Full text of DOI:10.1557/mrs2001.177

www.mrs.org/publications/bulletin
including effects of local stress, plastic
accommodation, and radiation-induced
transmutants and mass transport, are only
partially understood.
F
undamental Studies
The radioactive decay of Pu presents
another challenge to understanding this
unique material and its evolution toward
thermodynamic equilibrium. Through a
combination of ꢀ and ꢅ decay, Pu trans-
mutes to form both insoluble He and
impurities that likely stabilize (Am) and
possibly destabilize (U and Np) the fcc
of Plutonium Aging
Brian D. Wirth, Adam J. Schwartz,
Michael J. Fluss, Maria J. Caturla,
Mark A.Wall, and Wilhelm G. Wolfer
3,5
ꢁ phase. This ingrowth occurs at a rate
of approximately 41 at. ppm He per yr,
7
5 at. ppm Am per yr, 35 at. ppm U per yr,
3
and 6 at. ppm Np per yr. In addition to
transmutation, the decay of Pu produces
radiation damage and thus high popula-
tions of point defects (vacancies and self-
interstitial atoms) and defect clusters.
Radiation damage from ꢀ decay in pluto-
nium occurs at a rate of ꢀ0.1 dpa/yr (dis-
placements per atom per year).15 While
the majority of displaced atoms quickly
return to lattice sites, the remaining va-
cancies and self-interstitials drive micro-
structural evolution. Experience from the
nuclear-energy industry has demonstrated
the microstructural consequences of radia-
tion exposure, which include the nuclea-
tion and growth of extended defects such
as voids (swelling) and gas bubbles,
changes in dislocation structures, and ac-
celeration and alteration of normal alloy
phase-decomposition sequences. The cor-
responding consequences to mechanical
properties typically include hardening,
reductions in ductility and fracture tough-
ness, higher creep rates, lower creep rup-
ture times, and increased susceptibility to
various environmentally assisted cracking
Introduction
Plutonium metallurgy lies at the heart of
The remarkably less dense but more duc-
science-based stockpile stewardship.1 One
aspect is concerned with developing pre-
dictive capabilities to describe the prop-
erties of stockpile materials, including an
assessment of microstructural changes with
age. Yet, the complex behavior of pluto-
nium, which results from the competition
of its 5f electrons between a localized
–3
tile fcc ꢁ phase is stable between tempera-
7
tures of 360ꢂC and 463ꢂC. Not surprisingly,
the ꢁ phase is preferred by Pu metallur-
gists and can be stabilized down to room
temperatures by the addition of Group IIIb
8,9
elements like aluminum and gallium.
The binary phase diagrams have been ex-
tensively studied during the period of
7
–14
(atomic-like or bound) state and an itiner-
1950–2000.
Thus, it is well established
4,5
ant (delocalized bonding) state, has been
challenging materials scientists and physi-
that additions of Al, Ga, and Am promote
the stability of the ꢁ phase, while additions
of U and Np reduce the stability of the
ꢁ phase, although the underlying mecha-
nisms responsible for this behavior are not
6
cists for the better part of five decades. Al-
though far from quantitatively absolute,
electronic-structure theory provides a de-
scription of plutonium that helps explain
the unusual properties of plutonium, as
5
well understood.
While the binary phase diagrams have
been studied extensively, the kinetics re-
sponsible for transitions to thermodynamic
equilibrium are often sluggish at best, es-
pecially near room temperature, making
equilibrium difficult to determine. Indeed,
5
recently reviewed by Hecker. (See also the
article by Hecker in this issue.) The elec-
tronic structure of plutonium includes five
5f electrons with a very narrow energy
width of the 5f conduction band, which
results in a delicate balance between itin-
erant electrons (in the conduction band)
or localized electrons and multiple low-
energy electronic configurations with nearly
equivalent energies.4 These complex elec-
tronic characteristics give rise to unique
macroscopic properties of plutonium that
include six allotropes (at ambient pres-
sure) with very close free energies but
large (ꢀ25%) density differences, a low-
symmetry monoclinic ground state rather
than a high-symmetry close-packed cubic
phase, compression upon melting (like
water), low melting temperature, anoma-
lous temperature-dependence of electrical
8,14
16–20
Timofeeva and co-workers have recently
shown that the U.S. phase diagram for the
Pu-Ga system may not be correct and that
the Pu-Ga ꢁ phase is not stable at room
temperature, but instead is only meta-
stable. Timofeeva demonstrated that at
ꢀ100ꢂC, ꢁ-Pu-Ga undergoes a eutectoid
transformation to the monoclinic ꢀ phase
processes.
Our approach to studying Pu metal and
alloys involves coupling experimental
and modeling techniques to monitor the
aging of old material, in addition to fun-
damental studies of key defect production
and transport mechanisms. The ultimate
objective is to predict property changes
of Pu-Ga alloys during aging and thus
develop quantitative predictions of their
useful storage life. In this brief article, we
will focus on the development of a funda-
mental understanding of radiation damage
and defect accumulation obtained through
this close coupling of experiments and
modeling, as well as illustrate some ex-
amples of our recent work to characterize
Pu microstructures and phase stability
obtained through transmission electron
microscopy (TEM).
,5
and the intermetallic Pu Ga, and thus
3
thermodynamic equilibrium at tempera-
8
,14
tures below 100ꢂC is a two-phase mixture.
However, the ꢁ-to-ꢀ ꢃ Pu Ga transforma-
3
tion is diffusional in nature and in this case
is exceedingly slow. Timofeeva’s work ex-
perimentally confirms the true equilib-
rium phase diagram and Pu-Ga ꢁ-phase
metastability as previously calculated by
5
resistance, and radioactive decay. Addi-
tionally, plutonium readily oxidizes and is
toxic; therefore, the handling and funda-
mental research of this element is very
challenging due to environmental, safety,
and health concerns.
13
Adler from thermochemical data. Of
perhaps greater importance is a displacive
or martensitic phase transformation from
ꢁ to ꢀꢄ (note that ꢀꢄ is a slightly modified
monoclinic ꢀ phase containing insoluble
Ga), known to occur at low temperatures.
Yet, the kinetics of this transformation,
Unalloyed Pu has a dense, monoclinic
structure (ꢀ phase) at room temperature
that is extremely brittle and highly reactive.
Alpha Decay
Radiation damage, including the self-
induced damage of ꢀ decay, can signifi-
MRS BULLETIN/SEPTEMBER 2001
679

Fundamental Studies of Plutonium Aging
cantly alter the underlying material
microstructure and thus impact a wide
range of materials properties. Microstruc-
tural evolution results from the generation
of primary damage in spatially correlated
high-energy displacement cascades and the
subsequent diffusion, clustering, and long-
range transport of the vacancy and self-
interstitial defects, which are inherently
coupled to solute and impurity diffusion.
A key to understanding and predicting
radiation damage and its attendant con-
sequences is developing fundamental
knowledge of both the primary defect
production and the defect diffusion and
clustering kinetics. The primary source of
damage in Pu metal and alloys arises from
and self-interstitial) pairs (ꢀ350 Frenkel
simulation of an 80-keV displacement cas-
cade in lead, performed on highly parallel
ASCI (Accelerated Strategic Computing
Initiative) supercomputers with 10 million
atoms for a total time of 200 ps. The first
few collisions occur at very high energy,
approximately several kiloelectronvolts,
and the cascade splits into subcascades,
two of which can be observed in Fig-
ure 1b. The number of displaced atoms
and the energy density increase rapidly
within the cascade volume, reaching a
peak within about 10 ps, at which time the
center of the cascade (or subcascades) has
a liquidlike structure and then slowly
cools, due to the high atomic mass and low
melting temperature, over about 100 ps.
The final damage produced is highly
spatially correlated, unlike the ꢀ-particle
track, forming groups of self-interstitial
and vacancy clusters.
It is unlikely that we can experimentally
capture all of the dynamics of the recoil
cascade. However, it may be possible in a
self-irradiating material like Pu to validate
the cooled cascade structure by perform-
ing in situ TEM observations at ultralow
temperatures that essentially freeze the
cascade structure in place. A key to such
an experiment is the accurate TEM image
simulation of the MD predictions. This is
a capability that we are jointly develop-
ing with R. Schaublin and co-workers at
École Polytechnique Fédérale de Lausanne
(EPFL) and the Paul Scherrer Institute (PSI)
in Switzerland. Experiments built on this
capability may now be feasible and are
likely to provide additional insight into
the details of the important source-term
structure. One can also imagine “follow-
ing” the annealing of these features and
thereby validating and refining the sub-
sequent kinetics calculations that describe
the mass-transport processes and damage
accumulation.
pairs per decay, represented by the red
dots in Figure 1a) and the He atoms them-
selves (represented by the black dot in Fig-
ure 1a), which come to rest in the Pu lattice
ꢀ10 ꢆm from their birth in the decay
event.
On the other hand, the fast-moving
ꢀ86-keV U recoil has a negligible mass
difference with Pu and rapidly undergoes
a branching series of high-energy, nearly
elastic collisions, which in turn produces
secondary collisions and generates a cas-
cade of vacancy and self-interstitial de-
fects. High-energy displacement cascades
evolve over very short times (ꢀ100 ps) and
small volumes, with characteristic length
scales of 50 nm or less, and are directly
amenable to molecular-dynamics (MD)
simulations. The physics of primary dam-
age production in high-energy displace-
ment cascades has been extensively studied
ꢀ
decay, as illustrated in Figure 1.
The ꢀ decay of a plutonium atom pro-
duces an ꢀ particle of ꢀ5 MeV and a
3,15
uranium recoil atom of ꢀ86 keV. The
spatial distribution and character of the
damage produced by these two particles
are very different in nature, and yet it is
the eventual interaction and evolution of
these spatially uncorrelated primary de-
fects that may, over time, drive materials
evolution and aging processes, thus pro-
ducing microstructural changes over time
scales as long as many decades. The fast
moving ꢀ particle has a range of approxi-
mately 10 ꢆm, losing the majority of its
kinetic energy through electronic excita-
tions and undergoing a long track of low-
energy atomic collisions. Figure 1a shows
a calculated He ion track; each red dot
represents an atomic collision that trans-
fers ꢀ50 eV, and the black dot represents
the location of the He nucleus when it
comes to rest. The calculation was per-
formed within the binary collision approxi-
mation, using the simulation program
22–24
with MD simulations.
The key conclu-
sions of the high-energy (ꢇ20 keV) simu-
lations are (1) intracascade recombination
of vacancies and self-interstitials results in
ꢀ25% of the defect production expected
from displacement theory (e.g., ꢀ575 com-
15,25
pared to ꢀ2300);
(2) many-body colli-
sion effects produce a spatial correlation
(separation) of the vacancy and self-
interstitial defects; (3) substantial cluster-
ing of the self-interstitials and, to a lesser
extent, the vacancies occurs within the
cascade volume; and (4) high-energy dis-
placement cascades tend to break up into
lobes or subcascades, which may also en-
23–25
hance recombination.
The lack of reliable semiempirical inter-
atomic potentials and electronic-structure
methods hinders our ability to directly
simulate the high-energy U recoil in Pu;
we have used low-melting-temperature,
fcc metals as surrogates. Figure 1b shows
a sequence of “snapshots” from an MD
21
TRIM (TRansport of Ions in Matter). The
primary defects produced along the He
track are largely isolated Frenkel (vacancy
Isochronal Annealing
Defect diffusion and clustering kinetics
represent the other main ingredients re-
quired for the reliable prediction of micro-
structural evolution resulting from Pu
ꢀ
decay. These properties can be deter-
mined experimentally by monitoring the
change in electrical resistance during low-
temperature (ꢀ10 K) irradiation and sub-
sequent isochronal annealing. However,
the complex 5f electronic structure and
behavior of Pu make data analysis very
challenging, and thus we use kinetic
Monte Carlo (KMC) simulations of defect
evolution during a simulated isochronal
annealing history to complement data
interpretation.
Figure 1. Alpha decay in plutonium produces an ꢀ particle of ꢀ5 MeV and a uranium recoil
atom of ꢀ86 keV. (a) A representative He ion track, with each red dot representing a
low-energy atomic collision and the black dot representing the resting position at the end
of the approximately 10-ꢆm range. (b) Heavy uranium recoil produces a dense collision
cascade over a ꢀ50-nm region, shown in a series of molecular-dynamics “snapshots”
The annealing recovery and defect ki-
netics of fcc metals following low-
(red spheres represent self-interstitial atoms; green spheres are vacancy sites).
680
MRS BULLETIN/SEPTEMBER 2001

Fundamental Studies of Plutonium Aging
temperature electron irradiation have been
extensively studied over the past 30 years
and are defined by five recovery stages.26
Stage I occurs at low temperatures, typi-
cally a small fraction of the melting tem-
a sharp, well-defined recovery stage at
tion of slightly larger vacancy clusters in
the self-irradiated specimen. Again, the
KMC simulations capture the qualitative
behavior of the Stage V shift, although
additional modeling is clearly required to
further validate our understanding of both
the primary damage production and the
defect kinetics.
about 45 K, consistent with previous self-
irradiation and annealing recovery stud-
27
ies of Pu-Al alloys. We believe the sharp
recovery stage in the self-irradiated speci-
men is actually Stage II recovery and arises
from the rapid one-dimensional diffusion
and annihilation of the relatively large
self-interstitial clusters directly formed in
displacement cascades. The more sluggish
and extended recovery in the proton-
irradiated specimen is a consequence of
the evolution of isolated self-interstitials,
likely including strong interactions (trap-
ping and de-trapping) with impurities
and clustering, albeit with smaller cluster
sizes than those directly produced in cas-
cades, which undergo one-dimensional dif-
fusion. Similar (to the proton) Stage I/II
behavior is commonly observed in electron-
irradiated alloys and metals with large im-
perature, 0.02T , with the annihilation of
m
mobile self-interstitial atoms, largely
through recombination with immobile va-
cancies combined with a lesser degree of
self-interstitial clustering and trapping at
impurities. Electron irradiation of metals
produces predominately isolated Frenkel
pairs, and the smaller amount of recovery
observed in Stage II is generally attributed
to the migration of small self-interstitial
clusters and the de-trapping of self-
interstitials from impurities. At higher
Helium Accumulation
As described previously, the buildup of
He occurs at a rate of ꢀ41 at. ppm per yr.
It is reasonable to expect that the insoluble
He gas will quickly cluster to form bub-
bles. Indeed, atomistic studies of He va-
cancy interaction in nickel and other fcc
metals show that helium vacancy cluster
complexes are strongly bound.28 As a re-
sult, a critical helium bubble nucleus con-
sists of just two helium atoms and one or
two vacancies at these relatively high
helium generation rates. Figure 3 shows
the results of a model prediction for the
He bubble density and diameter during
aging, assuming that helium undergoes
substitutional diffusion, with an activa-
tion energy of 0.68 eV at 35ꢂC and only
considering the average bubble size. Ex-
perimental positron-annihilation lifetime
measurements by Colmenares and Howell
of 21- and 35-yr-old Pu-Ga alloys indicate
the existence of He-vacancy complexes
temperatures, ꢀ0.2T , the vacancies begin
m
to migrate, and Stage III recovery results
from vacancy annihilation and clustering.
Stage IV is attributed to vacancy inter-
actions and de-trapping from impurities.
Annealing is complete after Stage V at tem-
26
purity concentrations. KMC simulations
predict qualitatively similar Stage I/II
behavior for the proton and self-irradiated
specimens, respectively, and provide key
insight into the annealing kinetics. The
simulations and experimental results cur-
rently are in good agreement regarding
the annealing recovery temperatures, but
are not as good for the magnitude of re-
covery. Some of this discrepancy may be
due to specimen beam-heating during the
low-temperature proton irradiation, and
thus Stage I is difficult to precisely deter-
mine experimentally.
peratures of ꢀ0.45T , with the dissolution
m
26
of vacancy clusters.
Figure 2 shows the isochronal annealing
recovery (up to temperatures of ꢀ400 K)
for an irradiated Pu-Ga alloy measured
experimentally (Figure 2a) and predicted
by KMC simulation (Figure 2b) following
low-temperature (ꢀ10 K) irradiation by
both 3.8-MeV protons and self-induced
29,30
ꢀ
decay. The primary damage produced
with a ratio of ꢀ1 He/vacant site.
TEM
by high-energy proton irradiation is simi-
lar to that produced by electrons and con-
sists of predominately isolated Frenkel
pairs, while the self-induced ꢀ decay pro-
duces both isolated point defects from the
He ion track and a dense, high-energy col-
lision cascade. The difference in primary
damage state is clearly visible in the low-
temperature (Stages I and II) annealing re-
covery and the completion of Stage V.
The proton-irradiated specimen exhibits
a slow recovery from about 30 K to 150 K,
while the self-irradiated specimen shows
characterization of aged Pu-Ga alloys is
currently under way to verify the modeling
predictions of Wolfer. The ability to image
1-nm microstructural features in the TEM
is a challenging endeavor, made even more
difficult by the highly oxidizing surface of
a freshly prepared Pu specimen. A highly
refined specimen-preparation procedure,
coupled with an excellent glovebox at-
mosphere, essentially eliminates the intro-
duction of stress, temperature, and reactive
Stage III recovery associated with va-
cancy migration occurs at about 190 K in
both the proton- and self-irradiated speci-
mens and is also captured in the KMC
simulations. Annealing recovery is essen-
tially complete for the proton-irradiated
specimen at about 320 K, while a slight
shift to higher temperatures is observed
for the self-damaged specimen. The shift
of Stage V to higher recovery annealing
temperatures is consistent with a popula-
Figure 2. The isochronal annealing recovery of a Pu-Ga alloy following low-temperature
ꢀ20 K) irradiation with 3.8-MeV protons (green and red, representing two distinct
experimental measurements) and the self-damage from ꢀ decay (black). (a) Experimentally
measured change in resistance; (b) kinetic Monte Carlo simulation results of the change in
defect density, along with the annealing Stages I, II, III, and V (Stage IV not shown).
(
Figure 3. Model predictions of He
bubble number density and mean
diameter as a function of age.
MRS BULLETIN/SEPTEMBER 2001
681

Fundamental Studies of Plutonium Aging
gases, as evident by the atomic-resolution
image of ꢁ-Pu-Ga along a [001] zone axis
shown in Figure 4a.
Phase Stability
optical-light micrograph of ꢀꢄ martensite
laths formed in the ꢁ-phase matrix after
cooling to ꢉ118ꢂC and holding for 1000 s.
After this low-temperature excursion, the
average size of the martensite precipitates
is 20 ꢆm long by 2 ꢆm wide. Notably, we
observe both increasing martensite size
and increasing number fraction with in-
creasing time at temperature, thus con-
firming the isothermal nature of this
transformation. A number of variants are
observed within each grain and help to
reveal the crystallographic nature of the
displacive phase transformation. The TEM
image in Figure 5b shows two such ꢀꢄ par-
ticles imbedded in the ꢁ matrix. An in-
creased dislocation density is clearly visible
at the ꢀꢄ–ꢁ matrix interface, indicating that
coherency strains and plastic deformation
are important components of the accom-
modation of the large volume change.
The stability of ꢁ-phase Pu-Ga alloys
during aging is also a fundamental con-
cern, especially considering the evolving
microstructure resulting from radiation
decay and transmutants. The nondiffusive,
or displacive, martensitic phase transfor-
mation remains among the most challeng-
ing problems to understand in physical
metallurgy, and Pu-Ga alloys present one
of the more extreme examples. Here, the
ꢁ-to-ꢀꢄ phase transformation results in a
density change of approximately 20%, en-
suring that stress and plastic accommo-
dation play central roles in governing the
details of the transformation. Microstruc-
tural characterization and determination
of the phase-transformation hysteresis by
electrical resistometry will provide key
parameters required to model the influ-
ence of microstructural changes on the
hysteresis width. Our approach is to bal-
ance the thermodynamic driving force
We have characterized large, ꢀ35-nm
He bubbles in a 35-yr-old Pu-Ga alloy that
was annealed at 400ꢂC to coarsen the bubble
microstructure, as shown in Figure 4b. A
uniform dispersion of He bubbles is ob-
served, many of which appear attached to
one or more dislocations. Calculations are
currently under way to determine the He
pressure within the bubbles through eval-
uation of the dislocation curvature. Analy-
sis of numerous micrographs provides a
13
3
bubble number density of 6.5 ꢈ 10 /cm
with an average diameter of 35 nm and a
volume fraction of ꢀ0.15%. The influence
of age, and thus He content, on bubbles is
evident by the smaller bubble populations
observed in annealed Pu-Ga alloys of
1
7-yr-old material.31,32
Summary and Outlook
(Gibbs free energies of the phases calcu-
A coupling of theory, modeling, and ex-
periments is under way to elucidate the
fundamental mechanisms of radiation
damage and phase stability in plutonium.
Radioactive decay of Pu results in signifi-
cant local lattice damage, while at the
same time producing the ingrowth of
insoluble He and impurities that likely
stabilize (Am) and destabilize (U, Np) the
fcc ꢁ phase. The combination of lattice
damage and transmutation results in a
microstructure that evolves with age. A
combination of atomistic simulations, low-
temperature irradiation, and isochronal
annealing recovery experiments is under
way to develop a detailed understanding
of primary damage production, defect
clustering, and transport kinetics, which is
key to understanding and predicting ra-
diation damage and its attendant conse-
quences to microstructures and properties.
The spatial character of the high-energy
lated from the phase diagram) with the
mechanical properties of the ꢁ matrix (the
retarding force).
Both new and naturally aged alloys
have been cooled to various temperatures
to evaluate the influence of age on the
transformation temperature and the mar-
tensite volume fraction. Figure 5a is an
Figure 4. (a) Atomic-resolution
transmission electron microscope
(
TEM) image of a ꢁ-phase Pu alloy
along the [001] zone axis. The shortest
distance between white dots is 0.23 nm.
(
b) Bright-field TEM image revealing
helium bubbles and the associated
dislocation structure.
Figure 5. (a) Optical-light micrograph of a ꢁ-phase Pu alloy cooled to ꢉ118ꢂC to induce the
ꢁ-to-ꢀꢄ phase transformation. (b) Bright-field TEM image of two small ꢀꢄparticles in the ꢁ matrix.
682
MRS BULLETIN/SEPTEMBER 2001

Fundamental Studies of Plutonium Aging
U-recoil displacement cascades is evident
in the low-temperature (Stages I and II)
annealing recovery stages and a shift of
Stage V recovery to higher temperatures
than for proton irradiation. Finally, the
characterization of aged Pu-Ga alloys is
under way to identify small, ꢀ1-nm He
bubbles predicted by models and to quan-
tify the effect of age on ꢁ-phase stability.
Massalski, H. Okamoto, P.R. Subramanian, and
L. Kacprzak (ASM International, Materials
Park, OH, 1990) p. 1843.
Houston, 1990) p. 1.
20. G.R. Odette, B.D. Wirth, D.J. Bacon, and
N.M. Ghoniem, MRS Bull. 26 (3) (2001) p. 176.
21. J.F. Ziegler, J.P. Biersack, and U. Littmark,
The Stopping and Range of Ions in Solids, Vol. 1,
edited by J.F. Ziegler (Pergamon Press, New
York, 1985).
8
. S.S. Hecker and L.F. Timofeeva, in Los Alamos
Science, No. 26, edited by N.G. Cooper (Los
Alamos National Laboratory, Los Alamos, NM,
2
000) p. 244.
9
. O.J. Wick, ed., Plutonium Handbook: A guide
22. R.S. Averback and T. Diaz de la Rubia, in
Solid State Physics, Vol. 51, edited by F. Spaepen
and H. Ehrenreich (Academic Press, New York,
1997) p. 281.
to the Technology (American Nuclear Society,
La Grange Park, IL, 1980).
10. F.H. Ellinger, C.C. Land, and V.O. Struebing,
Acknowledgments
J. Nucl. Mater. 12 (1964) p. 226.
1. B. Hocheid, A. Tanon, S. Bedere, J. Depres,
S. Hay, and F. Miar, in Proc. 3rd Int. Conf. on Plu-
tonium, edited by A.I. Kay and M.B. Waldron
23. D.J. Bacon, J. Nucl. Mater. 251 (1997) p. 1.
24. R.E. Stoller, JOM 48 (1996) p. 43.
1
This work was performed under the
auspices of the U.S. Department of Energy
by the University of California, Lawrence
Livermore National Laboratory, under
contract No. W-7405-Eng-48.
25. M.J. Norgett, M.T. Robinson, and I.M.
Torrens, Nucl. Eng. Des. 33 (1975) p. 50.
26. P. Ehrart, in Atomic Defects in Metals,
Landolt–Bornstein New Series Group III/25, edited
by H. Ullmaier (Springer-Verlag, Berlin, 1991).
27. D.A. Wigley, Proc. R. Soc. London, Ser. A 284
(1965) p. 344.
(
Chapman & Hall, London, 1965) p. 321.
1
2. N.T. Chebotarev, E.S. Smotriskaya, M.A.
Andrianov, and O.E. Kostyuk, in Proc. 5th Int.
Conf. on Plutonium and Other Actinides, edited by
H. Blank and R. Lindner (North-Holland, New
York, 1975) p. 37.
13. P.H. Adler, Metal. Trans. A 22A (1991) p. 2237.
14. L.F. Timofeeva, in Proc. Int. Conf. on Ageing
Studies and Lifetime Extension of Materials, edited
by L.G. Mallinson (Kluwer Academic Publish-
ers, Dordrecht, 2000) p. 191.
15. W.G. Wolfer, in Los Alamos Science, No. 26,
edited by N.G. Cooper (Los Alamos National
Laboratory, Los Alamos, NM, 2000) p. 274.
16. G.R. Odette, J. Nucl. Mater. 85/86 (1979) p. 533.
17. L.K. Mansur and E.E. Bloom, J. Met. 34
(1982) p. 23.
18. B.N. Singh and A.J.E. Foreman, J. Nucl.
Mater. 122/123 (1984) p. 537.
19. P.L. Andresen, F.P. Ford, S.M. Murphy, and
J.M. Perks, in Proc. 4th Int. Symp. on Environ-
mental Degradation of Materials in Nuclear Power
Systems—Water Reactors (NACE International,
References
1
2
3
. S.D. Drell, Phys. Today 12 (2000) p. 25.
. R. Jeanloz, Phys. Today 12 (2000) p. 44.
. S.S. Hecker and J.C. Martz, in Los Alamos Sci-
28. J.B. Adams and W.G. Wolfer, J. Nucl. Mater.
166 (1989) p. 235.
29. C. Colmenares, R.H. Howell, D. Ancheta, T.
Cowen, J. Hanafee, and P. Sterne, First Positron
Annihilation Lifetime Measurement of Pu, Report
No. UCRL-ID-126003 (Lawrence Livermore
National Laboratory, Livermore, CA, 1996).
30. R.H. Howell (personal communication).
31. D.L. Rohr, K.P. Staudhammer, and K.A.
Johnson, Development of Plutonium Transmission
Electron Microscopy, Report No. LA-9965-MS
(Los Alamos National Laboratory, Los Alamos,
NM, 1984).
ence, No. 26, edited by N.G. Cooper (Los Alamos
National Laboratory, Los Alamos, NM, 2000)
p. 238.
4
. A.M. Boring and J.L. Smith, in Los Alamos
Science, No. 26, edited by N.G. Cooper (Los
Alamos National Laboratory, Los Alamos, NM,
2
000) p. 90.
5
. S.S. Hecker, in Los Alamos Science, No. 26, ed-
ited by N.G. Cooper (Los Alamos National Lab-
oratory, Los Alamos, NM, 2000) p. 290.
6
. For example, see Los Alamos Science, No. 26,
32. T.G. Zocco and D.L. Rohr, in Specimen Prepa-
ration for Transmission Electron Microscopy of Ma-
terials, edited by J.C. Bravman, R.M. Anderson,
and M.L. McDonald (Mater. Res. Soc. Symp.
edited by N.G. Cooper (Los Alamos National
Laboratory, Los Alamos, NM, 2000).
7
. D.E. Peterson and M.E. Kassner, in Binary
Alloy Phase Diagrams, 2nd ed., edited by T.B.
Proc. 115, Pittsburgh, 1988) p. 259.
ꢀ
Don't Miss These Symposium Proceedings
from the Materials Research Society
Scientific Basis for Nuclear Waste
Management XXIII
Volume 608-B
Microstructural Processes in
Irradiated Materials
Volume 540-B
$
$
$
83.00 MRS Member
95.00 U.S. List
110.00 Non-U.S.
$ 80.00 MRS Member
$ 92.00 U.S. List
$100.00 Non-U.S.
Scientific Basis for Nuclear Waste
Management XXII
Volume 556-B
Structure and Dynamics of
Glasses and Glass Formers
Volume 455-B
$
$
$
78.00 MRS Member
89.00 U.S. List
98.00 Non-U.S.
$ 65.00 MRS Member
$ 75.00 U.S. List
$ 86.00 Non-U.S.
For more information, or to order any of the proceedings volumes listed above,
contact the MRS Customer Services Department.
TEL
FAX
724-779-3003
724-779-8313
E-Mail info@mrs.org • www.mrs.org/publications/books/
www.mrs.org/publications/bulletin
MRS BULLETIN/SEPTEMBER 2001
683