simulated nuclear waste glasses have reported only
slight ( < 0.1%) volume changes (Bibler, 1982; Weber,
1988). Shelby (1980), however, reported up to 1%
compaction in commercial borosilicate glasses after
y-irradiation to 10” rad, which is comparable to the
compaction observed in some actinide-doped glasses
(Fig. 4). Sato et al. (1984) have reported swelling of
0.2% in a simulated nuclear waste glass after
y-irradiation; Manara et al. (1984) and Sato et al. (1984)
also observed swelling of up to 50% (from oxygen
bubble formation) after electron-irradiation. Based on
these results, it is impossible to determine conclusively
whether the volume changes observed in a-decay
studies are due to atomic displacements or ionization
R. C. Ewing et al.
ALPHA IONIZATION, 10” rad
,.o[ 0;5 1; l.,
0.8 -
0.6-
s!
-0.4 SON
56.30.2O.U2
-0.61 I I I
0 0.5 1 .o 1.5 2
CUMULATIVE DOSE, lo’* ALPHA DECAYS/g
Fig. 5. Volume changes versus dose in 238Pu-doped simulated European nuclear waste glasses. After Day et al. (1985).
ALPHA IONIZATION, 10’ 1rad
0.5 1 .o
CUMULATIVE DOSE, 1O’* ALPHA DECAYS/g
Fig. 6. Volume changes versus dose in 244Cm and Z38Pu-doped simulated Savannah River Plant HLW glass compositions.
After Weber (1988).
Radiation effectsin nuclear waste forms 77
??VG 9613
- 1.2t
I 1 I I I I I 1 I
0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.6
J
Alkali/Silicon Ratio
Fig. 7. Saturation volume change versus alkali/silicon ratio in simulated nuclear waste glasses.After Weber (1991a).
damage. However, the ion-irradiation studies of
Arnold (1985, 1988) have strongly indicated that
ionization processes are the dominant mechanism for
producing volume changes in borosilicate waste
glasses. Consequently, the volume changes in actual
waste glasses, if due to ionization damage, could
saturate within 10,000 yr for U.S. defense glasses and
within several hundred years for commercial waste
glasses in other countries. While the magnitude of the
above volume changes does not pose any problems for
the safe storage of nuclear waste, they indicate changes
in microstructure that may affect the integrity of
nuclear waste glasses and radionuclide release rates
(which are proportional to the surface area in contact
with solutions). Furthermore, differential volume
changes as a result of temperature dependencies,
which have not been investigated, and compositional
inhomogeneities may result in microcracking, residual
stresses, additional loss of integrity, and enhanced
radionuclide release.
4.2. Stored energy
The interactions of radiation with nuclear waste
glasses can result in the storage of latent energy
associated with the changes in structure and bonding
that is released as heat at elevated temperatures.
Studies on actinide-doped simulated nuclear waste
glasses at the Pacific Northwest Laboratory (Roberts
et al., 1976; Weber and Roberts, 1983), AERE-Harwell
(Hall et al., 1976), and under a collaborative European
Communities program (Malow et al., 1980) indicate
that a-decay can result in stored energy that is
generally less than 150 J g- ’and saturates at a dose of
0.1-0.3 x 10” a-decays g- ’ (10” rad). This is illustrated
in Fig. 8, where the total stored energy for
several simulated waste glass compositions has been
determined using differential scanning calorimetry
(Weber and Roberts, 1983). In general, the increase in
stored energy with dose followed exponential behavior
similar to that shown by the changes in volume;
however, the dose for saturation of stored energy is
considerably less than the dose required for saturation
of volume changes, which suggests different mechanisms
and associated defects. Primak and Roberts
(1984) have postulated that this stored energy is
associated with the breaking of glass network bonds
and may be due to ionization damage from the
a-particle. If this is the case, stored energy from
ionization damage associated with B-decay could
saturate in 30 yr. Stored energy data on neutronirradiated
waste glasses (Roberts et al., 1976; Antonini
et al., 1980; Cousens and Myhra, 1983) are in good
agreement with the results in Fig. 8. No data are
available on stored energy release in y- or electronirradiated
glasses to evaluate the role of ionization in
the absence of significant displacive energy deposition.
The stored energy in actinide-doped simulated
waste glasses is released over a broad temperature
range from 100 to 600°C (Roberts et al., 1976; Weber
and Roberts, 1983). Since the specific heat ofthe nuclear
waste glasses is of the order of 1 J (g.K)-‘, the
maximum temperature rise that could occur from the
rapid release of this stored energy is only 50-125 K.
Temperature excursions of this magnitude should
78 R. C. Ewing et al.
ALPHA IONIZATION, 10” rad
0 0.2 0.4 0.6
1501
I I I
w
n 50
kz
:
I
0 PNL 76-68
•I PNL 77-260
A PNL E-G
??PNL P-G
0 PNL 72-68-a
Fig. 8. Stored energy, measured by differential scanning calorimetry, as a function of dose in 244Cm-doped simulated nuclear
waste glasses. After Weber and Roberts (1983).
I I I I I
0 0.2 0.4 0.6 0.8
CUMULATIVE DOSE, lo’* ALPHA DECAYS/g
present no serious problems. Furthermore, the results
of Roberts er al. (1976) indicate that the stored energy
decreases almost linearly with increasing storage
temperature. Thermal recovery of the stored energy in
the PNL glasses (Fig. 8) was investigated by 7-day
isochronal anneals (Weber and Roberts, 1983). The
results show a single recovery stage that is similar for
all the waste glass compositions, with full recovery
occurring at 360°C (Fig. 9). These results are similar to
those of Roberts et al. (1976) for the effects of storage
temperature.
4.3. Transmutations
Nuclear transmutation is the formation of a new
chemical element as a result of nuclear reactions or
radioactive decay processes. The principal source of
transmutations in nuclear waste is b-decay of the
relatively abundant fission products cesium and strontium.
Transmutation of these two elements is accompanied
by changes in both ionic radius and valence.
Cesium (1 + ) decays to barium (2 + ) with a decrease in
ionic radius of 20%, and strontium (2+) decays to
yttrium (3 + ), which in turn decays to zirconium (4 + )
with a final ionic radius decrease of 29%. In an
accelerated study of cesium to barium transmutation
in five glasses and two ceramics, Gray (1982)
found no significant effect of the transmutations on
microstructure, chemical durability or density.
4.4. Helium accumulation
Helium, which results from the capture of two
electrons by an a-particle, must be accommodated
interstitially or substitutionally, diffuse to internal
defects, or be released at the glass surface. Based on a
number of studies (Hall et al., 1976; Turcotte, 1976;
Malow and Andresen, 1979; Malow et al., 1980) in
which helium diffusion and release in nuclear waste
glasses were measured, it can be concluded that the
helium generated will accumulate in the glass, and
only a small fraction can be expected to be released at
ambient temperatures. Observed decreases in diffusion
coefficients with increasing dose (Turcotte, 1976;
Malow and Andresen, 1979) suggest trapping of
helium at radiation-induced defects. Low-level helium
concentrations (up to 100 atomic parts per million,
appm) introduced both by (n, c() reactions under
neutron irradiation and by accelerator implantation
have resulted in bubble formation in simulated nuclear
waste glasses after post-irradiation annealing at 600°C
in one case (De et al., 1976a) and 750°C in another
(Malow and Andresen, 1979). When higher helium
concentrations (up to lo4 appm) are introduced by
(n, a) reactions, bubble formation was observed at
Radiation effects in nuclear waste forms 19
ambient irradiation temperatures (< 100°C) by De et
al. (1976a) and at slightly elevated irradiation temperatures
(~230°C) by Sato et al. (1988). The exact
nature of the bubbles was not determined in either
case, but they are assumed to be helium or helium/oxygen
bubbles. In a more recent study, Inagaki et al.
(1992) observed bubble formation in 238Pu- and
244Cm-doped glasses as a result of a-decay. The
change in glass density determined by observed bubble
size and abundance correlated well with the measured
change in density.
4.5. Microstructural changes
The principal radiation-induced changes in the
microstructures of nuclear waste glasses are phase
changes, phase separation, microfracturing and
bubble formation. Thermally-induced devitrification
(Jantzen et al., 1984; Spilman et al., 1986) and phase
separation (Tomozawa et al., 1979) in HLW glasses
are well known. Radiation-enhanced devitrification
of nuclear waste glasses has not been investigated
to date. Phase separation has been induced in
simulated waste glasses by electron-irradiation
(DeNatale and Howitt, 1984 and 1985). Crystalline
phases that often form in waste glasses from thermal
devitrification during cooling and storage (additionally,
crystalline phases form in the glass during
1.0
solidification from the melt) can undergo a radiationinduced
crystalline-to-amorphous transformation. In
a study of a Cm-doped waste glass containing
crystallites of Ca,(Gd, Cm),(SiO,),(PO,)O, (apatite
structure) and (Gd, Cm),Ti,O, (pyrochlore structure),
the differential expansion (5-8%) associated
with the crystalline-to-amorphous transformation of
these phases resulted in significant microfracturing as
a result of radiation-induced differential volume
changes as illustrated in Fig. 10 (Weber et al., 1979;
Weber and Roberts, 1983). The behavior of these two
particular phases is discussed in detail in Sections
5.2.1. and 5.2.6. This radiation-induced microcracking
can significantly increase the surface area for radionuclide
release.
As noted above, helium implantation in waste
glasses can result in the formation of helium bubbles.
Oxygen bubbles can also form and were first observed
by Hall et al. (1976) in simulated nuclear waste glasses
irradiated with electrons in a high-voltage electron
microscope at 200°C to simulate the effects of
b-irradiation. Since then, irradiation-induced oxygen
bubble formation has been observed by others in a
wide variety of glasses as a result of electron-beam
irradiations (Todd et al., 1960; Manara et al., 1982,
1984; Sato et a/., 1983; DeNatale and Howitt, 1984,
1985,1987; Arnold, 1985; DeNatale et al., 1986; Heuer,
1987; Heuer et al., 1986) ion-beam irradiations
0 PNL 72-66 a
fl PNL 77-260
V PNL 76-66
0 PNL E-G
m PNL P-G
0.0
0 200 400 600
ANNEALING TEMPERATURE, “C
Fig. 9. Thermal recovery of stored energy in “Yrn-doped simulated nuclear waste glasses. After Weber and Roberts (1983).
(DeNatale et al., 1986; DeNatale and Howitt, 1987;
Heuer, 1987), and y-irradiations (DeNatale and Howitt,
1985, 1987; Heuer, 1987; Howitt et al., 1991). The
presence of a gas phase in the bubbles has been
confirmed in the studies by Hall et al. (1976) and Heuer
et al. (1987). The formation of oxygen bubbles, which
results from radiolytic decomposition, has been correlated
with the ionization component of energy deposition
and not the ballistic (atomic displacement)
component (DeNatale et al., 1986), in agreement with
the results of Arnold (1985, 1988). Consequently,
ionization effects from a-, p- and y-radiation are more
important to nuclear waste glass behavior than
originally assumed. Although high-resolution microstructural
characterization was not performed in the
early studies of actinide-doped nuclear waste glasses,
Weber (1988, 1991a) postulated that similar bubble
formation from cc-particles emitted in actinide-doped
nuclear waste glasses may occur and could account for
the volume expansions observed in Figs 4-6. Recent
results by Inagaki et al. (1992) confirm that bubbles,
probably of He, with radii of 20&300 nm are formed
as a result of a-decay in a 244Cm- and 238Pu-doped
simulated nuclear waste glass and are primarily
responsible for the measured volume expansion.
4.6. Radiolytic decomposition
The early emphasis of radiation effects studies in
nuclear waste glasses on displacement damage has led
to some misconceptions concerning waste glass
stability. Although not originally anticipated, it is now
evident that complex borosilicate glasses, including
those proposed for the storage of nuclear wastes, can
decompose by an ionization-driven radiolytic process
that produces molecular oxygen and may lead to
oxygen bubble formation as described above. The
initial observation of oxygen bubble formation by Hall
et al. (1976) did not create much concern because the
electron dose required to produce the bubbles was
quite high, 1012-1013 rad, exceeding the dose equivalent
to a million years of storage. Nearly 10 yr later,
DeNatale and Howitt (1985) reported that gamma
irradiation of simulated nuclear waste glasses at
significantly lower doses (lo9 rad) will result in bubble
formation; this observation was of significant relevance
to nuclear waste glass behavior. Subsequent
y-irradiation studies (DeNatale and Howitt, 1987;
Heuer, 1987; Howitt et al., 1991) have concluded that
oxygen bubbles form (Fig. 11) in a wide variety of
glasses at low doses. However, in another study
(Tosten, 1989; Biber et al., 1990) of one of the same
Radiation effects in nuclear waste forms 81
Fig. 11. Bubble formation in a simulated Savannah River Plant HLW glass gamma irradiated to 1.3 x lo9 rad. Micrograph
courtesy of J. P. Heuer and D. G. Howitt, University of California-Davis.
glass compositions, y-irradiation-induced decomposition
or bubble formation was not observed, even at an
order of magnitude higher dose. These authors suggest
that bubble formation in y-irradiated glasses which
were reported by others may be an artifact of sample
preparation. Several quantitative studies have shown
that the kinetics of the bubble formation process are
consistent with the motion of alkali metal cations and
not oxygen (DeNatale and Howitt, 1984; Lava1 and
Westmacott, 1980). The formation of oxygen bubbles
is apparently brought about by radiolytic decomposition
of the ionic component of the glass, followed by
the migration of the cations into the glass and the local
precipitation of molecular oxygen (Howitt et al.,
1991). The bubbles observed in an actinide-doped
simulated waste glass (Inagaki et al., 1992) also may be
produced by this mechanism, although the role of the
helium (or-particles) and the composition of the
bubbles (oxygen and/or helium) has not been estab-
lished.
Irradiation studies of simulated waste glasses carried
out by DeNatale and Howitt (1987) and Heuer
(1987), utilizing y-radiation, ion beams, and electron
beams at similar temperatures, indicate a significant
dose-rate effect for glass decomposition and subsequent
bubble formation (Fig. 12). Decomposition
under y-irradiation occurs at much lower doses than
for either ion- or electron-irradiations. The process for
10s 10s 10’0 10” lo’* 10'3 10'4 10
Ionizing Dose, rad
5
Fig. 12. Dose and dose rate dependence of ionization-induced
bubble formation in simulated HLW glasses. Adapted from
DeNatale and Howitt (1987) and Heuer (1987).
82 R. C. Ewing et al.
y-induced decomposition under ambient conditions
(< 100°C) has an onset at about 10s rad, saturates at
lo9 rad, and results in a uniform distribution of
bubbles (Fig. 11) with an average diameter of 556 nm
(Heuer, 1987). Under electron irradiation, bubble
formation occurs over a broad temperature range with
a maximum in the rate of formation at 250°C
(DeNatale and Howitt, 1984; Heuer, 1987) as illustrated
in Fig. 13. Similar behavior is reported by
Manara etal. (1984) in electron-irradiated borosilicate
glasses; however, the peak in bubble formation shifts
with Na content. The temperature dependence of
radiolytic decomposition and bubble formation under
y-irradiation has not been investigated (although
DeNatale and Howitt (1985) reported enhanced
bubble formation at 50”(Z), but given the strong doserate
dependence in Fig. 12, a shift in the temperature
dependence from that shown in Fig. 13 is expected.
The implications of these results may be more
serious than other irradiation studies have suggested,
as this radiolytic decomposition process, if operable,
can be completed in relatively short time periods
(< 10 yr) for actual nuclear waste glasses. In addition,
the low-dose character of this effect underscores the
need for a better understanding of the temperature
dependence of this process, since the waste glasses will
be exposed to elevated temperatures and temperature
gradients during this time period. As noted above, two
studies (Sato et al., 1984; Manara et al., 1984) have
determined significant volume expansion (50%) from
100
50
10
the size and density of observed bubbles. The swelling
associated with this radiolytic process is sensitive to
alkali content (Manara et al., 1984) but cannot account
for the decrease in the volume expansion with alkali
content (Fig. 7). Concerns raised (Tosten, 1984; Bibler
et al., 1990) about whether radiolytic decomposition
and bubble formation actually occur under y-irradiation
emphasize the need for a more fundamental
understanding of this process in complex glasses.
4.7. Mechanical properties
In a recent study of a simulated nuclear waste glass
doped with 238Pu and 244Cm, Inagaki et al. (1993)
observed an exponential decrease in hardness and
Young’s modulus with dose, while the fracture toughness
increased exponentially with dose. The maximum
values of the relative changes reported for the hardness,
Young’s modulus, and fracture toughness in
theseglasses were -25%, - 30% and +45%, respectively.
In the same study, the changes in hardness and
Young’s modulus completely recovered during
annealing above 673 K with an activation energy of
0.28 eV; the recovery of the change in fracture
toughness was initially more rapid, with an activation
energy of 0.21 eV, then the rate of recovery slowed
significantly, and complete recovery of the fracture
toughness was never observed. The changes in fracture
toughness are strongly correlated to the behavior of
bubbles, which they observed previously in these
0 100 200 300 400 500
Temperature, “C
Fig. 13. Temperature dependence of the dose to induce bubble formation under electron irradiation in HLW glasses. Adapted
from DeNatale and Howitt (1984) and Heuer (1987).
Radiation effects in nuclear waste forms 83
Ionization DoseI lo” rad
w 2.5 O
1 2
i;j _
t
b~<~~~;~
Ial
,y-•
2 l.O-
3
+ Fracture Toughness
-5
t
A Hardness
OS0 ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’ ’
1 2 3 4
FLUENCE , 1Ol6 alphas/cm2
Fig. 14. Hardness and fracture toughness versus dose in an a-irradiated simulated HLW glass. After Weber and Matzke (1987).
glasses (Inagaki et al., 1992).The bubbles, which can
impede crack propagation, anneal with kinetics similar
to the recovery of the fracture toughness.
A similar decrease in hardness (24%) was observed
in an actinide-doped glass by Bonniaud et al. (1979). In
a study of a-irradiated simulated waste glass (Weber
and Matzke, 1987), the hardness decreased 15% and
the fracture toughness increased 80% with dose
(Fig. 14). In another study of cc-irradiated waste glass
(Routbort and Matzke, 1983), a 75% increase in
fracture toughness was reported. These large increases
in fracture toughness may also be associated with
bubble formation; however, because the microstructural
characterization was not completed, the
presence and role of bubbles cannot be confirmed.
Gamma-irradiation of a simulated nuclear waste glass
to 6.6 x lo* rad indicates a slight decrease in elastic
modulus and a slight increase in hardness and fracture
toughness (Weber, 1988). In a study of a commercial
borosilicate glass (Zdaniewski et al., 1983), y-radiation
did not appreciably affect the strength or fracture
toughness below 10s rad. The maximum doses in these
studies are low; consequently, the effect of y-radiation
on mechanical properties is unknown for y-doses
equivalent to more than 1 yr of storage.
4.8. Radionuclide release
Radiation can affect the release rate of radionuclides
from waste glasses by increasing the surface area for
radionuclide release (microfracturing) and by changing
the dissolution rate of the glass. The extent of the
radiation-induced microfracturing will depend on
differential volume changes, microstructure and mechanical
properties. Presently, there are insufficient data
and understanding currently available to predict the
extent of radiation-induced microfracturing in waste
glasses, except to note that it can be expected to occur,
especially as a result of the radiation-induced differential
expansion of the crystalline phases in the glass.
The dissolution rate of nuclear waste glasses may be
affected by the radiation-induced changes in
chemistry, microstructure and network bonding. As
noted in several reviews (Turcotte, 1981; Burns et al.,
1982a,b; Weber and Roberts, 1983) and discussed in
detail recently (Weber, 1988, 1991a), almost all data
concerning the effects of a-decay on dissolution (leach)
rates of waste glasses were determined prior to 1982 in
short-term tests based solely on weight loss. The data
obtained from these tests (Fig. 15) indicate no more
than a factor of 3 increase in the short-term dissolution
rates as a result of radiation effects from cc-decay. It is
now well-established that short-term tests based on
weight loss can underestimate the radiation-induced
increases in dissolution rate by a factor of 34 due to
precipitation of alteration phases on the surface of the
glass (Westsik and Harvey, 1981; Weber et al., 1985a;
Weber, 1988); consequently, based on this limited
database, it has been suggested by Weber (1991a) that
radiation effects are not expected to increase the
nuclear waste glass dissolution rates by more than a
factor of 10. Although this may overestimate (or even
underestimate) the actual effects, it is the best assessment
that can be drawn from the limited database and
understanding that currently exists. This assessment is
supported by a recent study by Eyal and Ewing (1993)
of a Th-doped borosilicate glass, where the Th
R. C. Ewing et al.
0 2 4 6
Cumulative Dose, 1018 Alpha Decays/g
Fig. 15. Increased leaching (dissolution) factor, L/L,, versus
dose for short-term tests. After Weber (1988).
daughter product release rate was increased from 15 to
45% due to incongruent dissolution along the recoil
tracks produced by the cc-recoil nucleus.
Ion-irradiation (Manara et al., 1982; Dran et al.,
1982; Primak, 1982) and neutron-irradiation (Cousens
and Myhra, 1988) studies of several HLW glasses
have shown irradiation-enhanced dissolution rates of
up to a factor of 4. Another study (Arnold et al., 1982)
indicated no significant increase in dissolution rates.
Ion irradiation, however, is known to introduce
compositional variations (Arnold et al., 1982), which
affect the interpretation of the results. Gammairradiation
of several HLW glasses up to doses of
10” rad have lead to increased dissolution rates of up
to a factor of4 (Grover, 1973; Bibler, 1982). In another
study (Bibler et al., 1990), where radiolytic decomposition
was confirmed not to occur, y-irradiation to
3.1 x 10” rad did not result in any measurable
increase in the leach rate. The effect of radiolytic
decomposition on dissolution rates remains unclear,
because it has not been determined whether radiolytic
decomposition occurred in the glasses for which
increased dissolution rates have been measured.
As alternatives to nuclear waste glasses, several
multi-phase, crystalline ceramic waste forms have been
proposed that are tailored to produce specific crystalline
phases as hosts for the different radionuclides.
Generally, fission products (such as cesium and
strontium) are confined to one or more crystalline
phases, and the actinides (uranium, neptunium, plutonium,
americium and curium) partition into other
crystalline phases. Synroc and other related titanatebased
ceramic waste forms have received the most
attention (Ringwood et al., 1979, 1981; Ringwood,
1982; Kesson et al., 1983; Newkirk et al., 1982). There
has been more limited study of radiation effects in
supercalcine, a silicate-based, tailored ceramic
(McCarthy, 1977; McCarthy et al., 1979) and glassceramics
(D&et al., 1976a,b). There have been far fewer
studies of radiation effects in the bulk crystalline
ceramic waste forms as compared to glass waste forms.
Details of the more extensive studies of radiation
effects in component phases are presented in
Section 5.2.
5.1.1. Synroc. Ringwood et al. (1980) have argued
that Synroc is stable with respect to a-decay damage,
based on an assessment of natural minerals of uranium
and/or thorium-bearing perovskite and zirconolite,
which are the actinide host phases of Synroc. Such
data can provide only qualitative results, as natural
and actinide-doped zirconolites do become amorphous.
The rate of amorphization and subsequent
changes in physical and chemical properties may vary
with modifications in composition. Therefore, quantitative
evaluation of Synroc as a waste form requires
radiation-damage testing of actual or simulated waste
forms.
In summary, depending on the type of corrosion In the early 198Os, a limited amount of accelerated
model used for waste glass dissolution, the radiation- irradiation-damage testing was performed on Synroc
enhanced radionuclide release rates may be expected and its constituent phases by Reeve et al. (1981) and
to increase in proportion to the increase in surface area Woolfrey et al. (1982) using fast neutrons. Cold(due
to microfracturing). The actual release rate will pressed and sintered (cps) and hot pressed (hp)
depend on geochemical conditions (e.g. presence or specimens of Synroc B and Synroc C, as well as cps
absence of silica-saturated solutions). Additionally, in specimens of their constituent minerals hollandite,
the absence of careful studies, a conservative factor of perovskite and zirconolite, were irradiated in the fastapproximately
10 is estimated for the enhanced neutron flux of the HIFAR reactor at Lucas Heights,
dissolution rate due to radiation-induced changes in for different lengths of time. (Synroc B does not
the atomic structure of the glass (e.g. Eyal and Ewing, incorporate simulated fission products; Synroc C
1993). Any temperature dependence of the radiation
effects could further affect this enhancement.
5. RADIATION EFFECTS IN CRYSTALLINE WASTE
FORMS
5.1. Multiphase ceramic wasteforms
Radiation effects in nuclear waste forms 85
does.) The largest dose reported on the Synroc
specimens (10 wt% waste loading) was 2.7 x
1O26n mm2 (E> 1 MeV), which corresponds to
0.7 dpa or 8 x lo’* cl-decays g- ’ and represents an
equivalent Synroc storage time of 6.5 x 10’ yr.
Volume, density, open porosity and unit cell parameter
changes were measured. The volume expansion
of Synroc specimens was generally larger than that of
the separately irradiated specimens of the constituent
phases. In addition, the results show that the volume
expansion is much larger for the Synroc-B (cps) than
for the Synroc-B (hp). The volume changes in
Synroc-C are similar for both hp and cps specimens.
The macroscopic volume expansion results for
Synroc-C (hp) are included in Fig. 16 (the neutron
dose has been converted to an equivalent cc-decay
dose). Based on these data, volume expansion at
saturation cannot be predicted.
Accelerated testing of Synroc-C and its constituent
phases has been completed using actinide-doping
techniques at Harwell (Evans and Marples, 1985;
Evans et nl., 1986). The radiation stability of hotpressed
Synroc specimens doped with 2 and 5 wt%
238Pu was investigated. X-ray diffraction analysis
indicated that the Pu-containing zirconolite and
perovskite phases were X-ray diffraction amorphous
after a bulk average dose corresponding to 2.8 x 10”
a-decay events g-’ (estimated to be 0.37 dpa in the
zirconolite and perovskite phases), which is equivalent
to a Synroc age of approximately lo4 yr for a 10 wt%
waste loading. The macroscopic volume expansion
(Fig. 16) was determined from the change in density
(Weber, 1990). The swelling exceeds 6% and shows
only a slight indication of approaching a steady-state
(saturation) value. The swelling of the Pu-doped
, , I ,
n-irradiated SYNROC-C
Pu-doped SYNROC-C
Cm-doped Titanate Ceramic
-0 2 4 6 8 10
Dose, IO’* alpha decays/g
Fig. 16. Microscopic volume expansion in Pu-doped (Evans
et al., 1986) and neutron irradiated SYNROC-C (Reeve et al.,
1981; Woolfrey et al., 1982), and a Cm-doped titanate
ceramic (Mitamara et al., 1990).
Synroc-C occurs at about twice the rate as the
equivalent neutron-irradiated Synroc-C. This may be
due to the relatively simple assumptions made by
Woolfrey et al. (1982) in estimating the equivalent
cc-decay dose for the measured neutron dose or may
indicate that neutron irradiation does not produce the
same radiation damage effects.
5.1.2. Titanate ceramics
The Japanese have investigated a titanate ceramic
consisting of perovskite, zirconolite, hollandite, freudenbergite
and loveringite for encapsulation of sodiumrich
high-level waste. This titanate ceramic was doped
with 0.69 wt% 244Cm, and the effects of x-decay on
density and dissolution rate were investigated (Mitamura
et al., 1990). The density decreased linearly with
dose. The volume change was + 1.0% after a dose of
0.7 x 10” cc-decays g-’ (Fig. 16), which was the
maximum dose reached. The release rates measured
for Cm and soluble non-radioactive elements (e.g. Na,
Cs, Sr and Ca) showed a general trend of increase with
increasing dose. The maximum dose reached was
relatively low compared to those in other studies,
corresponding to a storage time of 5000 yr.
Sandia National Laboratories investigated a titanate
ceramic to immobilize acidic high-level nuclear
waste. This titanate ceramic is a multi-phase assemblage
that consists of rutile (TiO,), an amorphous
silicate phase, perovskite, zirconolite, hollandite and a
U-Zr-rich phase believed to have a pyrochlore
structure (Dosch et al., 1984). Radiation effects from
a-decay have been simulated by irradiating TEM
specimens with Pbf ions (Dosch et al., 1985). Multiple
energies (4G250 keV) were used to produce a uniform
damage profile. Irradiation at single energies of
24G250 keV were also performed. Examination by
transmission electron microscopy of the irradiated
specimens indicated considerable damage in all phases
after a dose equivalent to 1 x 1019 x-decay
events cm- 3. All phases remained crystalline except a
phase with a chemistry typical of a pyrochlore, which
became amorphous.
5.1.3. Supevcalcine. The only radiation damage
study that has been performed on a supercalcine
formulation (SPC-2) was carried out at the Pacific
Northwest Laboratory (Rusin et al., 1979; Turcotte et
a/., 1982). The SPC-2 formulation was doped with
2 wt% 244Cm. Stored energy, X-ray diffraction analysis,
and density measurements were made initially and
at 3-month intervals, reaching a total bulk dose of
1.2 x 10” a-decay events g - I. Initial characterization
86 R. C. Ewing et al.
showed the supercalcine consisted of three major
phases: fluorite, apatite and a tetragonal phase.
Subsequent analysis suggested that curium predominantly
partitioned into the apatite and the tetragonal
phase.
Analysis by X-ray diffraction of the Cm-doped
supercalcine, as a function of dose, showed a gradual
transformation of the apatite from the crystalline to
amorphous state, in agreement with the results of
Weber et al. (1979) on the amorphization of apatite
crystallites in a nuclear waste glass. An expansion in
the unit cell of apatite was also observed with
increasing dose; the increase in unit cell volume
saturated at 2.5%. There was a slight increase in the
intensity of the diffraction maxima associated with the
tetragonal phase, but changes in the tetragonal unit
cell could not be determined. The fluorite peak was
used as a standard, as this phase did not appear to
contain Cm and was unaffected by external radiation.
Stored energy and density measurements were also
made as a function of dose for the Cm-doped
supercalcine. The stored energy reached a maximum
value of 42 J g- ’ at a dose of 0.5 x 10” a-decay
events g 1 and decreased slightly with further increases
in dose. Energy release was not complete at
600°C (the upper limit of the calorimeter used);
therefore, additional energy release may occur at
higher temperatures (> 600°C). The density decreased
exponentially with dose, resulting in a volume expansion
of 1.4% at a dose of 1.2 x lo’* a-decay events g - I.
The macroscopic swelling as a function of dose in
SPC-2 is shown in Fig. 17. The volume expansion at
saturation (1.4%) is partially attributed to the volume
expansion in the apatite as a result of radiationinduced
amorphization.
0.0 0.5 1.0 1.5 5.2. Crystalline phases
Dose, IO” alpha decays/g
Fig. 17. Volume changes in Cm-doped supercalcine, celsian
glass ceramic (Turcotte et al., 1982), and Pu-doped celsian
glass ceramic (Malow et al., 1980).
5.1.4. Glass-ceramics. A glass-ceramic is a finegrained
mixture of glass and ceramic phase ideally
derived from a homogeneous glass by heat-treating a
glass precursor at the temperature of maximum
nucleation rate for the ceramic phases(s), followed by a
high-temperature treatment to yield a maximum
growth rate. The celsian glass ceramic, developed at
the Hahn-Meitner Institut and named for the predominant
crystalline phase, is easy to fabricate, is
homogeneously crystallized, and contains a variety of
leach-resistant host phases for the radionuclides.
Celsian glass-ceramic as a waste form has been studied
for potential radiation effects both at the Pacific
Northwest Laboratory (Turcotte et al., 1982) in
cooperation with the Hahn-Meitner Institut, and as
part of a collaborative research program within the
European Community (Malow et al., 1980). In the
study at PNL, the effects of radiation damage in
244Cm-doped celsian glass-ceramic were investigated.
The density was observed to decrease exponentially to
saturation values at a dose of 0.9 x 1018 a-decay
events g-l. As shown in Fig. 17, the volume expansion
at saturation was 0.5%. Stored energy also increased
exponentially with dose to a value of 80 J g- ‘. Data
from the X-ray diffraction analysis revealed that a Cmrich,
rare-earth titanate phase with the pyrochlore
structure underwent a small volume expansion with
dose and eventually became X-ray diffraction amorphous
at the same saturation dose of 0.9 x 1018
u-decay events g- ‘. In the study by Malow et al.
(1980), the celsian glass ceramic was doped with
2.5 wt% 238Pu. The volume also expanded exponentially
with dose, as shown in Fig. 17. Saturation is
predicted at a dose of 1.6 x lOI8 cc-decay events g- ‘,
with a corresponding volume expansion of 0.5%, in
good agreement with the PNL results. The stored
energy was measured to be 43 f 5 J g 1after a dose of
1.l x 1Ol9 cc-decay events g-l. This is approximately
half the value determined at PNL, but it is not known
whether saturation was reached in this material.
Changes in leach rates were measured for specimens
stored at 23 and 170°C to a cumulative dose of
1.1 x 10’8 m-decay events g- ‘. The leach rate decreased
by 4% in the specimen stored at 170°C. The
fractional helium release was also determined to be 3%
after storage at 170°C to a cumulative dose of
1.1 x 1018 cc-decay events g-l.
There have been a number of studies of individual
synthetic phases and minerals that are structurally
and/or chemically analogous to phases found in
multiphase ceramic waste forms. These studies contri-
Radiation effects in nuclear waste forms
bute significantly to the understanding of radiation
effects in the multiphase ceramic waste forms by
providing a detailed understanding of the behavior of
each component phase. The effects of radiation on the
most important phases are summarized. Two of the
phases to be discussed, pyrochlore and apatite, have
also been observed as crystalline phases in nuclear
waste glasses.
p 1.0
.*E
0.8
a,
z 0.6
i?
x9 0.4
.-
5
2 0.2
5.2.1. Pyrochlores. Pyrochlore (Fd3m, Z= 8),
““‘A “‘B ‘vX,‘vY is a derivative of the fluorite structure
‘type’ (Chakoumakos 1984, 1986) in which the
A-site contains large cations (Na, Ca, U, Th, Y and
lanthanides) and the B-site consists of smaller, higher
valence cations (Nb, Ta, Ti, Zr, Fe3+). The essential
feature of the structure is sheets of corner-sharing, BX6
octahedra parallel to the (111) plane which are
arranged into three- and six-membered rings (the
hexagonal tungsten bronze structure). Actinides may
be accommodated in the A-site, and charge balance is
maintained by cation deficiencies in the A-site and
substitutions on the B-site. Rare-earth titanates with
the pyrochlore structure have been observed as
actinide-host phases in nuclear waste glasses (Weber er
al., 1979), in titanate ceramic waste forms (Dosch et
al., 1984) and in a glass ceramic waste form (Malow et
a/., 1980; Turcotte et al., 1982). In addition, the rareearth
titanates with the pyrochlore structure are
chemically and structurally related to zirconolite, an
important actinide host phase in Synroc and titanate
ceramic waste forms.
0.0
0 5 10 15 20
Dose, 1O’salpha decays/g
Fig. 18. Relative diffracted X-ray intensity as a function of
cumulative dose for Cm-doped Gd,Ti,O, (Weber et al.,
1986) and for natural minerals of the pyrochlore group
(Lumpkin and Ewing, 1988).
In Cm-doped nuclear waste glass (Weber and
Roberts, 1983) crystallites of the pyrochlore
(Gd, Cm),Ti,O, transformed from a crystalline to an
amorphous state as a result of a-decay of the
incorporated Cm. The volume expansion associated
with the amorphous transformation of this pyrochlore
phase and an apatite phase resulted in microfracturing
of the simulated waste glass. Similarly, in a study of a
Cm-doped celsian glass ceramic (Turcotte et al., 1982),
the originally crystalline pyrochlore phase,
(Nd,.&m 0.tS)2(TiI.,5Zr0.35 )O,, transformed to the
amorphous state due to self-radiation damage from
z-decay. The unit cell volume was estimated to expand
0.8% before amorphization was complete. Dosch et al.
(1985) reported that the pyrochlore phase in a titanate
ceramic waste transformed to an amorphous state as a
result of irradiation with Pb+ ions. More specific
details of the amorphization process could not be
obtained from the study of these multiphase waste
forms.
investigated (Weber et al., 1985b, 1986). The decrease
in diffracted X-ray intensity with increasing dose is
shown in Fig. 18 for Cm-doped Gd,Ti,O, (Weber et
al., 1986) and for minerals of the pyrochlore group
(Lumpkin and Ewing, 1988). The progression of the
amorphization process in Cm-doped GdzTi,O, was
also followed by transmission electron microscopy
and selected area electron diffraction (Fig. 19). At a
low dose of 0.6 x 10’s x-decays g-r the material
exhibits a strongly crystalline electron diffraction
pattern and clearly resolvable individual defects and
fission tracks. At increasing dose levels, the density of
radiation-induced defects increases to the point that
individual defect clusters are no longer readily distinguished
and a radial ring of diffuse intensity
associated with the presence of amorphous material is
observed in the electron diffraction pattern. At higher
doses, the material is nearly fully amorphous with
some residual crystallinity evident in the electron
diffraction pattern. At a dose of 3.1 x lo’* a-decay
events g - l, the material becomes fully amorphous
with no evidence of any residual crystallinity.
The macroscopic swelling, AVJV,, of Cm-doped
Gd,Ti,O, is shown in Fig. 20 as a function of dose, D,
and follows an exponential relationship given by the
expression:
(2)
where A, is the swelling at saturation, and B,,, is the
amount of material damaged per a-decay event. The
values of A, and B,,,have been determined to be 5.1%
and 6.8 x 10-r’ g, respectively (Weber et al., 1985b,
1986).
The effect of a-decay on the pyrochlore phase Wald and Weber (1984) studied the radiationGd,Ti,O,
doped with 244Cm has been extensively induced changes in dissolution kinetics of Cm-doped
R. C. Ewing et al.
Fig. 19. Microstructural changes in Cm-doped Gd,Ti,O, at (a) 0.6x 10’s u-decay events g-l, (b) 1.0 x lo’*
events g-l, (c) 2.0 x 10’s a-decay events g-l, and (dI) 3.1 x 1CP8 a-decay events g-’ (Weber et al., 1986).
a-dec:ay
Gd,Ti,O, by testing fully-damaged (amorphized distilled, deionized water at 90°C for 14 days. Weight
specimens and a second set of specimens that had been loss measurements indicated an increase in dissolution
fully recrystallized to the original structure by a 12 hr rate of a factor of 2.5 as a result of the radiationanneal
at 1100°C). Specimens were immersed in induced swelling and amorphization. More accurate
Radiation effects in nuclear waste forms 89
Dose, 1O’* alpha decays/g
Fig. 20. Macroscopic swelling of Cm-doped Gd,Ti,O, as a
function of cumulative dose. After Weber et al. (1985b).
radiochemical analysis indicated a radiation-induced
increase of a factor of 2CL-50in the dissolution rates of
the non-network (TiO, octahedra form the network)
Cm and Pu. The concentration of Ti in solution was
below the detection limit of the inductively coupled
plasma emission spectroscopy system, suggesting
some dissolution resistance of the titania network.
Analysis for non-network Gd was not performed.
These results suggest that the dissolution may occur
incongruently, selectively leaching the non-network
ions. Similar behavior is observed for zirconolite
(Wald and Weber, 1984). The increase in dissolution
rate due to amorphization is probably due to the
increase in the size of the network tunnels as the
structure disorders and expands.
The fracture toughness in Cm-doped Gd,T&O,
increased with cumulative dose to a broad maximum
and then decreased slightly (Fig. 21). This is similar to
the behavior of zirconolite (Clinard et al., 1985a;
Weber et al., 1986) and apatite (Weber and Matzke,
9
2.0
I ” ” I ‘, ” I “,
_E
Gd,Ti,O,
2 I- ,
0 1 2 3 4
Dose, 1O’salpha decays/g
Fig. 21. Fracture toughness of Cm-doped Gd,Ti,O, as a
function of cumulative dose. After Weber et al. (1986).
1986b). The increase in fracture toughness is attributed
to the composite nature of the microstructure. At low
to intermediate doses, the microstructure consists of
amorphous tracks in a crystalline matrix, and this
composite microstructure can inhibit crack propagation
and increase the fracture toughness. As the
amorphous phase becomes the dominant matrix at
high doses with remnant crystallites, the fracture
toughness decreases slightly as some of the internal
stresses are relieved. This is supported by the observations
of zirconolite that suggest a relaxation of
disorder at high doses (Foltyn et al., 1985) and the
analysis of strain accumulation in natural pyrochlores
(Lumpkin and Ewing, 1988).
Isochronal(l2 h) annealing of fully amorphous Cmdoped
Gd,Ti,O, shows a linear recovery of density
with temperature up to the temperature where recrystallization
begins (700°C). Recrystallization results in
a sharp recovery of the density with a peak in the
recovery rate at 750°C. Full recovery of the density and
recrystallization of the original pyrochlore structure
are essentially complete at 850°C. Thermal recovery
studies of natural minerals of the pyrochlore group
indicate that recrystallization is an exothermic reaction
that peaks in the temperature range from 650 to
700°C and releases 12&200 J g-’ of stored energy
(Lumpkin et al., 1986). The temperature range is in
reasonable agreement with the isochronal recovery
behavior of Cm-doped GdzTinO,, and the stored
energy released is similar in magnitude to that released
during recrystallization of zirconolite (Weber et al.,
1986; Foltyn et al., 1985) and apatite (Weber, 1983).
Systematic studies (Lumpkin and Ewing, 1985,
1986, 1988; Lumpkin et al., 1986, 1988) have been
completed of a-decay event damage of natural, isometric
pyrochlore structure types. Minerals of the
pyrochlore group are among the common rare element
accessory minerals occurring in a wide range of
igneous rocks: carbonatites, nepheline syenite pegmatites
and granitic pegmatites. Actinides may be accommodated
in the A-site, and in natural pyrochlores the
substitution of uranium can be extensive (up to
30 wt%). Charge balance is maintained by cation
deficiencies in the A-site and substitutions on the B-site
(typically, Nb, Ta and Ti). Natural pyrochlores are
classified on the basis of their major B-site cations:
microlite (Ta > Nb, Ti < 33 mol%); pyrochlore
(Nb> Ta, Ti < 33 mol%); betafite (Ti> 33 mol%)
(Hogarth, 1977). Most of the detailed work on
radiation damage effects has been completed on
microlite, in which the compositions may reach values
as high as 30 wt% UO, and 5 wt% ThO,. Depending
on the age of the sample, the total calculated dpa may
be as high as 80. Careful analysis of samples from the
90 R. C. Ewing ef al.
same deposit, e.g. the Harding pegmatite in New
Mexico which is 1300 million years old, allows one to
study radiation damage over a range of dpa (0.249) as
a function of the uranium concentration
(0.05-8.6 wt% UO,). This provides a good analogue
for long-term radiation damage effects in an actinidebearing,
crystalline waste form. In the remainder of
this section, we summarize the results of such a study
in which the periodic-to-aperiodic transition was
characterized by X-ray diffraction (XRD) analysis,
high-resolution transmission electron microscopy
(HRTEM) and extended X-ray absorption spectroscopy
(EXAFS) and X-ray absorption near edge
spectroscopy (XANES) as a function of increasing
u-decay event dose.
Figure 22 illustrates the change in the X-ray
diffraction pattern of microlites as a function of
increasing cc-decay event dose. With increasing dose
(#147, dpa=0.2; #153, dpa=49), the intensity of the
diffraction maxima decreases and the peaks broaden,
but not as quickly as in the actinide-doped material
(Fig. 18). The positions of the diffraction maxima do
not shift to lower values of 20 (as in the case of zircon
when the unit cell volume increases) because, although
unit cell expansion with increasing dose and the
accumulation of defects is commonly observed, the
a
I-_(222) Harding Microlites
147 1 A A A A
A 149 h L Iv A A
258 -
288
153
I I I I I I
Betafite 083
Pyrochlore 084
Microlite 291
20 60
I I I I
29.8 30.0
Fig. 22. X-ray powder diffraction patterns of pyrochlore group minerals. (a) Microlites from the Harding pegmatite, Taos
County, New Mexico. 0, (T.= 100 million yr) ranged from 0.014 x 1Or6 a-decay events mgg’ for sample #147 to 3 x 1016
u-decay events mg- 1for sample #153. (b) Peak profiles of the (222) diffraction maximum of the microlites. (c) Powder
patterns of three fully metamict samples. DC= 5 x lOI6 or-decay events mg -t for sample #083, 1 x lOI6 a-decay events mg-’
for sample #084, and 1.5 x lOI6 a-decay events mg - ’for sample #291. After Lumpkin and Ewing (1988).
Radiation effects in nuclear waste forms 91
increasing cl-decay dose in natural pyrochlore is the
result of the substitution of U for Ca and Na. The
contraction of the unit cell due to the chemical
substitution evidently compensates for the expansion
of the unit cell due to radiation damage effects. Volume
expansion has been well-documented in pyrochlores
(Wald and Offerman, 1982) in which the composition
remains constant. Debye-Scherrer patterns of microlite
single crystals show, with increasing dose, loss of
K,, and Ka2 peak splitting (0.5-1.0x 1019 u-decay
events g-l), loss of back reflections (1-5 x 1019
cc-decay events g- ‘), and progressive loss of low angle
peaks until none is discernible (5-15 x 1019 cl-decay
events g- ‘). The Bragg diffraction maxima remain
essentially symmetric throughout the transition from
the crystalline to the metamict state. A detailed
analysis (Karioris et al., 1982) of the decrease in Bragg
intensity can be used to estimate the amount of
material damaged by each cr-recoil (Lumpkin and
Ewing, 1988). The calculated spherical diameter of an
cc-recoil track was estimated to be 4.6 nm, in good
agreement with the value of 5.4 nm estimated by
Clinard et al. (1985b) for Pu-doped zirconolite.
The X-ray line broadening was analysed using
Williamson and Hall’s technique (1953), which provides
information on crystallite size and strain. The
maximum strain reached is 0.003, halfway through the
crystalline-to-metamict transition. Average crystallite
dimensions decrease from approximately 450 nm to
less than 15 nm just prior to the microlite becoming
X-ray diffraction amorphous. This suggests that the
early stage of damage accumulation is characterized
by unit cell expansion and strain caused by isolated
defects (predominantly due to a-particle damage).
During the second stage, strain is reduced as the
overlap of aperiodic regions (a-recoil tracks) increases.
With the increase in the volume of the aperiodic
domains, there is a reduction in the size of the
crystallites that remain in the still crystalline regions.
The slight rotation of these crystallites can contribute
to the reduction in strain.
The results of the high-resolution transmission
electron microscopy confirm the interpretation of the
X-ray diffraction data (Fig. 23). High-resolution TEM
images of highly crystalline samples recorded in the
[ 1lo] orientation show variations in contrast resulting
from the B,X, framework and channels parallel to
[llO]. A typical image is shown in Figs 23(a,b). Most
of the images (e.g. 23(c-g)) were taken in the (111)
orientation, exhibiting 0.6 nm lattice fringes which
probably result from the hexagonal tungsten bronze
layer stacking parallel to (111). Some images also
display 0.3 nm (222) lattice fringes. Effects of increasing
U-content on the microstructure are illustrated by
the series of HRTEM images in Fig. 23. The
U-contents have been converted to a corrected dose
assuming the mean life of an cc-recoil track is of the
order of 100 million years. The U-content and
equivalent corrected dose (a-decay events mg- ‘) for
each image are given in the caption of Fig. 23. Three
stages of damage are evident: (1) At an early stage of
damage (< lOI* a-decay events g- ‘), lattice fringes are
essentially continuous, but small variations in diffraction
contrast suggest the presence of isolated defects.
(2) At an intermediate stage of damage (1018-10’9
cl-decay events g- ‘), lattice-fringe free areas indicate
the increasing presence of aperiodic domains of
a-recoil damage tracks. As the dose increases the
proportion of aperiodic domains increases, and
‘islands’ of crystallinity decrease in size as the tracks
increase in number and begin to overlap. (3) At the
highest doses (> 1019 u-decay events g- ‘), no lattice
fringes are apparent and the material is electrondiffraction
amorphous. Electron diffraction maxima
are absent, having been replaced by a diffuse, X-ray
scattering halo.
Once long-range periodicity is lost, XRD and
HRTEM techniques provide little insight into the
results of further damage; however, spectroscopic
techniques provide a powerful means of examining
changes in the first and second coordination sphere
geometries of the aperiodic material. In a series of
studies, Greegor and colleagues (1984a,b, 1985a,b,
1986, 1987) applied EXAFS and XANES analysis to
partially and fully-metamict members of A,B,O,
oxides (e.g. AB,O,: euxenite, polycrase, priorite and
blomstrandine; A2B,0,: zirconolite; A,_,B,OsYo_l
nH,O; pyrochlore, microlite and betafite). Based on
these results, which are summarized in Ewing et al.
(1988), a good model of the structure of the metamict
state is one in which there is: (1) a slight distortion of
the nearest-neighbor coordination polyhedra; (2) a
decrease in the coordination number and bond length
in the primary coordination polyhedra; (3) a slight
increase in the mean second nearest neighbor distances;
(4) a loss of second nearest neighbor periodicity,
probably due to the rotation of coordination
polyhedra across shared corners and shared edges.
The structure of Ta,O,, with its large cell volume and
complex structure, is intermediate between the highly
symmetric periodic pyrochlore structure and the
aperiodic, metamict state; thus serving as a model for
the aperiodic atomic arrangement of fully-damaged
pyrochlore (Greegor et al., 1987).
Lumpkin and Ewing (1988) have noted that if suites
of pyrochlores (of different ages) are analyzed from
single localities, the calculated dose for the crystallineto-metamict
transition increases with the age of the
92 R. C. Ewing et al.
Fig. 23. HRTEM images of microhtes from the Harding pegmatite arranged in order of increasing a-decay dose. (a) Sample
#147, [llO] image, UO,=O.O5, D,=O.O14x 1Or6 a-decay events mg-r.
D,=O.l x 1Or6a-decay events mg-‘.
(b) Sample #149, [llO] image, UO,=O.9,
(c-e) Sample #149, (111) lattice images. UO,=O.9-1.5, D,=O.l4I).4x 1Or6a-decay
events mg-r. (f-g) Sample #288, (111) orientation. UO,=2.5-2.9, D,=O.74.8 x 1Or6a-decay events mgg ‘. (h) Sample
#288, fully-metamict, UO, = 3.6, DC= 1.2 x lOi a-decay events mg- I. Arrows in (b) and (c)indicate possible a-recoil tracks
(Lumpkin and Ewing, 1988).
Radiation effects in nuclear waste forms 93
geologic deposit (Fig. 24). This is clear evidence for the
annealing of a-recoil damage over geologic time. The
relationship between critical amorphization dose and
geologic age was modeled assuming ‘fading’ of the
a-recoil ‘tracks’ in a fashion similar to that used to
describe fission fragment track fading. Critical parameters
are the cc-decay event dose, the age of the
sample, the thermal history, and the mean life of an
a-recoil track. The variation of dose as a function of
age for natural pyrochlores can then be corrected by
the consideration of recoil track annealing. Annealing
under ambient conditions over long periods results in
a higher apparent dose being required for amorphization.
The actual dose versus time curves (Fig. 24) can
be modeled by an exponential equation which considers
the annealing to be described by the half-life of a
recoil-track, ra, in the material (Lumpkin and Ewing,
1988). Figure 25 illustrates such a correction when the
mean life of a track, z,, is assumed to be 100 million
years. A similar analysis of data for zircon and
zirconolite gave mean ‘track’ lives of 400 million years
and 700 million years, respectively. This is in contrast
to the mean lives determined by Eyal and coworkers
using a differential etching technique to determine the
mean life of damaged areas in monazite, UO,,
metamict betafite, thorite and metamict samarskite
METAMCT
0
: TRANSITION ZONE
P
t (109 yr)
Fig. 24. Variation of dose as a function of age for natural
pyrochlores. The upper curve and data represent the
saturation dose estimated from slightly crystalline and
metamict samples. The lower curve and data represent the
onset of detectable a-decay damage defined by an XRD
criterion of I/Z,=O.S-1.0. For a detailed explanation of the
samples used and the symbols see Lumpkin and Ewing
(1988).
(Eyal and Kaufman, 1982; Eyal and Fleischer, 1985;
Eyal et al., 1985, 1987; Lumpkin et al., 1988). For
phases which retain their crystallinity despite high
a-decay event doses (> 10 dpa), the mean life of an
E-recoil track was of the order of lo4 yr. Thus, the
natural samples provide clear evidence for annealing
over geologic time under ambient conditions. The final
state, crystalline or metamict, of the mineral depends
on the mean life of the cc-recoil track, a material
dependent property.
One of the important issues with any waste form is
the amount of stored energy associated with the
radiation damage. For natural pyrochlores thermal
recrystallization effects (measurement of the heat of
recrystallization and identification of phases formed)
have been determined on metamict members of the
pyrochlore group which have received cc-decay event
doses of up to 4 x 10” a-decay events g- ‘. The heats
of recrystallization range from 125 to 210 J gg ‘. The
temperature of the exotherm of recrystallization varies
between 400 and 700°C; the exact temperature is
determined by the sample composition (i.e. 40&6OOC
for pyrochlores and microlites; 65&700°C for betafites).
Release of energy decreases as a function of the
crystallinity (estimated on the basis of the intensity of
X-ray diffraction maxima), with the fully-metamict
samples approaching 200 J g-‘. Lower measured
values (4&125 J g- ‘) are the result of alteration of the
pyrochlores. Other metamict, complex oxides with
stoichiometries of ABO, and AB,O, have lower heats
of recrystallization (4G35 J g-l) and are easily distinguished
from pyrochlore group minerals. The
recrystallization energies for the pyrochlore group
minerals are consistently higher than those measured
in synthetic, Pu-doped zirconolite (50 J g- ’at doses of
3 x 1Ol6 a-decay events mg-’ = 15 x 10z6 a-decay
events mw3), synthetic Cm-doped zirconolite
(127 J g-r at doses of 5 x lOi a-decay events me3),
and natural zirconolite (4&50 J g at doses of 2 x lO”j
a-decay events mg- ’= 1.Ox 1O26 cc-decay events
m- 3). Activation energies of recrystallization range
between values of 0.29 and 0.97 eV, less than those
measured for Pu-doped zirconolites (Foltyn et nl.,
1985; Weber er al., 1986).
5.2.2. Zirconolite. The waste form phase zirconolite
(monoclinic CaZrTi,O,) is one of the three principal
phases of Synroc, and zirconolite is one of the most
extensively studied phases, as it is the primary actinide
host. Monoclinic zirconolite is a fluorite-derivative
structure closely related to pyrochlore. The (001)
layers of the corner-sharing arrangement of TiO,
octahedra are similar to the (111) layers in pyrochlore
94 R. C. Ewing et al.
I:
TRANSITION ZONE B
0
- 0 ‘c m
CRYSTAUfNE
I I ll1lll~ I I 111111~ I I IllIll
d.1 i
t(i09/yr)
Fig. 25. Variation of dose as a function of age for natural pyrochlores after correcting for annealing of cc-decay damage using a
mean track life, r,, of 100 million yr. The upper line is the saturation dose and the lower line is the onset of a-decay damage
detectable by XRD. The variable ‘corrected displacements per atom’ (cdpa) is equal to the dpa corrected for annealing
(Lumpkin and Ewing, 1988).
and the (001) layers in the hexagonal tungsten bronzes.
The A-site cations (Ca, Zr and actinides) are ‘sandwiched’
between the (001) sheets of B-site cations
(corner-sharing TiO, octahedra). There can be several
polytypes depending on the stacking arrangements of
the sheets. In this section, we discuss the damage
response as investigated by study of synthetic compositions
doped with the short half-life actinides, 23sPu
(87.7 yr, a-decay to 234U) and 244Cm (18.1 yr, cc-decay
to 240Pu), natural zirconolites damaged by cc-decay of
uranium and thorium, and finally ion-beam irradiated
zirconolite.
5.2.2.1. 23*Pu-substituted zirconolite. A zirconolite
(CaPuTi,O,) composition can be synthesized with the
substitution of PuO, for ZrO,. X-ray absorption
spectroscopy demonstrated that the 238Pu was incorporated
into the A-site (Greegor et al., 1988). The
effect of this chemical substitution is to convert the
zirconolite composition to an isometric, pyrochlore
derivative structure (Clinard et al., 1984). A major
change in CaPuTi,O, with increasing dose is the large
decrease in density (Fig. 26). Storage at ambient
temperature (approximately 350 K, due to self-heating)
resulted in significant swelling, up to a saturation
value of approximately 5.5 vol.% in material that had
reached an a-decay dose of 2-3 x 1O25 a-decay
events m -3 (for 200 days of storage time). Storage at
575 K results in less swelling, which required a longer
time (300 days) to reach saturation. When this
material was held at 875 K, swelling was minimal.
Transmission electron microscopy conducted after
various storage times and temperatures (Clinard et al.,
1984) showed that with sufficient E-decay damage, Puzirconolite
at 350 and 575 K converts from the
crystalline to the aperiodic, metamict state. Shortly
after the test material was fabricated, examination by
transmission electron microscopy of the Pu-zirconolite
at ambient temperature showed a typical crystalline
diffraction pattern (Fig. 27). At intermediate
states of damage, the microstructure consisted of a
mixture ofcrystalline and aperiodic domains (Fig. 28).
The observed microstructural heterogeneities are consistent
with a random distribution of spherical damage
tracks 5 nm in diameter, a microstructure that can be
predicted from the shape of the corresponding swelling
curve (Clinard, 1986). At high damage levels, electron
diffraction patterns (Fig. 29) show full conversion
from crystalline materials with sharp diffraction maxima
to an inner correlation ring characteristic of the
nearest neighbor interatomic spacing in an aperiodic
structure.
Microhardness indentation studies indicated that as
damage accumulated the zirconolite became softer
(Fig. 33), yet fracture toughness increased (Fig. 34)
(Clinard et al., 1985a). The first of these changes is
consistent with conversion to the metamict state, while
Radiation effects in nuclear waste forms 95
ALPHA DECAY DOSE ( x IO= a/m3 )
p- .
5 .
d2- . a’
2 . .***
I-...”
-1: . . . . . _. . . . . . . . . . . . . . .L.. . . . . 075 K
Y I
OO
I I 1 1 I I _
loo 300
STOR:: TIME (days)
400
IO3 IO4 IO5 5 x IO5
SYNROC AGE (years)
IO6
Fig. 26. Swelling of CaPuTi,O, as a function of storage time, temperature and cc-decay dose. Equivalent age for the
SYNROC-C waste form containing 10 wt% commercial waste calcine is also shown (Clinard et al., 1984).
Fig. 27. Bright-field image and corresponding diffraction pattern from Pu-substituted zirconolite after 4 days storage near
room temperature. The diffuse rings in the diffraction pattern are from the carbon substrate (Chnard et al., 1984).
the latter was interpreted as resulting from the was densification up to a value of 4 vol.% over a
toughening effect of the two-phase microstructure (i.e. period of 400 days. Transmission electron microscopy
mixed periodic and aperiodic domains). confirmed that this recovery corresponded to recrysWhen
the Pu-zirconolite which had been damaged tallization from the metamict state (Clinard et al.,
at ambient temperature was heated to 875 K, the result 1985b). Differential scanning calorimetry measure-
96 R. C. Ewing et al.
Fig. 28. Dark-field image and corresponding diffraction pattern from Pu-substituted zirconolite after 100 days storage near
ambient temperature (Clinard et al., 1984).
1 0.4.m
J
Fig. 29. Bright-field image and corresponding diffraction pattern from Pu-substituted zirconolite after 460 days near ambient
temperature. Diffraction spots representative of crystalline material are no longer present (Clinard et al., 1984).
ments (Foltyn et al., 1985) showed that, as expected,
recovery of metamict material to the crystalline state
was an exothermic process. The amount of energy
released initially increased with damage dose, but then
peaked and subsequently decreased to about half the
maximum value (Fig. 30). This suggests that redamage
of previously-damaged material results in a reordering
of the damage state to one of lower energy.
Modeling calculations (Fig. 31) were completed to
determine the fraction of material that was undamaged,
damaged by a single-track, and damaged by
multiple overlap of tracks as a function of irradiation
dose (Coghlan and Clinard, 1988). Comparison of
Fig. 31 with Fig. 26 shows that swelling occurs under a
constantly changing damage condition, not just with
respect to the proportion of damaged and undamaged
Radiation effects in nuclear waste forms 91
ALPHA DECAY DOSE (x102%.h3)
5 IO 15
I I I 1
& 25
0:
: o,0 400 800 1200
STORAGE TIME (days)
Fig. 30. Energy release from Pu-substituted zirconolite as a
function of aging time and corresponding a-decay dose
(Foltyn et al., 1985).
material, but also with respect to the extent of
redamaged material.
A comparison of stored energy and its corresponding
damage state at different damage levels is shown in
Table 2 (Coghlan and Clinard, 1988). The greatest
amount of stored energy, 100 J g-i, was measured in
material that contained the largest fraction of singletrack
damaged material. At higher doses stored energy
decreased (Table 2), although both the amount of
redamaged material and the total amount damaged
have increased. From these results, Coghlan and
Clinard (1988) determined that the stored energy for
the single track damaged condition is of the order of
200 J g-l. Probable sources of the energy which is
released on annealing of irradiation-damaged
CaPuTi,O, are: (1) the atomic disorder that is
induced; (2) the accompanying strain that results from
formation of aperiodic domains. An estimate was
made of the maximum contribution from strain, 18%,
(Klemens et al., 1987); therefore, by difference, the
decrease in stored energy can be primarily attributed
to atomic-scale rearrangement of the aperiodic structure.
The width of the inner correlation ring observed
by TEM is another indication of the extent of disorder.
Measurements of CaPuTi,O, (Clinard, 1986) showed
that this parameter varied with history of the test
sample (Fig. 32). The greatest width of the inner
correlation ring (diffuse scattering halo) resulted from
material that was held for 460 days at 350 K. The
extent of disorder decreased with irradiation; thus, it is
not surprising that holding the material for 957 days
caused a narrowing or sharpening of the correlation
ring. Further sharpening was observed when the
material was stored at 575 K for 430 days, or reheated
to 875 K for a few days after being self-irradiated to the
fully-redamaged aperiodic structure (625 days at
350 K). Thus, measurements of the inner correlation
ring are consistent with the understanding of damage
and redamage processes obtained by other techniques.
EXAFSjXANES studies completed on 238Pusubstituted
zirconolite held at 350 K for at least 500
days helped to identify details of the damage state
(Greegor et al., 1988). For comparison, measurements
were made on essentially undamaged material (that is,
substituted with 23gPu with a half-life of 24,100 yr as
compared to that of 238Pu of 87.7 yr). In the
heavily-damaged material, the Pu/O nearest neighbor
bond length contracted by 0.003 nm, while expansion
was observed for the more distant coordinating atoms,
corresponding to a 67% overall volume expansion.
This observation is close to the measured volume
swelling of 5.5% (Fig. 26).
5.2.2.2. 244Cm-doped zirconolite. Self-damage
studies of zirconolite have also been carried out by
doping with 244Cm, an a-active isotope with a half-life
of 18.1 yr. Weber et al. (1986) doped CaZrTi,O, with
3 wt% 244Cm and assessed changes in volume,
microstructure, hardness, fracture toughness, leachability
and recrystallization behavior. The Cm-doped
zirconolite retained the monoclinic structure; whereas,
the 23sPu-doped zirconolite was isometric (although
up to concentrations of 5 mol% PuO,, the monoclinic,
zirconolite structure was retained (Clinard et al.,
1984)). More importantly, the composition of the Cmdoped
zirconolite was not greatly changed by the
substitution; whereas, the Pu-substituted structure
was significantly changed by replacement of ZrO, by
PUO,.
Swelling data at 350 K were obtained to nearsaturation,
and by use of an exponential relationship, a
saturation swelling value of 6.0 vol.% was calculated
for a dose of 2 x 10z5 E-decay events rnm3. These
results are in good agreement with those for the Pusubstituted
material, which exhibited saturation swelling
of 5.5 vol.% at nearly the same damage level
(Clinard et al., 1984). TEM images and diffraction
patterns confirmed that Cm-doped zirconolite
amorphized by the accumulation of damage tracks
produced by the 240Pu recoil ions that are emitted
during cc-decay of 244Cm. A smaller amount of
additional damage resulted from spontaneous fission
of 244Cm, which introduced a low density of fission
tracks.
Changes in hardness and fracture toughness as a
function of increasing a-decay event dose are shown in
Figs 33 and 34 combined with the results for Pu-doped
zirconolite. Both materials showed a decrease in
hardness and an increase in fracture toughness; the
98
Fig. 31. Calculated fraction
0.8
0.6
R. C. Ewing et al.
ALPHA DECAY DOSE (~10~~ a/m3)
0 2 4 6 8 10
I I I I I
MORE TIMES
DAMAGED TWO OR
MORE TIMES
DAMAGEDONCE
200 400 600
STORAGE TIME (days)
of undamaged, once-damaged, and multiply-damaged Pu-substituted zirconolite as a function of
storage time and damage dose (Coghlan and Clinard, 1988).
former change is expected as a result of the radiationinduced
transformation from a periodic to an aperiodic
state, while the latter may result from the
heterogeneous microstructure (Clinard et al., 1985a).
Differential thermal analysis of the Cm-doped
zirconolite after self-irradiation to 2.5 x 10z5 a-decay
events m 3 showed a strong exothermic reaction
during recrystallization, with a sharp recovery stage of
114 J g- ’at 68&73o”C. These values of stored energy
and recovery temperature agree closely with those
obtained from Pu-substituted zirconolite, although
the latter material showed a halving of stored energy at
higher radiation doses as a result of damage beyond
the dose required for amorphization (Foltyn et al.,
1985). Measurements on the Cm-doped material
showed that the temperature of the energy release peak
increased with heating rate; this behavior allowed the
calculation of an activation energy of recrystallation of
5.8 eV.
The dissolution (leaching) characteristics of the
Cm-doped zirconolite were determined (Weber et al.,
1986) by testing fully-damaged (amorphized) specimens
and a second set of specimens that had been fully
recrystallized to the original structure by a 12 hr
anneal at 1100°C. Samples were immersed in distilled,
deionized water at 90°C for 14 days. The results,
shown in Table 3, indicate a notable increase in leach
rate with amorphization; especially significant is the
factor of 10 increase for the x-decay product Pu. The
high final pH is believed to limit the solubility of the
Cm to the value shown. The low concentrations of Ti
as compared to that of Ca suggest that the dissolution
is incongruent, with the Ti ions (which are part of the
TiO, octahedra comprising the network structure)
being more resistant to leaching than the interlayer
cations (e.g. Pu and Ca). The mechanism for the
enhanced leaching of the radiation-induced amorphous
state may be associated with the broken bonds
and the increased free volume (a more open framework)
of the amorphous state.
5.2.2.3. Ion beam-irradiated zirconolite. Radiation
damage effects in zirconolite have also been investigated
by bombardment with energetic heavy ions
(Weber et al., 1992). TEM of samples irradiated with
3 MeV Ar+ ions showed damage build-up similar to
that observed in material damaged by a-decay of
238Pu or 244Cm, namely, accumulation of small
defects (alpha-recoil tracks) to the point where the
material became globally disordered. Recent studies
(White et al., 1994) of the temperature dependence of
amorphization in zirconolite have shown an increase
in amorphization dose with temperature, in general
agreement with the results of Clinard et al. (1984) for
23*Pu-substituted zirconolite.
Natural zirconolites have also been used in ionbeam
irradiation simulations of a-decay event-induced
amorphization (Ewing and Wang, 1992). Recrystallized
samples (1130°C for 8 hr in air) were irradiated
with 1.5 MeV Kr+ ions using the HVEM-Tandem
Facility at Argonne National Laboratory. The Kr+
Radiation effects in nuclear waste forms 99
WIDTH OF INNER
CORRELATION RING (PEAK)
FOR METAMICT Ca PuTi,O,
FULL WIDTH
AT HALF-MAX.
I 1 i I I
350 K FOR 350 KFOR 575 K FOR I
460 DAYS 957 DAYS 430 DAYS 350 FOR’626 DAYS,
THEN 675 K FOR:
1.7 DAYS 6 DAYS
Fig. 32. Widths of inner correlation rings from electron diffraction patterns of metamict CaPuTi,O, after conversion to
intensity peaks. Values were determined at half-maximum and 10% of full height for five damage conditions (Clinard, 1986).
dose rate during the irradiation was 3.4 x 10” ions
cm - ’s - I, which is a damage rate 2 x 10” times higher
than that which has occurred in the natural zirconolite
due to decay of uranium. In situ transmission electron
microscopy was completed during the ion irradiation.
After a Kr+ doseof x 1014ionscm~z (=0.3 dpa), the
zirconolite grains were amorphized. The dpa level for
ion beam induced amorphization is six times lower
than that calculated for the natural a-decay induced
process. Some of the a-decay event damage may have
recovered by point defect migration over geologic
periods of time. It is interesting to note that inclusions
of thorianite, ThO,, in the zirconolite did not become
amorphous at equivalent doses. This is consistent with
the already mentioned apparent radiation resistance of
UO, . The resistance of thorianite to amorphization
compared to zirconolite is consistent with a mode1 of
damage accumulation in insulators in which the
critical dose for amorphization is inversely proportional
to the topologic and chemical complexity of the
phase (Wang et al., 1991). Thorianite is isometric with
only one unique cation site. Zirconolite is monoclinic
with three unique cation sites. Such criteria may be
useful in evaluating the response of waste form
ceramics to a-decay event damage.
Dose-dependence of damage in zirconolite has been
studied by Headley er al. (1982) using Pb+ ions of
4&240 keV at ambient temperature. Damage pro-
100 R. C. Ewing et al.
10
--N .
\
9- \
\
\ CaPuTiZ07
\
J
m 6-
\
\
&
\
.
I
$ 7-L.. ==__
E
u
b 6
I
CazrTi,077-Q--Q
5-
4-
3-““““““““““”
0 5 10 15 20
Cumulative Dose, 1O24 Alpha Decays/m3
Fig. 33. Hardness of Cm-doped CaZrTi,O, and Pu-substituted zirconolite as a function of cumulative dose
(Weber et al., 1986).
0.5
c
CaPuTiP07
011” 8 ” ” ” ” ” ” ” ” ”
0 5 10 15 20
Cumulative Dose, 1O24 Alpha Decays/m3
Fig. 34. Fracture toughness of Cm-doped CaZrTi,O, and Pu-substituted zirconolite as a function of cumulative dose
(Weber et al., 1986).
Radiation effects in nuclear waste forms 101
Table 2. Damage level, stored energy and corresponding damage fractions for CaPuTi,O,
self-damaged near 300 K (Coughlan and Clinard, 1988)
Damage level Stored energy
(days) (kJ kg-‘)
Amount Amount Amount
undamaged damaged once redamaged
(“/) (“/) (“/)
154 100 18 32 50
224 92 10 23 61
558 59 0 2 98
1198 51 0 0 100
Table 3. Summary of 9O”C,14 days leach test data (Weber et al., 1986)
Material State
Normalized elemental mass loss (g mm2)
Ca Ti Cm Pu
Percent
weight
loss
Final
PH
CaZrTi,O, Crystalline 0.72 to.02 0.001 0.02 0.02 6.4
CaZrTi,O, Amorphous 5.99 10.02 0.001 0.22 0.05 7.3
ceeded in four stages: isolated tracks, track overlap,
domination by aperiodic domains, and finally complete
amorphization. Results were compared with
those for natural zirconohte self-irradiated to various
damage levels by a-decay over geologic times. Both the
nature and extent ofdamage were similar, regardless of
the large difference in damage rates. Later, Clinard et
al. (1984) extended the comparison to include
23sPu-substituted zirconolite and showed that
damage at the intermediate rate characteristic of
zirconolite was similar to that observed in natural and
Pb+ ion-bombarded material.
5.2.2.4. Natural zirconolite. Although zirconolite
occurrences are rare, selected specimens have been the
subject of extensive studies (Ewing et al., 1982a; Ewing
and Headley, 1983; Ewing, 1983; Oversby and
Ringwood, 1981; Sinclair and Ringwood, 1981; Sinclair
and Eggleton, 1982). Samples from Sri Lanka
(550 million yr old) typically can have concentrations
of UO, of 2 wt% and of ThO, of 20 wt%, thus
reaching doses of 1016 a-decay events mgg ’ (= 1O26
a-decay events mm3). This is a dose equivalent to
nearly two displacements per atom.
X-ray diffraction analysis of fully-metamict zirconolite
shows no diffraction maxima confirming that
zirconolite is X-ray diffraction amorphous. Because of
the paucity of samples which span a range of a-decay
event doses, it is not possible to follow the change of
diffraction maxima intensity and position with
increasing a-decay event dose. A zirconolite annealed
at 1130°C did give a diffraction pattern with sharp
peaks which could be refined on the basis of a
monoclinic unit cell (Lumpkin et al., 1986b). Referring
to the work of Wald and Offerman (1982) on a
244Cm-doped zirconolite, the expansion of the unit
cell was anisotropic with the greatest expansion
perpendicular to the sheets of TiO, octahedra
(Fig. 35) reaching a maximum value of 2.1% volume
expansion at a dose of 1O23a-decay events me3.
High-resolution transmission electron microscopy
showed no evidence for crystallinity (Lumpkin et al.,
1986b), and in approximately 10% of the grains
examined there were subspherical microvoids ranging
in diameter from 100 to 400 nm (Ewing and Headley,
1983). The microvoids are attributed to He-gas
accumulation which is a result of the radioactive
decay. Annealed samples (1100°C) recrystallized producing
crystals up to 2000 nm in length. The crystallites
were often twinned at the unit cell scale as a result
of different stacking sequences of the TiO, sheets, thus
resulting in laminar intergrowths of the various
zirconolite polytypes (Lumpkin et al., 1986b). The
same texture of fine-scale twinning has been observed
in synthetic zirconolite in Synroc C (White et al., 1984)
and synthetic zirconolite of Ca,$m,,,ZrNb,.,
Mg,,,Ti,.,O, composition (White, 1984). The defects
along the twin boundaries may become sites at which
actinide elements are incorporated. In some instances
(Lumpkin et al., 1986b), there was evidence for the
formation of crystals (55100 nm) with d-spacings
similar to those of the fluorite structure type. Lumpkin
et al. (1986b) suggested that metamict zirconolite may
recrystallize initially with a disordered, fluorite-type
structure. Continued heating to higher temperatures
(100&12Oo”C) appears to favor the highly twinned,
monoclinic zirconolite structure. Differential thermal
analysis (3G12OO”C) showed exothermic recrystallization
at 780°C with a heat of recrystallization of
102 R. C. Ewing et al.
Fig. 35. Polyhedral representation of zirconolite (CaZrTi,O,) crystal structure. (Upper) (010) projection (four unit cells are
outlined) with an ellipsoid superimposed to illustrate the relative expansion of the structure due to a dose of 1O25 a-decay
events me3. The octahedra represent TiO,; the small spheres, Zr; the large spheres, Ca. (Lower) (001) projection (four unit
cells shown as in upper diagram). The shading is darkest for the top layer, and lightens with increasing depth into the structure.
No physical significance should be attached to the sphere sizes used. (Lumpkin et al., 1986.)
Radiation effects in nuclear waste forms 103
50 J g-’ for a zirconolite with a total dose of > 10z6
cc-decay events rne3. This is similar to the value
obtained by Foltyn et al. (1985) for synthetic
CaPuTi,O, at nearly the same a-decay event dose of
10z6 cc-decay events me3 after 1200 days. Although the
measured energy release for the natural zirconolites
and Pu-doped zirconolites of equivalent dose are
nearly identical, both of these values are distinctly
lower than the value (84 J g-‘) measured for Pudoped
zirconolite that had experienced a smaller dose
(approximately 10z5 a-decay events mm3) (Foltyn,
1985) or for Cm-doped zirconolite (126 J g-‘) at a
dose of 2.0 x 10” cc-decay events mm3 (Weber et al.,
1985a, 1986). Foltyn et al. (1985) have suggested that
the decrease in released energy with increasing dose is
due to a redamaging process that results in a
reordering of atoms in the metamict state.
Detailed EXAFS and XANES studies have been
completed on fully-aperiodic zirconolite (Lumpkin et
al., 1986b; Greegor et al., 1984a; Farges et al., 1993).
These studies, in many cases on the same sample from
Sri Lanka (#l 11.35),were on the Ca, Th, U, Zr and
Ti EXAFS spectra, that is all of the major cations of
zirconolite. Principal conclusions included: (1) The
loss of long range atomic periodicity occurs by
changes in the metal-oxygen-metal (M-O-M) angles
across mainly edge-sharing coordination polyhedra.
The tilting of nearest neighbor coordination polyhedra
occurs without significant changes in the first
and second neighbor distances. Thus, the secondneighbor
coordination geometries for Ca, Ti and Zr
may not be very different in metamict and crystalline
zirconolite, despite the loss of long-range periodicity.
This is illustrated in Fig. 36 which shows that a slight
change (5”) in the mean Zr-O-M angle correlates
with a relatively large variation (0.01 nm) in the
Zr-M distance. Thus, variations in the mean
M-O-M angle do not have to be large in order to
have a significant effect on long-range periodicity and
the medium-range order in this radiation-damaged
phase. (2) Although there is some evidence for
reduction in coordination number and an associated
decrease in bond-length with increasing cc-decay event
dose (Greegor et al., 1984), the short-range (nearest
neighbor oxygen) environments of Zr4+, Th4+, Ca2+
and Ti4+ are little changed by increasing radiation
damage. (3) The structural environments around
Zr4’ and Th4+ in metamict zirconolite (7-coordinated
and g-coordinated, respectively) are clearly
different from those measured in silicate and borosilicate
glasses in which they are commonly
6-coordinated with oxygens. This suggests that there
are fundamental structural differences in cation
environments between radiation-induced aperiodic
(metamict) phases and glasses quenched from
melts (Sales and Boatner, 1990).
5.2.3. Perovskite. Perovskite, CaTiO,, is a mineral
that may assume a wide range of compositions as
stable solid solutions. Perovskite is also a highly
refractory phase that has been stable in near surface
environments for billions of years. Thus, CaTiO, is an
excellent candidate for immobilization of radioactive
waste elements, particularly actinides, and is a major
constituent phase in Synroc and other titanate ceramic
waste forms. Perovskite is orthorhombic and consists
of a three-dimensional network of corner-sharing
TiO, octahedra with Ca occupying the large void
space between the octahedra (the corner-sharing
octahedra are located on the eight corners of a slightly
distorted cube). The actinides, rare earths, and other
large cations readily replace Ca in the structure, while
smaller-sized cations replace titanium.
There are limited data on radiation effects in
perovskite. Sinclair and Ringwood (1981) have studied
natural perovskites of varying age and uranium and
thorium contents. Natural perovskite samples that
have received a dose of up to 2.6 x 1Ol8 u-decay
100 120 140 160 180 200 220 240 MO
Zr-O-cation angle (“)
Fig. 36. Plot of mean Zr-metal distance (A) versus Zr-0-M
angle (“) for polyhedra of M-cations (M=Ti, Zr, Th, Ca)
arranged in the edge-shared configuration in zirconolite. The
stippled region shows the normal range of distance and angle
values in crystalline zirconolite. (Farges et al., 1993.)
104 R. C. Ewing et al.
events g-’ (equivalent to 20,000 years storage for
Synroc) showed a high degree of atomic periodicity
with no evidence of amorphization. The unit-cell
volume expansion at this dose was 1.8% due to the
accumulation of simple structural defects. Loparite is
an isometric perovskite that contains large amounts of
Ce and Nb and is also highly crystalline. The unit-cell
volume expansion is 1.5-1.7% after a dose of l-2 x
10’s cc-decay events g-’ (Ringwood et al., 1980;
Ringwood and Sinclair, 1981).
Cold-pressed and sintered specimens of synthetic
CaTiO, have been neutron irradiated in the fastneutron
flux of the HIFAR reactor at Lucas Heights,
for different lengths of time (Reeve et al., 1981;
Woolfrey et al., 1982). The highest dose reported was
2.5 x 1O26n mm2 (E> 1 MeV), which represents an
equivalent Synroc storage time of 5.8 x lo5 yr and
corresponds to 0.7 dpa or 8 x 10” cc-decayevents g- ‘.
The volume expansion reached a value of 6% at the
maximum dose achieved, but showed no indication of
approaching a saturation value (Fig. 37). X-ray
diffraction analysis indicated sharp diffraction maxima,
suggesting strong resistance to amorphization,
but shifts in specific peaks were not analyzed. Similar
behavior has been observed in Pu-doped CaTiO, but
at lower doses (Fig. 37).
While the above studies indicate that the perovskite
structure is fairly resistant to radiation-induced amorphization,
studies by Karioris et al. (1982) indicate
that CaTiO, readily amorphizes under 3 MeV Ar+ ion
irradiation. Similarly, the perovskite-type compound
CmAlO, has been reported to transform from a
crystalline to a fully amorphous state after a dose of
1.4 x 10’s u-decay events g-l, which was achieved in
7.5 days from the decay of 244Cm (Mosley, 1971). The
01 ’ ’ ’ ’ ’ ’ 1
0 2 4 6 6 10
Dose, 10” alpha decays/g
Fig. 37. Volume expansion of Pu-doped (Evans et al., 1986)
and fast neutron-irradiated (Woolfrey et al., 1982) perovs-
kite.
crystal structure CmAlO, was fully recovered after
annealing at 1000°C for 30 min. Questions regarding
the radiation stability of perovskite are being
addressed in a new study of a-decay damage in Cmdoped
perovskite. Preliminary results indicate a linear
increase in volume with dose (0.75% at 5 x 10”
a-decay events g- ‘), as well as an increased leach rate
(Mitamura et al., 1994).
5.2.4. Barium hollandite. Barium hollandite,
BaAl,Ti,O,,, is a major phase in several different
Synroc formulations (Ringwood et al., 1979, 1981;
Ringwood, 1982; Newkirk et al., 1982). The structure
of barium hollandite consists of a framework of TiO,
octahedra that are linked (edge-sharing and cornersharing)
to form square channels parallel to the c-axis.
These channels can accommodate a wide variety of
large cations, including fission products, such as Cs.
Although hollandite will not contain actinides, it will
experience irradiation from a-particles emitted in
adjacent actinide-containing phases. The cumulative
a-particle fluence incident on barium hollandite as a
function of waste storage time is shown in Fig. 38.
Weber (1985) investigated the effect of a-particle
irradiation on the structure of barium hollandite
utilizing cc-particles emitted isotropically from an
effectively semi-infinite 238Pu02 source. The energy
spectra of the cr-particles emitted from this source
closely approximate that expected to be incident on
the barium hollandite phase in Synroc. Irradiation
with the a-particles resulted in a large expansion of the
unit-cell volume that saturates at 2-2.5% and also
induces a displacive transformation from the tetragonal
structure to a lower symmetry monoclinic
structure. The anisotropic unit-cell expansion is primarily
along the a- and b-axes and causes a small
increase of the angle p from 90” in the tetragonal
structure to 91”in the deformed monoclinic structure.
A similar tetragonal-to-monoclinic transformation
has been observed as a result of heavy-ion and electron
irradiations (Barry et al., 1983). The anisotropy of the
unit-cell expansion caused an increase in the size of the
channels along the c-axis, which could significantly
affect the ability of the barium hollandite structure to
retain channel-site atoms, such as Cs, in an aqueous
environment. The implied increase in dissolution rates
of non-network atoms as a result of radiation-induced
volume expansions has been previously described for
both the pyrochlore and zirconolite structures (Wald
and Weber, 1984) which also consist of a TiO,
octahedra network.
Comparisons of the E-particle irradiation results of
Weber (1985) with the fast-neutron irradiation results
Radiation effects in nuclear waste forms 105
Basis:
0 Commercial HLW
0 3.3% Enriched UO, Fuel
?? 33,000 MWdltonne
?? 25 wt.% Waste Loading
?? 60 wt. % ActinideContaining
Phases
lo’s,
I I I I,,88 I ,,,,I I I ,,,I I1 II
10 102 10s 104
Waste Storage Time, years
Fig. 38. Cumulative a-particle fluence incident on barium hollandite in SYNROC. After Weber (1985).
of Woolfrey et al. (1982) indicate a significant irradiation
source dependence. Both studies were completed
using specimens from the same batch preparation of
hollandite. The volume changes under cr-particle
irradiation occur at an order of magnitude lower dose,
calculated as displacements per atom, than the volume
changes under fast-neutron irradiation (Fig. 39). This
is significant since an order of magnitude in dose
corresponds to two or three orders of magnitude in
time. The a-particle irradiation primarily produces
isolated interstitials and vacancies along the path of
the a-particles that give rise to a large unit-cell
expansion; whereas, the fast-neutron irradiation,
according to Ball and Woolfrey (1983) produces
planar interstitial clusters and isolated vacancies, with
only a small increase in unit-cell volume.
While there was no evidence for amorphization in
the x-irradiated barium hollandite, heavy-ion irradiation
of barium hollandite has been reported to cause
amorphization (Headley et al., 1982; Karioris et al.,
1982). These results clearly demonstrate that the
proper selection of irradiation source is important if an
accurate simulation of waste form behavior is desired.
5.2.5. Zircon. Zircon (ZrSiO,, Z4,/amd, Z=4)
occurs in silica-rich tailored ceramics, and its structure
type is the high-temperature form of the heavy rare
earth phosphates (the low-temperature form has the
monazite structure type). Pu can substitute for Zr, and
radiation damage studies have been completed on
238Pu-doped zircon up to concentrations of 8.1 mol%
238Pu (Weber, 1990, 1991b). Additionally, zircon is a
phase extensively used in U/Pb radiometric agedating,
as it may contain uranium in concentrations up
to 5000 ppm. Early workers (Holland and Gottfried,
1955) investigated radiation damage effects in an effort
to use the degree of damage as a measure of the age of
the zircon. Theirs was the first study to quantitatively
correlate the decrease in density, decrease in refractive
index, and expansion of the unit cell parameters with
increasing a-decay event dose. Based on their study of
zircon, Holland and Gottfried proposed the first
model of defect accumulation and damage in-growth
for a complex ceramic. This early history of the study
of radiation damage effects in metamict zircons is
summarized by Ewing (1994).
There are data on radiation effects in natural zircons
(irradiation time: 570 million yr (Murakami et al.,
1986, 1991; Woodhead et al., 1991a,b)), Pu-doped
zircon (irradiation time: 6.5 yr (Exarhos, 1984; Weber,
1990, 1991b)) and most recently ion-beam irradiated
zircon (irradiation time: 30 min (Babsail et al., 1991;
Wang et al., 1993; Weber et al., 1994)). The damage
rate for the 238Pu-doped zircon is 6 x 10” Bq g- ‘, and
the damage rate for natural zircons is < lo3 Bq g- ‘.
This extraordinary data set with a dose rate range of
106 R. C. Ewing et al.
> 108, allows one to investigate the efficacy of using
accelerated techniques (actinide-doping and ion-beam
irradiations) to simulate the long-term effects of
a-decay event damage.
The study of natural zircons culminated with the
very detailed analysis (Murakami et al., 1991; Woodhead
et al., 1991a) of damage in-growth in suites of
zircons from Sri Lanka (570 million yr old) which had
experienced variable x-decay event doses due to
variable U and Th contents (0.06 x 1015-6.8 x 1015
a-decay events mg- ‘). In addition to the decrease in
density and refractive indices with a-decay event dose
as determined by Holland and Gottfried in 1955, later
workers applied detailed X-ray diffraction analysis
(Murakami et al., 1986, 1991), electron microprobe
analysis (Murakami et al., 1991; Chakoumakos et al.,
1987), high-resolution transmission electron microscopy
(Murakami et al., 1991; McLaren et al., 1994),
infrared spectroscopy (Woodhead et al., 1991b), and
X-ray adsorption spectroscopy (Farges and Calas,
1991; Farges, 1994) to the study of this radiationinduced
transition to the aperiodic (metamict) state.
Most recently, the decreases in harness (40%) and
elastic moduli (60%) over the dose range of
0.15-0.65 dpa were determined in zoned zircons using
a mechanical properties microprobe (Chakoumakos
et al., 1991) and increases in stored energy have been
measured as a function of increasing x-decay event
dose (Ellsworth et al., 1994). The enthalpy of annealing
at room temperature varies sigmoidally as a
function of radiation dose and reaches a saturation
plateau at doses of 5 x 10” cl-decay events rng- i. The
large magnitude of the enthalpy of annealing plateau,
-59 kJ mol-‘, suggests that the damage to the
structure is pervasive on the scale of nanometers.
Based on X-ray diffraction analysis and highresolution
transmission electron microscopy, Murakami
et al. (1991) described three stages (Fig. 40) of
damage in-growth. Stage-I (~3 x lOi cc-decay
events g- ‘) is characterized by sharp Bragg diffraction
maxima with a minor contribution from the diffusescattering
component. Electron diffraction patterns
were sharp. Damage is dominated by the accumulation
of isolated point defects which cause unit-cell
expansion and distortion that account for most of the
decrease in density. These defects may partially anneal
over geologic time. Stage-II (3-8 x lOi cc-decay
events g-i) is characterized by significant decreases in
the intensity of the Bragg diffraction maxima which
become asymmetric due to increased contributions
from the diffuse-scattering component. High-resolution
transmission electron microscopy revealed that
the microstructure consisted of distorted crystalline
regions and amorphous ‘tracks’ caused by a-recoil
nuclei. With increasing a-decay dose, damaged crystalline
regions are converted into aperiodic regions,
but with no further significant expansion of the unit
cell in the remaining crystalline regions. Stage-III( >
Dose, dpa
Fig. 39. Volume expansion in a-irradiated (Weber, 1985) and neutron-irradiated (Woolfrey et al., 1982) barium hollandite.
6
BaAI,Ti,O,, Type B
5Neutron
Irradiation
P 4-
s
‘B
5
:: 3 Alpha Irradiation
w
E
2
2 2
Radiation effects in nuclear waste forms 107
Stage I
L
Stage II
!
Stage III
1
(b)
0 0 0
0 0
0 0
a 0
a 0 0
Structure
-_-A_I
,L
-
219angle -
Electron
diffraction
pattern
X-ray
diffractometer
pattern
Fig. 40. Schematic representation of the change in X-ray diffraction pattern (right column), electron diffraction pattern (center
column), and the structure (left column) with increasing a-decay event dose (top to bottom). After Murakami et al. (1991).
8 x lo’* a-decay events g- ‘) consists of material which
is entirely aperiodic as far as can be determined by
X-ray or electron diffraction. Figure 40 is a schematic
representation of the changes in X-ray diffraction and
electron diffraction pattern that occur at each of the
three stages. These stages of damage are illustrated in
HRTEM micrographs over the range of
0.0025~.50 dpa (Fig. 41).
Finally, there was no HRTEM or XRD evidence for
the formation of ZrO, or SiO, as final products during
the last stage of metamictization: however, recent
work based on microcalorimetry (Ellsworth et al.,
1994), X-ray absorption spectroscopy (Farges, 1994)
and SIMS and electron microscopy (McLaren et al.,
1994) has postulated the presence of zirconium- and
silicon-rich domains which are less than 10 w in size.
These domains have been identified at higher temperatures
by neutron diffraction experiments (Mursic et al.,
1992).
Based on modeled density changes, aperiodic
regions continued to experience a change in structure
as they were redamaged. During Stage-II of the
process, the modeled density of aperiodic regions
changed from 4.5 g crne3 to 4.1 g cmm3. The amorphization
process is consistent with a model for the
multiple overlap of displacement cascades (Carter and
Webb, 1979; Webb and Carter, 1979, 1981), suggesting
amorphization occurs as a result of defect accumu-
108 R. C. Ewing et al.
Fig. 41. HRTEM micrographs of self-radiation damage in natural zircons showing an increased degree of amorphization with
increasing dose: (a) 0.0025 dpa, (b) 0.091 dpa, (c) 0.32 dpa, and (d) 0.50 dpa. The spacing marked in (c)is 0.33 nm [the (200)
d-spacing]. The slight rotation of the remaining crystallites in the amorphous matrix in (c) is caused by local strain fields
induced by amorphization. (Weber et al., 1994.)
lation rather than directly within a single displacement
cascade. Figure 42 compares the fraction of amorphous
material (estimated from density and optical
data) for Pu-doped and natural zircons (solid and
dashed lines, respectively) to the calculated values for a
double overlap model (i.e. the defect density for
amorphization requires the overlap of three cascades
as discussed below). Note the incubation dose
(0.05 dpa) required for the natural zircon which
reflects annealing of the isolated point defects. Comparison
of the results for natural zircon with those for
Pu-doped zircon (Weber, 1991b) showed that doserate
variations (even as great as a factor of lo*) had no
substantial effect on the damage accumulation pro-
cess.
Unit cell parameters and volume increased (Fig. 43)
and density decreased more for the Pu-doped zircon
than for natural zircon in the early stages of damage
accumulation (< 3 x lo’* u-decay events g-r), suggesting
that annealing of point defects in the early
stages of the damage accumulation process occurs in
natural zircon under ambient conditions. Evidence for
annealing under ambient conditions was noted by
Holland and Gottfried (1955) and the mean lifetime of
an a-recoil track was estimated to be 400 million yr by
Lumpkin and Ewing (1988). The annealing accounts
for the distinct sigmoidal shape of the damage curves
for natural zircon and the apparent incubation period
before the onset of amorphization (see Fig. 42). Thus,
radiation damage from cl-decay events in zircon results
in the simultaneous accumulation of simple structural
defects and amorphous (or aperiodic) regions that
exhibit loss of long-range order (Chakoumakos et al.,
1987; Murakami et al., 1991; Farges and Callas, 1991).
The amorphous regions eventually overlap to produce
a completely amorphous state that may be structurally
unique. The local amorphization process is believed to
result from the spontaneous local collapse of the
Radiation effects in nuclear waste forms
Pu-Doped Zircon
Natural Zircon
109
Dose (dpa)
Fig. 42. Amorphous fraction, f,, as a function of dose in Pu-doped and natural zircons. The solid curve for Pu-doped zircon is
based on a fit of the double-overlap model (Webb and Carter, 1979,198l) to the data of Weber (1990). The dashed curve for
natural zircon is based on a fit of the double-overlap model, with an incubation dose (0.05 dpa), to the optical and density data
of Holland and Gottfried (1955). After Weber et al. (1994).
6, I
Dose (dpa)
Fig. 43. Unit cell volume expansion, A V,,/Vo, as a function of dose in Pu-doped and natural zircons. The data of Weber (1990)
for Pu-doped zircon fit a simple model (Nellis, 1977; Weber, 1981) for the accumulation of immobile defects,
AV,,/V, = A{ 1-exp[ -LSD]}. Thedata ofHolland and Gottfried (1955) and Murakami et al. (1991)fit anempirical expression
of the form AV,,/V,-A{1 -exp[-(BD)“]}. After Weber et al. (1994).
crystal structure and disappearance of long-range
order due to a local high-defect concentration; amorphization
is triggered when the free energy of the highdefect
region becomes equal to the free energy of the
amorphous state. The local defect concentration
necessary to trigger amorphization can occur directly
in the displacement cascade (particle track) of a single
heavy particle (recoil nucleus in a-decay) or from the
overlapping cascades of individual light and heavy
particles (a-particle and recoil nucleus).
Several models of both direct and multiple-overlap
processes have been developed to describe amorphization
in ion-irradiated semiconductors (Carter and
Webb, 1979; Webb and Carter, 1979). These models
have been used to gain some insight into the process
for amorphization in zircon due to internal a-decay
110 R. C. Ewing et al.
events (Weber et al., 1994). As shown previously
(Weber, 1990, 1993) amorphization in Pu-doped and
natural zircons is consistent with amorphization
occurring by the catastrophic collapse of local structure
through an increase in local defect concentration
by the double overlap of displacement cascades. The
fraction of amorphous phase, f., is given by the
expression:
f.=l-[(1+RD+R20’/2)exp(-RD)], (3)
for the double-overlap model (Carter and Webb, 1979;
Webb and Carter, 1979), where R (with dpa as the unit
of dose, D) is the average ratio of total atoms to
displaced atoms within the cascades (tracks) produced
by the cc-decay event particles (cc-decay events per unit
volume or mass). The amorphous fractions have been
determined previously for Pu-doped zircon from
macroscopic swelling and unit-cell expansion data
(Weber, 1990) and for natural zircons by Holland and
Gottfried (1955) from both refractive indices and
density data. The results and model fits are shown in
Fig. 42. The amorphous fractions for Pu-doped and
natural zircons exhibit very similar behaviors as
functions of cumulative displacement dose, except for
an offset at low dpa due to the incubation dose in
natural zircon that is caused by the point-defect
annealing over geologic time as discussed above. In the
case of Pu-doped zircon, non-linear regression analysis
yields a value of R=9.51 dpa- ‘. An incubation
dose, D,, ,must be assumed in the case of natural zircon
for the double-overlap model to provide a reasonable
fit to the results of Holland and Gottfried (1955). Nonlinear
regression analysis yields values of R= 11.67
dpa- 1and D, = 0.05 dpa for natural zircon. The value
of R is in good agreement with the value determined
for Pu-doped zircon, and the incubation dose is
consistent with the incubation dose determined from
both the XRD intensity data and the onset for
expansion along the a-axis. Physically, the doubleoverlap
model suggests that some limiting local defect
concentration that corresponds to the overlap of three
displacement cascades is required to trigger amorphization;
this is consistent with the dependence of
amorphous fraction on defect concentration reported
earlier, with the observed effect of defect annealing,
and with the observations of some intermetallic
compounds.
If it is assumed that the crystalline fraction contains
both isolated defects and disordered regions (recoil
tracks) that give rise to the measured unit-cell
expansion, then, due to the composite nature of the
radiation-induced microstructure the total macroscopic
swelling, A VJV, , of zircon can be expressed as:
where f, is the mass fraction of crystalline phase, f, is
the mass fraction of amorphous phase, AVJV,, is the
average unit-cell volume change of the crystalline
phase, and AVJV, is the relative volume change
associated with the amorphous phase (Weber, 1993).
The unit-cell volume change, AV,,.V,, is given by
equation (6) with the fit parameters determined
directly from experimental measurements (Fig. 43). It
can be inferred that the volume change associated with
amorphization, AVJV,, is constant and equal to the
value of the macroscopic swelling at saturation,
(AV,/V,= 16.6% for the Pu-doped zircon and
AVa/Vo = 18.4% for natural zircons). Consequently, as
f, is known and f,= 1-f,, the total macroscopic
swelling, A V,/ V,, can be calculated using equation (4)
and is shown in Fig. 44, along with the experimental
data on total macroscopic swelling and the calculated
crystalline, f,AV,,/V,, and the amorphous, f,AV,/V,,
components. The calculated macroscopic swelling,
based on equation (4), provides an excellent fit to the
experimental data. The dominance of the unit-cell
expansion at low doses is clearly evident for both the
Pu-doped and natural zircons.
Most recently, particle irradiations (2 MeV He+
(Babsail et al., 1991); 0.8 MeV Nef, 1.5 MeV Ar+, 1.5
MeV Kr+, 0.7 MeV Kr+, 1.5 MeV Xe+ (Wang et al.,
1993)) have been used to induce amorphization in
zircon, and the results compared (the time scales in this
comparison range from 0.5 hr to 570 million yr!) to the
results of a-decay event damage (Weber et al., 1994).
High-resolution transmission electron microscopy
confirms that the damage in-growth processes are
similar for natural and ion-beam irradiated zircons
(Fig. 45). The defects are produced in varying concentrations
along each particle track, depending on mass
and energy of the particle. Figure 46 shows the
dependence of amorphization dose in zircon at room
temperature with the average energy transferred to
recoils per ion. These results show that the amorphization
dose for the heavier ions (Kr+ and Xe+ dose rates
of approximately 10-4-10- 3dpa s-‘) is nearly identical
to that determined from the 92 keV U+ recoil
emitted during a-decay in 238Pu-doped zircon with a
dose rate of 3 x lOmy dpa s-l. Thus, the amorphization
process is nearly independent of the damage
source (a-decay events versus heavy-ion beam irradiations)
and is consistent with models based on the
multiple overlap of particle tracks, indicating amorphization
occurs as a result of a critical defect
concentration.
Using in situ TEM to determine the critical amorphization
dose of zircon irradiated by 1.5 MeV Kr + as
Radiation effects in nuclear waste forms 111
20 I I
Pu-Doped Zircon
Natural Zircon
- 16
s
p 12
m
5
5 8
2 4
0.2 0.4 0.6 0.8
Dose (dpa)
Fig. 44. Contributions of crystalline fraction volume change, f,A Vu,/ V,,, and amorphous fraction volume change, f,A V,/ VO,to
the total macroscopic swelling, A V,/V,,, in Pu-doped and natural zircons based on a composite model of the radiation-induced
swelling. The solid curve is the sum off,AV,,/V, andfaAVJVO and provides an excellent fit to the experimental data. After
Weber et al. (1994).
a function of temperature provides an insight into the
competing roles of defect recovery (annealing) and
damage accumulation in a waste form ceramic
(Fig. 47). The dose for complete amorphization in
zircon increases with temperature in two stages. This
increase in amorphization dose with temperature is
believed to be due to a decrease in the average cascade
size (intracascade defect density) or shrinkage of the
amorphous regions (in overlapped cascades) as a
result of the thermal annealing of irradiation-induced
defects. The first stage (Stage-I), which occurs below
300 K, may be associated with the intracascade
recombination of closely-spaced defects. As the temperature
increases above 300 K, the Stage-I annealing
process in zircon is complete (all Stage-I defects have
recombined), and the dose for complete amorphization
from the remaining (unannealed) defect structure
within the cascade is constant, independent of temperature,
from 300 to 473 K. Above 473 K, the dose
for complete amorphization again increases with
temperature in a second stage (Stage-II), which may be
due to additional intracascade recombination from the
thermal migration of remaining point defects or
irradiation-enhanced epitaxial recrystallization (point
defect annihilation at the crystallineeamorphous inter-
face).
Because both the evolution of the amorphous state
and the dose for complete amorphization in the
1.5 MeV Kr+ irradiated zircon at room temperature
are similar to that observed above for the Pu-doped
and natural zircons, it is assumed that amorphization
under heavy-ion irradiation is also due to the multiple
112 R. C. Ewing et al.
Fig. 45. HRTEM images of the crystalline-to-amorphous transformation in zircon irradiated with 1.5 MeV Kr+ ions at 300 K:
(a) unirradiated, (b) 0.057 dpa, (c) 0.17 dpa, and (d) 0.34 dpa. Complete amorphization is achieved at 0.55 dpa. The spacing
marked in (c) and (d) is 0.33 nm [the (200) d-spacing]. Note the similarity of the images in (b), (c) and (d) to those of selfradiation
damaged zircon in Figs 41(a), (b) and (c). (Weber et al., 1994.)
t -2 MeV He+
300K
?? Zircon
. Pu-Doped Zircon
_._
0 1000 2000 3000 4000 5000
Energy to Recoils (eV/nm/ion)
Fig. 46. The dependence of amorphization dose in zircon at room temperature on the average energy transferred to recoils per
ion. The results show that the amorphization dose for the heavier ions (Kr+ and Xe’) is nearly identical with that determined
for the 92 keV U+ recoil emitted during a-decay of ‘s*Pu-doped zircon. (Weber et al., 1994.)
Radiation effects in nuclear waste forms 113
0.8
0.8
0.0 1
0
I
200 400 600 800 1000
Temperature (K)
Fig. 47. Temperature dependence of the dose for complete amorphization under 1.5 MeV Kr + irradiation. The data show two
stages associated with defect annealing that decrease the rate of amorphization. For Stage-I (T i 300 K), D, is estimated from a
data fit to be 0.35 dpa. For Stage-II (75473 K), Do is estimated to be 0.51 dpa. (Weber et al., 1994.)
overlap of cascades. The effects of temperature on
amorphization can then be modeled by kinetic processes
with activation energies, E,, that decrease the
average cascade size (or average defect density) and,
consequently, the rate at which the fraction of material
damaged directly by the cascades accumulates at
constant flux. Weber et al. (1994) have determined the
activation energy, E,, associated with each annealing
stage. The activation energy for Stage-I defect annealing
is 0.02 If:0.01 eV, and the activation energy determined
for Stage-II defect annealing is 0.31 kO.03 eV.
In addition to the activation energies, the critical
temperature, T,, above which complete amorphization
does not occur due to the dominance of Stage-II
defect annealing over damage production, can be
determined. The critical temperature, T,, for zircon
under these irradiation conditions is estimated to be
1100 K. The critical temperature for complete amorphization
in zircon for other ions and energies may be
different, as has been observed for Ca,La,(SiO,),O,
(Weber and Wang, 1994). The low value of the
activation energy for Stage-I suggests that this annealing
stage in zircon is associated with an intracascade,
close-pair recombination process at low temperatures
that is active up to 300 K where annealing is
apparently complete and above which it is assumed
that all close pairs recombine instantaneously. This
process, which is insufficient to prevent complete
amorphization, reduces the number of defects that
survive each cascade at temperatures above 300 K.
The activation energy for Stage-II annealing in zircon
is similar to activation energies reported in other
materials for an irradiation-enhanced epitaxial recrystallization
process associated with irradiationenhanced
defect mobility; however, because the
amorphization process in zircon is apparently due to
the accumulation of a critical defect concentration
from cascade overlap, the Stage-II annealing could
also be due to additional defect recombination and
annealing from the mobility of one or more defects
above 473 K. Although the exact nature of the secondstage
annealing process is not yet known, it competes
with the irradiation-induced amorphization process
and dominates above T,, thus preventing complete
amorphization. The activation energy for the second
annealing stage, 0.31 eV, is considerably less than the
activation energy, 3.6 eV, for thermal track annealing
(epitaxial thermal recrystallization) in zircon (Sandhu
et al., 1990) or the activation energies, 5.1 and 6.6 eV,
for bulk thermal recrystallization of fully amorphous
zircon (Weber et al., 1994). This type of thermal
annealing analysis is essential in the study of crystalline
ceramic waste forms in order to determine their
final damage state as a function of the thermal history
of the waste form.
Because of the important implications in determining
the U/Pb systematics of zircons for geologic age
dating, there are data on the increased loss of nuclides
as a function of increasing degree of damage (Pidgeon
et al., 1966; Ewing et al., 1982b; Tole, 1985; Suzuki,
1987; Sinha et al., 1992), as well as efforts to relate
these data to nuclear waste containment (Gentry et al.,
1982; Ludwig et al., 1984; Sinha et al., 1991). In a study
of a suite of natural zircons with varying contents of
uranium and thorium, Ewing et al. (1982b) found that
over a dose range of 1019-1021 a-decay events g-‘,
114 R. C. Ewing et al.
there was an increase of one order of magnitude in the
wt% of zircon that is dissolved and an increase in the
leach rate from 2.9x lo-* to 2.3 x lo-’ gem-’
day- ‘. For completely metamict zircons (dose 2 10”
cc-decay events g-‘) the leach rate was as high as
1.8 x 10m6 g cm-’ day-‘. These results are consistent
with the observation that loss of nuclides or increased
solubility is, at least in part, related to the increasing
level of radiation damage (Pidgeon et al., 1966; Tole,
1985; Sinha et al., 1992).
This unusually complete data set for damage effects
in zircon provides a basis for confirming the use of
accelerated damage processes (actinide-doping of
synthetic phases and heavy-particle irradiations) to
study the long-term behavior of ceramics in a radiation
field. Similar studies should be possible for members of
the pyrochlore group as these minerals occur abundantly
in nature and can be synthesized with a wide
variety of actinide-bearing compositions (Chakoumakos
and Ewing, 1985).
5.2.6. Apatite. Several rare-earth silicates with the
apatite structure have been observed or proposed as
actinide host phases in a nuclear waste glass (Weber et
al., 1979), supercalcine (McCarthy, 1977; McCarthy et
al., 1979), a glass ceramic waste form (De et al.,
1976a,b), and a cement waste form (Jantzen and
Glasser, 1979). The apatites in nuclear waste solids are
rare-earth silicate isomorphs of natural apatite,
Ca,,(PO,),(F, OH),, and generally have the composition:
Ca,_,RE,+,(SiO,),_,(PO,),O, (where RE=
La, Ce, Pr, Nd, Pm, Sm, Eu and Gd). The actinides
readily substitute for the rare earths in this hexagonal
crystal structure when it is present in the waste form.
As noted above in the study of a Cm-doped nuclear
waste glass (Weber et ul., 1979; Weber and Roberts,
1983) crystallites of Ca,(Gd, Cm),(SiO,),(PO,)O,
transformed from a crystalline to an amorphous state
as a result of a-decay of the incorporated Cm. The
volume expansion associated with the amorphous
transformation resulted in microfracturing of the
simulated waste glass. Similarly, in a study of
Cm-doped supercalcine (Rusin et al., 1979), the
originally crystalline apatite phase, Ca,(Re,
Cm),(SiO,),O,, transformed to an amorphous state
due to a-decay. The unit cell of this phase was
estimated to expand 2.5% before the amorphization
process was complete. More specific details of the
amorphization process in these phases, structural
data, and the effects of amorphization on the integrity
of these apatite phases could not be obtained from the
study of these multiphase waste forms. Several natural
apatite minerals with significant amounts of rare
earths and silica are known to have an appreciable
thoria content. These minerals are reported to be
partially or fully metamict (amorphous) due to selfradiation
damage from the alpha decay of the Th
(Lindberg and Ingram, 1964; Haaker and Ewing,
1981), but detailed structural data on the effects of
radiation in these minerals have not been reported.
Studies of fission track dating and annealing (Koul,
1979) and ion-beam-induced amorphization (Wang et
al., 1994) in naturally occurring apatites have also
been reported.
In order to understand the effects of self-radiation
damage from cc-decay on apatite phases in nuclear
waste forms, an extensive study of single-phase,
244Cm-doped Ca,(Nd),(SiO,),O, has been carried
out (Weber, 1982,1983,1993; Wald and Weber, 1984;
Weber and Matzke, 1986a,b). Radiation damage from
x-decay in this material results in the simultaneous
accumulation of simple structural defects (i.e. vacancies
and interstitials on both the cation and anion
sublattices) and amorphous regions that exhibit loss of
long-range order. The amorphous regions eventually
overlap to produce a completely amorphous state
(Fig. 48). Consequently, as the radiation damage
process proceeds, the fraction of crystalline material,
f,, decreases and the fraction of amorphous material,
f,, increases. For this material, Weber (1993) has
shown that amorphization primarily occurs directly in
the cascades of the recoil nuclei emitted during a-decay
of Cm and that the increase in amorphous fraction
with dose, D, is given by the expression:
f,= 1-exp(-B,D), (5)
where B,, the mass of amorphous material produced
per decay event, has been determined to be 4.42 x
10-i’ g. The change in amorphous fraction with dose
for Cm-doped Ca,Nd,(SiO,),O, is shown in Fig. 49.
The crystalline fraction, f, = 1-I,, is consistent with
the observed decrease in diffracted X-ray intensity
(Weber, 1993).
The accumulation of structural defects in the
crystalline fraction of material results in an increase in
unit-cell volume with cumulative dose. The unit-cell
volume expansion, AV,,/V,, which is shown in
Fig. 50, follows the characteristic exponential behavior
predicted by models (Nellis, 1977; Weber et al.,
1991a) for the accumulation of isolated structural
defects and is given by the expression:
Av,,=A,,Cl-exp(-B,,D)l, (6)
where A,, is the relative unit-cell volume expansion at
saturation and B,, is the rate constant (per unit dose)
for the simultaneous recombination of structural
defects during irradiation. The values for A,, and B,,,
Radiation effects in nuclear waste forms 115
Wf 0.1
Fig. 48. Microstructural changes in Cm-doped CazNd,(SiO&O, at (a) 0.76 x lo’*, (b) 1.32 x 1018, (c) 2.02 x lo’*, and
(d) 2.55 x 10” m-decays g-’ (Weber, 1993).
116 R. C. Ewing et al.
Dose, 1O’* alpha decays/g
Fig. 49. Amorphous fraction,f,, in Cm-doped Ca,Nd,(SiO,),O, as a function of dose. After Weber (1993).
: Ca,Nd,(SiO,),O,
4Dose,
1O’* alpha decays/g
Fig. 50. Unit-cell volume expansion, AVJVo, in Cm-doped Ca,Nd,(SiO,),O, as a function of dose. After Weber (1993).
have been determined to be 4.95% and 5.08 x 10-l’ g
(Weber, 1993). The unit-cell expansion is anisotropic
with the expansion of the unit cell along the cc-axis
about twice that along the c-axis. This anisotropic
behavior is similar to that observed in ion-irradiated
Ca,Nds(SiO,),O, and shown to be consistent with
topological constraints on the changes in the polyhedral
connectivity within this structure (Weber et al.,
1991).
The overall macroscopic swelling, A V,/V, , which
has contributions from the crystalline and amorphous
fractions of material, is given by equation (4) (Weber,
1993). The overall macroscopic swelling determined
experimentally from changes in immersion density is
shown in Fig. 51 along with the predicted contributions
from the crystalline and amorphous fractions.
The solid curve is the sum of the crystalline and
amorphous components (equation (4)) and provides
an excellent fit to the data. The nearly equal contribution
of the unit-cell expansion and amorphization
process to the total macroscopic swelling at low doses
is clearly shown as is the dominance of the amorphization
process at high doses.
The normalized dissolution rate of Cm-doped
Ca,Nd,(SiO,),O, at 90°C in deionized water
increased by about a factor of 10 as a result of
Radiation effects in nuclear waste forms 117
- Ca,Nd,(SiO,),O,
8-
0 1 2 3 4 5
Dose, 1O’* alpha decays/g
Fig. 51. Contributions of crystalline and amorphous components to the overall macroscopic swelling in Cm-doped
Ca,Nds(SiO,),O, as a function of dose. After Weber (1993).
N 1.4
FE
2 1.2
H
-8 1.0
i! 0.8
r
ge 0.6
f 0.4
5
E 0.2
IL
0.0
I ’ ” 8 II ’ ”
_ Ca,Nd,(SiO,),Op
Dose, 1018alpha decays/g
Fig. 52. Fracture toughness in Cm-doped Ca,Nd,(SiO,),O, as a function of dose. After Weber and Matzke (1986b).
radiation-induced swelling and amorphization
(Weber, 1983; Wald and Weber, 1984). The dissolution
appears to occur congruently in this apatite phase.
The mechanism for the enhanced dissolution of the
amorphous phase is most likely associated with the
high degree of disorder, broken bonds, and more open
structure.
The fracture toughness in Cm-doped Ca,Nd,
(SiO,),O, increased with cumulative dose to a broad
maximum and then decreases slightly (Fig, 52). This is
similar to the behavior discussed above in pyrochlores
and zirconolite (Clinard et al., 1985; Weber et al.,
1986a). The increase in fracture toughness is attributed
to the composite nature of the microstructure. At low
to intermediate doses, the microstructure consists of
amorphous tracks in a crystalline matrix, and this
inhibits crack propagation and increases the fracture
toughness. As the amorphous phase becomes the
dominant matrix at high doses with remanent crystallites,
the fracture toughness decreases slightly.
The activation energy for recrystallization from the
fully amorphous state in Cm-doped Ca,Nd,
(SiO,),OZ has been rigorously determined by Weber
(1983) to be 3.lkO.2 eV. Recrystallization of the
118 R. C. Ewing et al.
0 100 200 300 400 500 600 700
Temperature, K
Fig. 53.Temperature dependence of the dose for complete amorphization in Ca,La,(SiO,),O,. After Weber and Wang (1994).
amorphous state occurs in a single exothermic recovery
stage that releases 130 J gg ’of stored energy. Full
recovery of the density and original apatite crystal
structure occurs during this single-stage recrystallization
process.
The temperature dependence of amorphization in
single crystals of the apatite phase, Ca,La,(SiO,),O,,
has been determined in recent ion-irradiation studies
by Weber and Wang (1993, 1994). The dose for
complete amorphization in Ca2La,(Si0,),0z is independent
of recoil-energy spectra at low temperatures
(< 15 K) and increases with temperature in two stages
for both 0.8 MeV Ne+ and 1.5 MeV Krf irradiations
(Fig. 53). The low-temperature stage (below 200 K)
has an activation energy of 0.01 f 0.003 eV and may be
associated with close-pair recombination. The hightemperature
stage (above 200 K) has an activation
energy of 0.13 f0.02 eV and may be associated with
irradiation-enhanced defect mobility. The critical
temperature above which complete amorphization
does not occur increased from 360 K for Ne+ to 710 K
for Krf. The amorphization dose below room temperature,
0.3-0.5 dpa for the heavier Kr+ irradiations,
is similar to the dose, 0.4 dpa, determined for full
amorphization of Cm-doped Ca,Nd,(SiO,),O, at
room temperature. The dose for complete amorphization
increases more rapidly for the lighter Nef
irradiations because of the larger fraction of recoverable
defects and smaller cascades that are produced by
the low-energy recoils.
5.2.7. Monazite. Monazite is unique, as it has been
proposed as a single-phase ceramic to incorporate a
wide variety of nuclear wastes, particularly those rich
in actinides (Boatner et al., 1980; Boatner and Sales,
1988). The mineral monazite is a mixed lanthanide
orthophosphate, LnPO, (Ln = La, Ce, Nd, Gd, etc.),
that often contains significant amounts of Th and U
(up to 27 wt% combined). As a result, many natural
monazites have been subjected to significant cc-decay
event damage over geologic time (some samples are up
to 2 billion yr of age). In spite of the large radiation
doses received by natural monazites, Ewing (1975) and
Ewing and Haaker (1980) have noted that monazite is
always found in the crystalline state. The apparent
resistance of natural monazites to radiation-induced
amorphization was an important factor in the initial
proposal of monazite ceramics as potential candidates
for the immobilization of nuclear wastes.
Karioris et al. (1981) have established that natural
monazite (Ce, La, Y)PO, can be readily transformed
to an amorphous state by irradiation with 3 MeV Ar+
ions to moderate doses. Gowda (1982) determined
that under 3 MeV Ar+ ion irradiation the damage
cross-section for amorphization was 0.54 f 0.10 nm2
for natural monazite and 0.56 f 0.08 nm2 for synthetic
monazite, CePO, . In natural monazite that was
irradiated to a partially amorphous state (diffraction
intensities were reduced by one-third), complete
recovery of the fully crystalline state occurred after
annealing at 300°C for 20 hr (Karioris et al., 1981).
The apparent stability of natural monazite is attributed
to this relatively low temperature of recovery of
the radiation damage. Differential scanning calorimetry
(DSC) measurements of natural (Gowda, 1982)
and synthetic (Gowda, 1982; Ehlert et al., 1983)
Radiation effects in nuclear waste forms 119
monazites irradiated with 3 MeV Arf ions have been
performed. In natural monazite, the stored energy
released during recrystallization was 33 J g- ‘, the
peak in the release rate occurred at 450°C and the
activation energy associated with the recovery process
was estimated to be 2.77 eV. For synthetic monazite,
CePO,, the stored energy release was 30.5 f 1.3 J g- ‘,
the temperature for maximum release was slightly
lower at 370°C and the activation energy was
estimated to be 2.7 + 0.3 eV.
Results by Sales et al. (1983) suggest that the leach
rate of the host matrix of a synthetic monazite, LaPO,,
containing simulated waste remains low even after the
material has been transformed to an amorphous state
by irradiation with 250 keV Bif ions. However, work
by Eyal and Kaufman (1982) on natural monazite
indicate that there is a preferential dissolution of the
radionuclide daughter-products, 234U, ‘j”Th and
228Th, by factors of between 1.1 and 10 relative to the
structurally incorporated parent isotopes 238U and
232Th. This isotopic fractionation is attributed to
radiation damage in the tracks of the recoil nuclei
emitted during cc-decay of the parent isotopes. While
there have been some concerns regarding these results
(Boatner and Sales, 1988), the increases in dissolution
rates are similar to those observed in other actinidehost
phases.
5.2.8. Sphene (= titanite). Sphene, CaTi(SiO,)O
(the proper mineral name is titanite) is a prominent
phase in the sphene-glass ceramic. Sphene is an
actinide-host (the structure will not accommodate Cs;
but will accommodate minor concentrations of Sr),
thus damage due to a-decay of actinides is an issue of
concern.
Although there have been studies using heavy
particle irradiations, 3 MeV Ar+ and 280 keV Bi2+
(Vance et al., 1984; Stevanovic et al., 1989) and studies
using y-radiation, 200 keV electrons (Vance et al.,
1986), most work has focused on natural sphenes
which contain U and Th (Vance and Metson, 1985;
Hawthorne et al., 1991). Up to doses of2 x 10z7 e m-*
there was no evidence of structural damage. The Ar+
irradiation indicated that a fluence of 7 x 10” a-decay
events g - ’ was required to produce complete metamictization.
This is in good agreement with the value
of 5 x 1Or8 a-decay events gg ’ required to reach the
metamict state in natural sphenes. Leaching experiments
on damaged sphenes indicated that the increase
in dissolution rate was no more than a factor of five;
however, leaching experiments on Bi’+ implanted
sphene and sphene-ceramics to doses of 6 x lOI ions
-2 resulted in an order of magnitude increase in the
Ltch rate (Stevanovic et al., 1989).
The study of natural sphenes by Hawthorne et al.
(1991) was unusual in that a wide variety of techniques
(electron microprobe analysis, powder X-ray diffraction,
single-crystal X-ray diffraction, infrared spectroscopy,
Mossbauer spectroscopy, magic-angle spinning
nuclear magnetic resonance spectroscopy, high-resolution
transmission electron microscopy, extended
X-ray absorption spectroscopy and X-ray absorption
near edge spectroscopy) were used to monitor structural
changes with increasing a-decay event dose. The
results substantiated the model for damage in-growth
outlined by Murakami et al. (1991) and illustrated in
Fig. 40. The work did suggest that the damage process
is accompanied by the reduction of Fe3+ to Fe’+,
which resides mainly in the aperiodic domains.
Valence changes of this type can have significant effects
on the solubility of certain actinides (e.g. uranium and
plutonium) particularly when they are partitioned into
the aperiodic domains.
6. FUTURE RESEARCH NEEDS
The issues surrounding the permanent disposal of
nuclear waste are both political and technological. The
delays in winning public acceptance for current
nuclear waste policy and technology threaten shutdown
of current commercial nuclear power plants and
hamper acceptance of new nuclear power reactors,
both of which could play a vital role in the U.S.
strategy to minimize human impact on global processes,
such as CO, emissions. Thus, the successful
demonstration that nuclear waste can be disposed of
safely is an important step in the development of
nuclear power. One of the critical barriers to radionuclide
migration to the biosphere is the nuclear waste
form; however, the physical and chemical properties of
the waste form may change dramatically as a result of
irradiation damage. Additionally, radiation-solid
interactions might also affect important geochemical
processes that are relied upon for nuclide retardation
(e.g. sorption of radionuclides on clays, radiolysis of
transporting fluids, radiolytic oxidation and corrosion
of UO,, and radiation damage in secondary phases
which by their solubilities control radionuclide concentrations
in the ground water). Unfortunately, due
to the state of understanding and limited research in
this area (Weber et al., 1991b), the nature, dimensions,
and implications of these radiation effects remain
elusive, particularly with regard to how these effects
impact regulatory criteria for radionuclide release.
More specifically, certain topics require immediate
attention:
120 R. C. Ewing et al.
1. 7’he efSect of ionizing radiation on borosilicate
waste glasses, particularly the possibility and impact of
radiolytic decomposition and bubble formation, are
neither well-understood or under current investigation.
Studies are needed to determine the temperature
and dose rate dependence of these phenomena and
their impact on radionuclide release rates. A thorough
understanding of the synergistic effects of ionizing and
displacive radiation on waste glass behavior is needed.
This should include the effects of radiolytic decomposition
and microstructural evolution on increased
surface area and leachability, which determine the
release rate of radionuclides under unsaturated conditions.
Also, there are only limited (and in most cases
questionable) data on the effect of displacive events
(e.g. a-recoil damage) on the leachability of borosilicate
waste glasses.
2. Computer simulation of displacement processes.
The final state of a radiation-damaged glass or ceramic
is a result of the temperature-dependent balance
between damage and annealing kinetics. The initial
damage stage can only be ascertained by models of
collision cascade formation (these events occur in
periods of lo-“- lo-l2 set). Very little is known
about the process of collision cascade formation in
glasses and ceramics, as compared to the state of
knowledge for metals. As an example, we have only
very limited data on the displacement energies of
cations in complex, multi-atom ceramics which are
typical of nuclear waste form assemblages.
3. Radiation-induced periodic-to-aperiodic transformations
(amorphization). Little is known of the
mechanisms of crystalline-to-amorphous transformations,
particularly in ceramics. There has been no
systematic evaluation of the role of structure-type,
bond-type, bond-strength and structural topology on
this phenomenon. This knowledge is necessary for the
efficient evaluation of crystalline ceramic waste phases
and their response to different types and doses of
radiation. This mechanistic understanding of the
amorphization process is particularly important in
extrapolating the long-term behavior of ceramics and
glasses in a radiation field.
4. Structure of the aperiodic state. Preliminary
research has demonstrated that there are many
possible structures and topologies in aperiodic materials.
These structural variations may reflect the mode
of amorphization (e.g. radiation-induced, pressureinduced
or quenched from a melt), and the aperiodic
structure may change with increasing level of damage.
These changes can have significant effects on the
behavior (e.g. leach rate, devitrification and mechanical
properties) of aperiodic phases, such as the
borosilicate glass waste form.
Acknowledgements-The authors gratefully acknowledge the
support that they have received from the Office of Basic
Energy Sciences of the U.S. Department of Enerav (WFC:
contract #W-740%ENG-36; RCE: grant #DE-&3-
93ER45498: WJW contract #DE-AC06-76RL0 18801. The
BES suppoit of basic, interdisciplinary research on radiation
effects in minerals and ceramics has contributed greatly to the
present understanding of radiation damage effects in nuclear
waste forms. F. W. Clinard, Jr expresses his appreciation for
the contributions of his coworkers in the Materials Science
and Technology Division at Los Alamos National Laboratory;
W. J. Weber acknowledges the contributions of his
coworkers at the Pacific Northwest Laboratory; R. C. Ewing
thanks his colleagues and students for their collaboration in
the field of radiation damage effects in minerals and ceramics
and the good times at El Patlo. All errors of omission and
interpretations of the data in the literature remain those of the
authors. Finally, R. C. Ewing thanks his colleagues in the
Department of Earth Sciences of Aarhus University for their
support during his visit in 1994 during which most of the
writing of this review was completed.
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