4.06

Radiation Effects in Refractory Metals and Alloys

K. J. Leonard
Oak Ridge National Laboratory, Oak Ridge, TN, USA

Published by Elsevier Ltd.

4.06.1

Introduction

181

4.06.2
4.06.2.1
4.06.2.2
4.06.2.3
4.06.3
4.06.3.1
4.06.3.2
4.06.3.3
4.06.4
4.06.4.1
4.06.4.2

Niobium and Nb-Base Alloys
Introduction and History of Nb and Nb Alloys
Radiation-Induced Swelling of Nb and Nb-Base Alloys
Mechanical Properties of Irradiated Nb and Nb Alloys

Tantalum and Ta-Base Alloys
Introduction and History of Ta and Ta Alloys
Irradiation-Induced Swelling of Ta and Ta-Base Alloys
Mechanical Properties of Irradiated Ta and Ta-Base Alloys
Molybdenum and Mo-Base Alloys
Introduction and History of Mo and Mo Alloys
Irradiation-Induced Swelling and Physical Property Changes in Mo and
Mo-Base Alloys
Mechanical Properties of Irradiated Mo and Mo Alloys
Tungsten and W-Base Alloys
Introduction and Irradiated Properties Database for W and W Alloys
Irradiation-Induced Swelling and Physical Property Changes in W and W Alloys
Irradiated Mechanical Properties of W and W Alloys
Outlook

182
182
183
185
188
188
189
190
194
194

4.06.4.3
4.06.5
4.06.5.1
4.06.5.2

4.06.5.3
4.06.6
References

Abbreviations

Symbols

bcc
C-103
Cb-752
DBTT

D
n
T
Tirr
Tm
a
DV/V
w

Body-centered cubic
Nb–10Hf–1Ti alloy
Nb–10W–2.5Zr alloy
Ductile–brittle transition
temperature
FS-85
Nb–10W–28Ta–1Zr alloy
HP

High purity
JIMO
Jupiter icy moons orbiter
LCAC
Low-carbon arc cast
NERVA
Nuclear experiment for rocket vehicle
applications
ODS
Oxide dispersion strengthened
RIS
Radiation-induced
segregation
SNAP
Systems nuclear auxiliary
power
T-111
Ta–8W–2Hf alloy
TZM
Mo–0.5Ti–0.1Zr alloy
UTS
Ultimate tensile strengths
UWMAK-III University of Wisconsin Madison
fusion reactor concept

194
197
206
206
206

207
209
211

Thermoelectric power
Neutron particle
Temperature
Irradiation temperature
Melting temperature
Alpha particle
Volume fraction swelling
Fluence

4.06.1 Introduction
Refractory metals and alloys offer attractive and
promising high-temperature properties, including
high-temperature strength, good thermal conductivity, and compatibility with most liquid metal coolants, many of which are suitable for applications in
nuclear environments. Though many of the refractory alloys have been known for more than 60 years,
there are significant gaps in the materials property
database for both unirradiated and irradiated
181


182

Radiation Effects in Refractory Metals and Alloys

conditions. In addition, significant issues related
to low-temperature irradiated mechanical property
degradation at even low neutron fluences restrict the

use of refractory metals. Protection from oxidizing
environments also restricts their use, unless suitable
protection or a liquid metal coolants is used.
Much of the early research on refractory metal
alloys was centered on applications in aerospace as
well as cladding and structural materials for fission
reactors, with particular emphasis on space reactor
applications. Reviews concerning the history of
these programs and the development of many of
the alloys whose irradiated properties are discussed
in this chapter can be found elsewhere.1–5 Due to
cancellations and reintroduction of new mission criteria for these space reactor programs, the materials
database shows similar waves in the gains of intellectual knowledge regarding refractory alloy and irradiated property behavior. Unfortunately, as seen in
the subsequent sections of this chapter, much of the
irradiated property database for refractory metals consists of scoping examinations that show little overlap
in either material type, metallurgical conditions (i.e.,
grain size, impurity concentrations, thermomechanical
treatments), radiation conditions (i.e., spectra, dose
and temperature), or postirradiation test conditions
or methods.
The irradiation behavior of body-centered cubic
(bcc) materials is known. Irradiation-induced swelling
because of void formation in the material lattice is
typical for temperatures between 0.3 and 0.6 Tm,
where Tm is the melting temperature. Maximum
swelling in refractory metals is <10% for displacement damage levels up to 50 dpa (displacements per
atom), but typical values for fission reactor applications are <4%. Alloy additions can further reduce the
sensitivity to swelling, for example, rhenium additions
to molybdenum or tungsten. These levels of swelling
are manageable through the appropriate engineering

design of components.
The generation of dislocation loops and point
defects provide significant irradiation-induced
strengthening or hardening of refractory metals and
alloys. This in turn creates reductions in the ductility
and fracture toughness of the material. This is most
pronounced at temperatures <0.3 Tm, where defect
mobility is reduced. The increase in the yield
strength of the material because of the irradiationinduced defects can exceed the fracture strength
of the material, leading to brittle behavior. These
degradations in material property can begin to
occur at neutron fluences as low as 1 Â 1020 n cmÀ2,

or $0.03 dpa3 and increase in severity with dose.
As irradiation temperatures increase, dislocation loop
and void sizes increase, whereas their densities are
reduced, providing improvements in ductility, though
at a reduced strength of the material. At high enough
temperatures, recovery of properties to levels close to
that of the unirradiated values is possible, though
changes in material properties may be further influenced by microstructural changes such as segregation
or precipitate formation of solute and transmuted species or recrystallization, which can lead to further deterioration of properties. Detailed information on the
effects of radiation on materials is presented in Chapter
1.03, Radiation-Induced Effects on Microstructure,
and in Chapter 1.04, Effect of Radiation on Strength
and Ductility of Metals and Alloys. In general, the use
of refractory alloys in radiation environments is not
recommended at temperatures <0.3 Tm. However,
new research work, particularly on molybdenum
and its alloys, has shown that control over interstitial

element contamination levels, grain size, and morphology, as well as the introduction of oxide dispersion
strengthening, can lead to improvements in the lowtemperature irradiation behavior. This is discussed in
detail in this chapter.
The following sections of this chapter deal with
the irradiated properties database of niobium, tantalum, molybdenum, and tungsten, as well as their
alloys. While vanadium may sometimes be considered a refractory metal, its melting temperature is
considerably lower than that of the other materials
mentioned. However, its radiation effects database
is considerable and well advanced relative to some
refractory metals and it is therefore discussed separately in Chapter 4.12, Vanadium for Nuclear Systems. The irradiated properties database for refractory
alloys is particularly thin, especially involving fracture toughness properties, irradiation creep effects,
and combined radiation effects with high thermomechanical loads such as those experienced in plasma
facing components or spallation target materials.
Where needed, a comparison of the unirradiated and
irradiated properties of a material is given.

4.06.2 Niobium and Nb-Base Alloys
4.06.2.1 Introduction and History of Nb
and Nb Alloys
The push for higher operating temperatures in turbine engines, as well as in reactor designs for both
terrestrial and space applications, has frequently


Radiation Effects in Refractory Metals and Alloys

governed the periodic scientific examinations of
refractory alloys. The historical examination of Nb
and its alloys is typical of this, with early studies of
the irradiation properties exploring the potential uses
of these alloys in fusion energy and fission type space

reactor. While these alloys have favorable properties,
such as elevated temperature capability and compatibility with liquid alkali metals for energy applications,
and attractive physical properties such as thermal
conductivity, much of the work on Nb and Nb-base
alloys has examined the nonirradiated properties.
It is worth putting into perspective the relatively
small commercial market for niobium-base alloys.
Approximately, 75% of all niobium metal is used
as minor alloying additions in steel, and only 1–2%
is produced in the form of niobium-base alloys.
The total market for niobium-base alloys in the
mid-1990s was <105 kg yearÀ1 (100 metric tonnes
yearÀ1).6 By comparison, 60 tons of niobium was
used in 1961 for the SNAP-50 reactor program
alone, and substantial additional quantities were
used for other research projects such as the NERVA
(Nuclear Experiment for Rocket Vehicle Applications) program.7 Throughout the history of the various space reactor programs, dozens of alloys were
examined, with several brought to near-commercial
production. However, today only the Nb–1Zr and
C-103 (Nb–10Hf–1Ti) alloys remain commercially
available for use in the sodium vapor lamp and rocket
or turbine engine exhaust nozzles.
Nb–1Zr has historically been considered the
only niobium-base alloy with a sufficiently mature
database (mechanical properties including thermal
creep, chemical compatibility, fabrication, and welding knowledge) to be considered a near-term candidate for radiation environments.7–11 Developed for
high ductility and good weld characteristics, the alloy
shows less-than-desirable thermal creep strength at
elevated temperatures compared to other refractory
alloys. Though the C-103 alloy has greater short-term

elevated temperature strength than that of Nb–1Zr,
its long-term properties show no improvement over
Nb–1Zr.12 Nonetheless, Nb–1Zr is the only Nb-base
alloy with a significant radiation effects database.
Despite the periodic programmatic interest in the
use of Nb and Nb-base alloys, no clear fundamental
study of the irradiated properties for a specific application has been performed or completed. Much of the
data available on the irradiated properties is scattered
and easily spans a time frame of several decades,
which can lead to misinterpretations of results on
the basis of either the limited scientific knowledge

183

of the time, lack of understanding of the sensitivity
of properties on impurity concentrations, or aging
effects. Radiation effects data are limited to the
examination of swelling and tensile properties, with
no information regarding fracture toughness or irradiation creep performance.
The following sections deal with radiation effects
on the properties of Nb and Nb–1Zr specifically.
While some initial scoping examinations have
been performed on other Nb-base alloys, these are
relatively inconsequential and based on the less-thandesirable ductility, thermal stability, or welding capabilities of these alloys.
4.06.2.2 Radiation-Induced Swelling of Nb
and Nb-Base Alloys
Like all group V transition metals, the affinity of Nb
for, and its ability to dissolve, high concentrations of
interstitial atoms such as hydrogen, oxygen, nitrogen,
and to a lesser extent, carbon can strongly influence

the properties of the metal through defect–impurity
interactions. Hydrogen, carbon, and oxygen impurities
have a strong effect on the tensile Ductile–brittle
transition temperature (DBTT) of pure Nb (reviewed
by Hahn et al.13), with hydrogen levels near 10 ppm
increasing the DBTT to 173 K and over 273 K at
levels >100 ppm (DBTT of high purity Nb and Nb
alloys is near 73 K14,15). The effects of oxygen and
carbon were less severe, but influential at levels
of 100 ppm and greater. The effect of nitrogen on
embrittlement also appears to be as severe as that of
oxygen, though some uncertainty exists as to whether
solubility limits have been exceeded in the data.1
The effect of interstitial impurities on the irradiated
properties of Nb and Nb-base alloys is significant
and has been examined, though the overall database
for irradiated properties is limited.
The interplay between the radiation-created defects
and the interstitial impurity elements was investigated by Igata et al.16 for pure Nb (70 wppm oxygen
and 30 wppm nitrogen) irradiated to 3.4 Â 1020 n cmÀ2
(E > 1 MeV) at temperatures below 413 K and postirradiation annealed up to 973 K. Increases in yield
strength over the as-irradiated values following
annealing were measured at 473 and 673 K, attributed to the interplay of the defect clusters trapping
oxygen and nitrogen atoms, respectively. Above 773 K,
no difference between the annealed and as-irradiated
yield stress was observed.
Hautoja¨rvi et al.17 and Naidu et al.18 examined the
interaction between vacancies and interstitial



Radiation Effects in Refractory Metals and Alloys

impurities in irradiated Nb through positron annihilation studies. In high-purity Nb, vacancy clustering
within the collision cascades is observed, starting as
low as 160 K, with vacancy migration peaking around
250 K, but in materials with higher hydrogen content,
the vacancy migration stage shifts to temperatures
close to 400 K.17 The irradiation exposure at intermediate temperatures (0.3–0.6 Tm) can lead to void
swelling, irradiation creep, and helium embrittlement
through processes involved in (n,a) reactions or
impurity gas atoms. Naidu et al.18 examined the effect
of He and its interaction with vacancies in pure Nb,
leading to the development of bubbles through
a-irradiated specimens. At temperatures between
623 and 1023 K, bubble growth occurs through the
addition of He atoms and vacancies, followed by
migration and coalescence at higher temperatures,
eventually leading to the annealing out of the He
bubbles and vacancy complexes above 1173 K.18
The irradiation-induced swelling of pure Nb generally appears at temperatures between 673 and 1323 K
with peak swelling near 873 K (0.32 Tm), though these
limits are not clearly defined and are based on the very
limited data available, compiled by Wiffen19 and
Pionke and Davis.1 A maximum swelling of 4.8% following irradiation to 2.5 Â 1022 n cmÀ2 at 858 K was
reported.1 However, the magnitude of swelling shows
considerable scatter in the literature, possibly reflecting the influence of impurity concentrations and differences in irradiation conditions and microstructural
interpretation of the materials.19 Fischer20 reported
that void concentration increased four to seven times
for a fourfold increase in flux for the same total fluence.
This produced a reduction in void size with flux and

therefore a reduction in the total swelling.
Loomis and Gerber21–23 examined the influence
of oxygen and substitutional binary alloy additions on
the swelling of 3 MeV 58Ni+ ion-irradiated Nb up
to $50 dpa. Void formation and characteristics in size
and morphology were found to be dependent on
temperature, oxygen concentration, and the type of
substitutional alloy addition. The average void diameter was found to increase with temperature as well as
oxygen up to 0.02 at.%. Higher oxygen concentrations resulted in a decrease in void diameter to
0.1 at.% O, above which void diameters showed no
significant changes. The number density of voids was
found to decrease with temperature, but increase
with oxygen concentration to $0.06 at.%, above which
the number density showed no significant change.
As the volume fraction of swelling (DV/V) is proportional to both the void number and the cube of the

void diameter, the volume fraction is observed to
increase with temperature and oxygen concentration
to $0.04 at.%, followed by a decrease and plateau of
the volume fraction above 0.1 at.%. The dependence
of DV/V on temperature and oxygen concentration
is illustrated in Figure 1. Microstructural examination revealed an ordering of the voids into a latticetype structure in the material irradiated at 1050 K
to $40 dpa and oxygen concentration !0.039 at.%
oxygen. The higher temperature of the maximum
swelling as compared to the neutron irradiation data
is believed to be associated with the higher displacement damage rate of the ion-bombarded material,19
though the higher impurity levels may also provide
an influence.
The effect of dilute ($2.4 at.%) substitutional
alloy addition on the swelling of 0.06 at.% oxygendoped Nb was also examined for 3 MeV 58Ni+ ion

irradiation at 1225 K. The DV/V was determined to
increase through the addition of Ta, but decreased
with increasing effectiveness by the addition of Ti,
Zr, V, and Hf. The addition of the reactive alloying
elements to Nb suppresses void formation through
the gettering of interstitial impurities that act as void
nucleation sites. The DV/V was determined to be
unaffected by the addition of Ni or Fe. The dependence of DV/V on temperature, oxygen, and substitutional addition is also shown in Figure 1.

14

30 MeV 58Ni+ irradiation
Tirr =1225 K, 50 dpa

12
10
DV/V (%)

184

8
1240 K

6

1225 K

4

1240

1200 K

2
0
0.00 0.04

0.08 0.12 0.16 0.20
Oxygen concentration (at.%)

Nb + 2.4 at.% Ta
Nb + 2.4 at.% Ni
Nb + 2.3 at.% Fe
Nb + 2.3 at.% Ti

Tirr (K)

1220
1200

Nb + 2.4 at.% Zr
Nb + 2.4 at.% V
Nb + 2.4 at.% Hf
Nb + 2.4 at.% Mo

Figure 1 The dependence of void volume fraction (DV/V )
in 3 MeV 58Ni+ ion-irradiated Nb on the concentration of
oxygen and dilute solute additions. Reproduced from
Loomis, B. A.; Gerber, S. B. J. Nucl. Mater. 1983, 17,
224–233.



Radiation Effects in Refractory Metals and Alloys

Swelling in Nb–1Zr has been examined, though
only scattered data are available in the examination of
temperature and flux dependence. The available
swelling data on Nb–1Zr, compiled by Powell et al.24
and Watanabe et al.25 presented in Figure 2, show the
lack of data on the temperature range in which peak
swelling appears. The swelling data shown in the
figure were measured through electron microscopy,
with the exception of the data by Powell et al.24 and
Wiffen.26 Alloy impurity chemistry, in addition to
interpretation and measurement error, may account
for the scatter associated with the lower temperatures. The work of Watanabe et al.25 and Garner
et al.27 indicates that irradiation-induced swelling is
dependent on the thermomechanical history of the
material. In that material, cold-working followed by
solution anneal and aging exhibited swelling, while
material not given the preirradiated cold-working
showed some densification. The changes in density
of the material are dependent on the phase-related
transformations involving precipitation.
Swelling in Nb–1Zr appears to be centered over a
more narrow temperature range than in Nb, with a
peak near 1073 K that is higher than that of the pure
metal. While the addition of Zr to Nb appears to delay
nucleation of voids to higher temperatures, the voids
that form are of larger size than those appearing in
pure Nb under comparable conditions. For example,


2.5

DV/V (%)

2.0
1.5
1.0
0.5
Irr 400
ad 600
iat
ion 800
te 1000
m
pe 1200
ra
tu 1400
re
(K
)

20
0

Powell et al.24
Jang and Moteff145
Sprague et al.146
Wiffen28


40
dpa

60

80

100

Wiffen26
Michel and Smith147
Watanabe et al.25
Garner et al.27

Figure 2 Swelling as a function of irradiation temperature
and dose for neutron-irradiated Nb–1Zr from available
literature compiled by Powell et al.24 and Watanabe et al.25

185

following irradiation to 2.5 Â 1022 n cmÀ2 (E > 0.1
MeV) at 1063 K, the diameter, concentration, and
volume fraction of voids in Nb–1Zr was 57.5 nm,
1.8 Â 1020 mÀ3, and 2.2%, respectively,1 whereas
under similar conditions, the same void parameters in
pure Nb were 18.6 nm, 2.8 Â 1021 mÀ3, and 1.04%.
While void formation and swelling in Nb and
Nb–1Zr occurs, the total swelling is generally <5%
and within engineering limits, even for high neutron
exposures >10 dpa.3 The addition of Ti to Nb was

found to increase void resistance and has been found
to suppress void formation in V at concentrations as
low as 3%.29 The combination of reactive alloy elements and Nb in the C-103 alloy may suggest a
greater void formation resistance than in pure Nb
and Nb–1Zr.
4.06.2.3 Mechanical Properties of
Irradiated Nb and Nb Alloys
Little coverage of the changes in mechanical properties following irradiation has been given to Nb and
Nb alloys, with the majority of the data for temperatures below 800 K. Some preliminary experimental
work on the irradiated mechanical properties of Nb
alloys Cb-752 (Nb–10W–2.5Zr)30 and FS-85 (Nb–
10W–28Ta–1Zr)31 is available. However, these alloys
are not commercially produced and have shown
indications of thermal aging instabilities, leading to
grain boundary embrittlement.12,32,33 The irradiated
mechanical properties of these alloys show similar
radiation hardening as in the pure metal, but with
mechanical properties more sensitive to thermal
aging conditions. The bulk of the irradiated mechanical properties data is for the Nb–1Zr alloy as well as
the pure metal, and is covered in this review.
The irradiated mechanical properties of Nb and
Nb–1Zr are strongly governed by irradiation temperature, determining whether the mechanical properties
are controlled by dislocation loops or a combination of
loops and voids in the microstructure. As cavity formation can be delayed or suppressed by higher irradiation temperatures in Nb–1Zr, mechanical property
comparisons between the alloy and the base metal will
reflect their irradiated microstructure. For Nb and
Nb–1Zr irradiated to 3 Â 1022 n cmÀ2 at $728 K, the
pure metal contains both dislocation loops and voids,
while the alloy exhibits no void formation.19 A comparison of the tensile properties of Nb and Nb–1Zr irradiated under similar conditions is shown in Figure 3.
The irradiated strength of both materials shows an

increase in tensile strength above the unirradiated


Radiation Effects in Refractory Metals and Alloys

120

Niobium

Nb–1Zr

Stress (1000 psi)

100
80
60
40
20

Elongation (%)

0
60

Control
3.7 ϫ 1022 n cm–2
At 450 ЊC

Control
3.0 ϫ 1022 n cm–2

At 460 ЊC
Ultimate

800

600

Yield and
ultimate

Yield

400
Ultimate

Ultimate

Yield

Yield

Stress (MPa)

186

200

0
Total


Total

40
20
0
0

Uniform

Uniform
Total

Total
Uniform
200

400

600
0
200
Test temperature (ЊC)

400

600

Figure 3 Comparison of tensile properties between Nb and Nb–1Zr tested under similar irradiation conditions.
Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space Nuclear Power Applications, CONF-8308130;
Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak Ridge, TN, 1984; pp 252–277.


800

Niobium
Ttest = 298 K
3.0 ϫ 1026 n cm–2
Tirr = 733 K

600
Stress (MPa)

condition, with Nb–1Zr showing a greater sensitivity
to irradiation. As the mechanical properties of Nb–1Zr
are dominated by the dislocation loop structures, yield
instability is observed in the material, leading to the
early onset of necking. This results in <0.2% uniform
elongation, though total elongation near 10% is still
achieved. The yield instability is associated with dislocation channeling, in which deformation dislocations
will create defect-free channels along their slip plane,
following the annihilation of the loop structures. This
occurs only after enough applied stress is achieved to
overcome the obstacles, but the net effect is a nonuniform plastic deformation through channels that allow
for the movement of deformation dislocations at
reduced stress.
The irradiated Nb samples whose properties are
shown in Figure 3 contain, in addition to dislocation
loops, voids that limit dislocation channeling by
providing added obstacles to deformation, resulting
in some measure of uniform elongation and work
hardening upon yielding. The microstructure dependence on the tensile properties can best be illustrated

by the comparison shown in Figure 4 of Nb irradiated
at 328 and 733 K. The higher irradiation temperature
results in the development of microstructural voids
and thus the significant differences in the tensile
curves. The lower irradiation temperature results in
dislocation channeling following yield and the

400

7.5 ϫ 1024 n cm–2
Tirr = 328 K

200

0
0

4

12
8
Elongation (%)

16

Figure 4 Comparison of tensile curves between Nb
irradiated at 328 and 733 K. Yield instability is seen at 328 K
due to channeling of deformation dislocations through the
irradiated dislocation loop structures. The higher irradiation
temperature resulted in the development of small voids

providing a barrier to dislocation movement. Reproduced
from Wiffen, F. W. In Refractory Alloy Technology for
Space Nuclear Power Applications, CONF-8308130;
Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National
Laboratory: Oak Ridge, TN, 1984; pp 252–277.

associated work softening during necking to failure at
around 11% total elongation. While the higher irradiation temperature sample was irradiated to a higher
total fluence, the effect of dose is observed only on the


Radiation Effects in Refractory Metals and Alloys

relative strength increase over the unirradiated condition. The higher irradiation temperature produced
voids in the microstructure, providing additional obstacles to deformation and higher uniform elongations
and modest work hardening.
Little is known with regard to the aging properties
of Nb–1Zr or the combined thermal and radiation
effects. The addition of 1 wt% Zr to Nb creates a
dispersion-strengthened alloy, in which the Zr combines with interstitial impurities creating fine precipitates throughout the material. The development of
these fine precipitates on aging at 1098 K can increase
the tensile strength between 50 and 100 MPa over the
annealed condition and provide an effective strengthening greater than that observed through modest
irradiation31 (Table 1).
Irradiation of Nb–1Zr to 0.9 dpa at 1098 K showed
a modest increase in yield and ultimate tensile
strength to 135 and 192 MPa, respectively, over the
annealed condition. This increase in tensile strength
either through aging or irradiation results in a corresponding decrease in uniform elongation from 15%
to 3.5% and total elongation from 25% to 15%.

Aging at temperatures above 1098 K produces little
effective hardening as the precipitates coarsen in the
microstructure.33 Irradiation to 0.9 dpa at 1248 and
1398 K of Nb–1Zr showed only a modest increase
in the yield strength over the aged and annealed
specimens, though ultimate tensile strength and elongation were unchanged or less. Irradiation to 1.88 dpa
at 1223 K resulted in weaker tensile properties
Table 1
of Nb–1Zr

187

compared to the 0.9 dpa sample, believed to be due
to further precipitate coarsening. The time under
irradiation conditions for the 1.88 dpa sample was
near 1100 h and produced similar tensile properties
as that of the aged-only material.
As discussed in the preceding paragraphs, the
irradiated properties of Nb and Nb–1Zr are governed
by their microstructure and are influenced by temperature, displacement damage rate, and neutron spectrum. The tensile properties of neutron-irradiated
Nb–1Zr for damage levels between 0.1 and 5 dpa
(Horak et al.34 and Wiffen35) summarized by Zinkle
and Wiffen3 are shown in Figure 5. At temperatures
below 800 K, a large increase in the tensile strength
from irradiation is observed with the corresponding
low uniform elongations. At higher temperatures,
uniform elongation increases because of the presence
of voids in the microstructure. However, the data
plotted in Figure 5 show uniform elongations remaining low up to 1100 K, while radiation hardening is
relatively moderate, suggesting that impurities are

the source of the reduced elongation values.
No irradiated fracture toughness data exist for Nb
or Nb–1Zr, though comparisons can be made from the
larger irradiated vanadium alloy database, in which
fracture toughness embrittlement becomes a concern
when tensile strength exceeds 600–700 MPa and therefore at temperatures below 400 K for Nb–1Zr.36 However, if a conservative value is assigned to the critical
stress to induce cleavage fracture of $400 MPa (40%
lower than that observed in vanadium alloys),

Tensile property comparison illustrating the effects of aging and irradiation on the mechanical properties

Test/aged/irradiated
temperature (K)

Yield strength
(MPa)

As-annealed condition
298
185.00
1073
102.33
1223
89.00
1373
83.00
1100 h aged
1073
185.67
1223

111.67
1373
82.50
Irradiated: 2.04 Â 1021 n cmÀ2 (E > 0.1 MeV), 0.93 dpa
1073
134.50
1223
149.50
1373
102.00
Irradiated: 4.13 Â 1021 n cmÀ2 (E > 0.1 MeV), 1.88 dpa
1223
104.50

Ultimate tensile strength
(MPa)

Uniform
elongation (%)

Total elongation
(%)

281.00
191.00
215.00
157.00

18.60
16.03

15.05
10.50

34.60
23.90
23.70
44.80

248.67
156.67
130.50

8.07
8.07
6.80

13.27
13.27
28.65

192.50
182.00
127.50

7.10
3.55
9.55

18.55
17.80

33.65

166.00

8.10

20.90

Source: Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/05-38;
Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005.


Radiation Effects in Refractory Metals and Alloys

Ultimate tensile strength (MPa)

700

35
Irrad. UTS

600

30
Irrad. eU
25

500

20


400
300

Unirrad. UTS

15

200

10

100

5

0
200

400

600

800
1000 1200
Temperature (K)

1400

Uniform elongation (%)


188

0
1600

Figure 5 Unirradiated and irradiated (0.1–5 dpa) ultimate tensile strengths (UTS) and uniform elongation (eU) of Nb–1Zr.
Irradiated data represented by solid symbols and unirradiated by open symbols. Figure reprinted with permission from Zinkle,
S. J.; Wiffen, F. W. Radiation effects in refractory alloys. In STAIF 2004, AIP Conference Proceedings; El-Genk, M. S., Ed.;
2004; Vol. 699, pp 733–740. Copyright 2004, American Institute of Physics.

fracture toughness becomes a concern at temperatures below 800 K for Nb–1Zr.3 While irradiated tensile strength above 800 K is close to the unirradiated
values, uniform elongation values remain low until
irradiation temperatures >1000 K. Therefore, a conservative approach towards engineering design needs
to be taken with this alloy.
The mechanical properties of irradiated refractory
alloys can be influenced by the formation of He
developed through the (n,a) reactions, leading to the
grain boundary formation of bubbles and the eventual
embrittlement of the material. Some scoping investigations on the effect of He on the irradiated mechanical
properties of Nb–1Zr have been performed. Wiffen37
investigated the high-temperature mechanical properties of 50 MeV a-irradiated Nb–1Zr. In tensile tests
conducted at 1273 and 1473 K, no significant effect of
He on the strength or ductility of Nb–1Zr was observed
for samples containing 2–20 appm He. Later analysis
of the creep ductility reductions was found to be
dependent on the observed precipitate phase development through the pick-up of oxygen during implantation.38 He-implanted Nb–1Zr through 100 MeV
a-irradiations at 323 and 873 K by Sauges and Auer39
found no significant effect on ductility up to 80 appm
He. Wiffen19 observed that uniform elongations stayed

around 1% between test temperatures of 723 and
1073 K on 130 appm 10B doped Nb–1Zr irradiated in
a fast reactor between 723 and 1223 K up to 6 Â 1022
n cmÀ2. These were slightly higher than those observed
in undoped material; this is believed to be due to the
formation of He bubbles in the grains of the material

acting similar to voids in generating obstacles to
dislocation channeling. In general, no detrimental
effects on mechanical properties were reported for
accelerator-injected He between 1273 and 1473 K for
He concentrations <200 appm.37,40

4.06.3 Tantalum and Ta-Base Alloys
4.06.3.1 Introduction and History of
Ta and Ta Alloys
Tantalum and its alloys have historically been examined for high-temperature nuclear applications, particularly in the various space reactor programs. For
reasons similar to those of Nb and its alloys, various
alloying combinations of Ta were examined, particularly in the late 1950s to 1960s. Much of this effort
emphasized the development of solid solution (W and
Re additions) and dispersion-strengthened (Hf addition) alloys. While Ta-alloys pay a penalty in higher
density over, for example, Nb, and decreases the low
temperature density-compensated strength to comparable values on Nb-base alloys. The higher melting
temperature of Ta (3290 K) results in better strength
retention above 1000 K and in density-compensated
creep strength.12,41
Early work on substitutional solid solutionstrengthened Ta–10W for aerospace applications42
led to limited examination of this alloy for irradiation
environments. The improved strengthening by addition of a maximum of 10 wt% allows the retention of
suitable nonirradiated ductility and weldability.43,44



Radiation Effects in Refractory Metals and Alloys

However, the use of Ta–10W in space reactor applications where liquid alkali coolants are considered
was unacceptable because of the lack of oxide gettering elements such as Hf that form stable dispersionstrengthened structures. The T-111 (Ta–8%W–2%
Hf) alloy, with its demonstrated compatibility with
liquid alkali metals and improved strength over pure
Ta while retaining ductility and weldability, has been
a lead candidate alloy in space reactor systems
since the 1960s.45 Though a considerable effort has
been made on the Ta–10W and T-111 alloys, the
irradiation properties database is very small. Irradiated mechanical property behavior follows typical
bcc alloys in which radiation hardening effects
including limit ductility appear and are expected at
temperatures $0.3 Tm (987 K).3
4.06.3.2 Irradiation-Induced Swelling
of Ta and Ta-Base Alloys
Swelling data for Ta and its alloys are limited to a few
studies.19 Void formation in pure Ta was experimentally observed through TEM examination of material
irradiated to 2.5 Â 1022 n cmÀ2 (E > 0.1 MeV) at temperatures between 673 and 1273 K.46 An empirical
estimation of the bulk swelling taken from microstructural void size density data of that study is shown in
Figure 6. Void concentrations in the material were
highest at the peak swelling temperature and decreased
with higher irradiation temperature with an associated
increase in cavity size. Ordering of the voids at the peak
3.0
Bates and Pitner47
Wiffen46
2.5


ΔV/V (%)

2.0

Neutron fluence
2.5 ϫ 1022 n cm–2
(E > 0.1 MeV)

1.5

1.0

0.5

0.0
200

400

600

800

1000 1200 1400 1600

Irradiation temperature (K)
Figure 6 Swelling data for pure Ta measured through
microstructural void density measurements by Wiffen46 and
from immersion density measurements by Bates and Pitner.47


189

swelling condition was reported to occur along the
{110} planes in the bcc structure. A subject of considerable theoretical debate, the mechanisms of void
ordering that have appeared in bcc and fcc metals
have been examined,48–50 since the first reported
occurrence in irradiated Mo.51 Disordered void structures in the microstructure of the higher temperature
irradiated Ta appear as the size of the voids increase,
though some rafting, or grouping, was reported.46
The swelling data of Wiffen46 derived from microstructural analysis correlate well with the immersion
density data of Bates and Pitner47 (Figure 6), from
which an empirical equation for percent swelling as a
function of temperature, T (K), and fluence, F (in
units of 1022 n cmÀ2, E > 0.1 MeV), was developed,
which is as follows:
DV
¼ ðFÞ0:4 f1:69 exp½Àð0:018T À 16:347Þ2 =aŠg
V
where
a¼

14:87 þ 44:57 exp½0:09ðT À 1338:71ފ
1 þ exp½0:09ðT À 1338:71ފ

½1Š

The broader width of the swelling peak as a function
of irradiation temperature for the calculation represented by eqn [1] compared to the microstructural
data of Wiffen46 is believed to be associated with

errors in the accurate irradiation temperature of
these early measurements. Experimental evidence
of decreased swelling at higher fluences was reported
by Murgatroyd et al.52 and attributed to the transmutation of Ta to W, resulting in a shift in the lattice
constant. Similar effects have been more closely
examined in Mo and TZM alloys, and attributed to
impurity segregation at void surfaces leading to shrinkage of the voids.53
Swelling measurements in Ta–10W and T-111
alloys are limited specifically to work by Wiffen,
from which a later summary was given.19 For irradiations at 723 and 873 K to a fluence of 1.9 Â 1022 n cmÀ2
(E > 0.1 MeV), no swelling in T-111 was observed,
though a possible densification of up to 0.36% may
have occurred as evidenced in length measurements.
In companion irradiations to that of pure Ta already
discussed, involving irradiations to 4.4 Â 1022 n cmÀ2
(E > 0.1 MeV) at temperatures between 698 and
1323 K,46 samples of Ta–10W were included with
postirradiation examination involving TEM analysis.
The microstructure of the irradiated Ta–10W contained fewer voids than the companion Ta samples,
with a lower swelling assumed in the Ta–10W alloy
but with values not accurately quantifiable.19


190

Radiation Effects in Refractory Metals and Alloys

and more recently by Byun and Maloy.56 In the
first, irradiation to 0.13 dpa (where irradiation to
0.76 Â 1022 n cmÀ2, E > 0.1 MeV is $1.0 dpa in pure

Ta57) at 673 K resulted in increased yield strength,
though no significant loss in ductility occurred over
the unirradiated control. However, work softening
following the yield drop was observed.
Irradiation to higher displacement doses in pure
Ta by Wiffen19 showed the potential lower operating
temperature limitation of Ta. Following irradiation to
1.97 dpa at 663 K, yield and ultimate tensile strengths
increased to near 600 MPa with a corresponding drop
in ductility to <0.3% uniform but with total elongation near 10%. The observed plastic instability, attributed to the lack of uniform elongation following
yielding, resulted from dislocation channeling. Some
recovery of ductility is observed following irradiation to 913 K, which correlates with temperatures
approximating the maximum swelling temperature
(Figure 6) and a change in the dominating microstructural features influencing deformation behavior in
the metal. The tensile data are presented in Figure 8,
along with the irradiated properties of T-111, which
are discussed later.
The recent work of Byun and Maloy56 investigated tensile behavior as a function of fluence for
pure Ta, Ta–1W, and Ta–10W, establishing deformation mode maps for pure Ta and Ta–1W that outline
the conditions in which brittle failure and uniform
and unstable plastic deformation occur. Following
fast-reactor exposures at temperatures <373 K, a
progressive hardening and gradual loss in ductility
are observed in the tensile properties of pure Ta,
leading to a near doubling of the yield stress by
0.14 dpa over the unirradiated value (Figure 9(a)).

4.06.3.3 Mechanical Properties of
Irradiated Ta and Ta-Base Alloys
The overall mechanical property data for irradiated

Ta and Ta-base alloys are very limited, with most
studies involving irradiation at temperatures <1073 K.
In general, the behavior of Ta and its alloys is similar
to that of other bcc materials in that radiation hardening is observed with significant reductions in elongation at temperatures <0.3 Tm (Tm ¼ 3290 K, pure Ta).
As is discussed in this section, the addition of solute
strengthening elements creates an increased sensitivity
to radiation hardening of the material. In addition to
the lack of high-temperature irradiation behavior,
impact and fracture toughness data for irradiated Ta
and Ta alloys are also limited.
As with all refractory metals, the mechanical behavior of pure Ta is highly dependent on the impurity levels in the material. This may explain the
observed differences between the work of Brown
et al.54 and Chen et al.,55 of 800 MeV proton irradiations up to 11 dpa at temperatures <673 K (Figure 7).
While chemical analysis quantifying the purity of
Ta was not reported in the former, irradiation to
0.26 dpa resulted in a yield strength increase from
350 to 525 MPa over the unirradiated value with a
corresponding drop in ductility below 2%. Flow instability following yield was characteristic of samples
irradiated to 0.26 and 2.9 dpa.54 Tensile properties of
high-purity Ta irradiated to 0.6–11 dpa tested at room
temperature and 523 K showed similar increases in
tensile strength, while the uniform elongation remained near 8% following irradiation to 0.6 dpa or
higher.55
The tensile properties of neutron-irradiated Ta
were reported by Claudson and Pessl,30 Wiffen,19

Stress (MPa)

800
600


(a)

Tirr 658-673 K
Ttest = 298 K

2.9 dpa

(b)

0.26 dpa

200

0

Control

Ta

5

High-purity
Ta (99.99%)

0.6 dpa

Control

400


Tirr < 473 K
Ttest = 298 K

11 dpa

10

15

20

0
Strain (%)

10

20

30

Figure 7 Stress–strain curves of pure Ta irradiated by 800 MeV protons (a) for lower purity Ta, and (b) higher purity Ta.
(a) Reproduced from Brown, R. D.; Wechsler, M. S.; Tschalar, C. In Influence of Radiation on Material Properties: 13th
International Symposium, ASTM STP 956; Garner, F. A., Henager, C. H., Jr, Igata, N., Eds.; ASTM: Philadelphia, PA, 1987;
pp 131–140 and (b) Reproduced from Chen, J.; Ullmaier, H.; Floßdorf, T.; et al. J. Nucl. Mater. 2001, 298, 248–254.


Radiation Effects in Refractory Metals and Alloys

240


Test temperature (K)
900

700

Stress (psi)

Tantalum
Control
200
1.97 dpa at 663 K
2.5 dpa at 913 K
160

500

700

900

T-111

1600
Control
2.5 dpa at 688 K
2.5 dpa at 913 K

Ultimate
Yield


1200
Ultimate
Yield

120

800

Yield and ultimate
Ultimate

80

Stress (MPa)

500

191

Ultimate
Yield

Ultimate

40

Yield

400


Yield

Elongation (%)

0
60

0
Total

40

Total

Uniform

Uniform

20

Total

Total

Total

Uniform

Total


Uniform

0
0

200

400

0
200
600
Test temperature (ЊC)

400

600

Figure 8 Comparison of tensile properties of neutron-irradiated Ta and T-111. Uniform elongations of <0.3% for the
663 K irradiations are not shown in the figure. Reproduced from Wiffen, F. W. In Refractory Alloy Technology for Space
Nuclear Power Applications, CONF-8308130; Cooper, R. H., Jr, Hoffman, E. E., Eds.; Oak Ridge National Laboratory: Oak
Ridge, TN, 1984; pp 252–277.

An early onset of necking or plastic instability was
observed in Ta at doses above 0.0004 dpa. The lower
elongation strains in the pure Ta compared with the
work by Chen et al.55 is believed to be due to the
higher oxygen content in the material.56
The introduction of 1 wt% W resulted in a near30% increase in unirradiated strength over pure Ta

(Figure 9(b)). The Ta–1W alloy showed greater sensitivity to radiation hardening than the pure metal.
The tensile properties as a function of dose were
similar to those of the pure Ta. However, above
0.004 dpa, plastic instability becomes more predominant in the Ta–1W alloy and occurs immediately
following yielding. For Ta–1W irradiated from 0.7 to
7.5 dpa in a mixed proton and neutron irradiation from
the same study, hardening was saturated with little
change in ductility (insert shown in Figure 9(b)).
Macroscopic deformation mode maps produced
for Ta and Ta–1W by Byun and Maloy56 are a graphical way of predicting the performance of a material in
an irradiation environment. The deformation mode
map for pure Ta is shown in Figure 10(a), while that

of Ta–1W is shown in Figure 10(b). The yield and
plastic instability stress were directly obtained from
tensile data, while the fracture stress was calculated
through a linear strain hardening model for necking
deformation, assuming that during instable deformation, the strain hardening rate remains unchanged and
is approximately the plastic instability stress. The fracture and plastic instability stresses are independent of
dose, with a ratio between the stresses of $2 for the
materials studied. The fracture strength decreases with
dose if the material becomes embrittled, for example,
through interstitial segregation or secondary phase
precipitation at grain boundaries, though this was not
observed in their work. The yield strength is highly
dose dependent, though the yield stress was significantly lower than the fracture strength in Ta–1W,
suggesting that the material may show limited ductility to even higher displacement doses. The effect of
increasing test temperature for each material further
increases the boundaries for uniform deformation
behavior. This increase was found to be greater in

pure Ta.


Radiation Effects in Refractory Metals and Alloys

Engineering stress (MPa)

1000

Ta
Neutron irradiated
Tirr 333–373 K

800

600

Engineering stress (MPa)

192

0.14 dpa

0.04

400

0.004
0.0004


200

Unirr.

0.00004

1000

Ta–1W
Neutron irradiated
Tirr 333–373 K

800

Proton and neutron irradiated
Tirr 323–433 K

800

7.5 dpa
4.6

600

0.14 dpa

400

600


0.04

0

400

200

2.0

0.004

Unirr.
0.7

200

0

10

20

30

0.00004

40

0.0004

Unirr.

0
(a)

0

10

20
30
Elongation (%)
1600

40

10

0

20
30
Elongation (%)

(b)

25.2 dpa

40


50

Ta–10W
Proton and neutron irradiated
Tirr 323–433 K

1400
Engineering stress (MPa)

0

1200
7.5
1000
800

4.4

Unirr.

0.7
2.0

600
400
200
0

(c)


0

5

15
10
Elongation (%)

20

Figure 9 Room temperature tensile curves for irradiated (a) pure Ta, (b) Ta–1W, and (c) Ta–10W. Reproduced from
Byun, T. S.; Maloy, S. A. J. Nucl. Mater. 2008, 377, 72–79.

The room temperature unirradiated tensile strength
of Ta–10W is nearly double the value of the Ta–1W
and triple that of pure Ta in the material investigated
by Byun and Maloy,56 and also shows an increased
sensitivity in radiation hardening over the pure metal
(Figure 9(c)). This sensitivity is also clearly apparent
at higher irradiation temperatures near 673 K, as
shown in the comparison of tensile curves that were
compiled by Ullmaier and Carsughi58 of earlier work
(Figure 11). Near room temperature irradiation of
Ta–10W to the mixed proton and spallation neutron
exposure by Byun and Maloy56 to doses between 2 and
25.2 dpa showed prompt necking following yielding.
Total elongation values of <3% were observed for
doses between 2 and 7.5 dpa, with near-zero ductility
observed at 25.2 dpa. Fast neutron irradiation studies of
Ta–10W by Gorynin et al.59 observed brittle failure

after 0.13 dpa in materials irradiated and tested near
600 K. Less than 5% total elongation was measured
following 1.97 dpa irradiation at 700 K, despite a near

doubling of the yield stress over the unirradiated material. Limited ductility was also observed following
2.63 dpa exposure in materials irradiated and tested
at 1073 K, with a yield strength increase from 240 to
315 MPa over the unirradiated control. While lowtemperature embrittlement following exposure to
0.13 dpa was reported in the neutron-irradiated materials59 and limited ductility following mixed proton
and neutron exposure,56 the interstitial concentrations on the behavior of these materials may be
more influential than the irradiation spectrum.
Similar to Ta and Ta–10W, very limited data exist
on the irradiated properties of T-111. The most
referenced base-line study is that by Wiffen,19
shown as in Figure 8. Large increases in yield and
ultimate tensile strengths are observed following irradiations to 1.9 Â 1022 n cmÀ2 (E > 0.1 MeV), 2.5 dpa,
at 688 and 913 K. The increase in radiation hardening
is substantially greater than that observed in pure
Ta irradiated under similar conditions. Yield and


Radiation Effects in Refractory Metals and Alloys

2000

1000

Fracture region

Fracture region


800

1500
True stress (MPa)

True stress (MPa)

193

600
Plastic instability region
400
Elastic region

200

1000

500

Uniform plasticity (523 K)
Uniform plasticity (Trm)

0
0.0

0.0001

0.001

0.01
Dose (dpa)

(a)

Plastic instability region

Elastic region
Uniform plasticity (523 K)
Uniform plasticity (Trm)

0.1

0
0.0

1
(b)

0.0001 0.001 0.01
0.1
Dose (dpa)

1

10

Figure 10 Deformation mode map of (a) pure Ta and (b) Ta–1W for room temperature irradiations, illustrating fracture,
plastic instability, uniform plasticity, and elastic regions as a function of stress and displacement dose. The increases in the
uniform plasticity region for temperatures of 523 K are superimposed. Reproduced from Byun, T. S.; Maloy, S. A. J. Nucl.

Mater. 2008, 377, 72–79.

400

Tirr~673 K

360
0.39 dpa

320
Stress (MPa)

280

Ta-10W
Control

240
200

Ta
Control

160

0.13 dpa

120
80
40

0

0

2

4
6
Strain (%)

8

Figure 11 Comparison of the radiation hardening of Ta
and Ta–10W irradiated at $673 K to displacement doses
of <0.39 dpa. Adapted from Ullmaier, H.; Casughi,
F. Nucl. Instr. Methods Phys. Res. B 1995, 101, 406–421.

ultimate tensile strengths of around 1250 MPa are
reported for irradiation at 688 K, with uniform and
total elongation of <0.3% and 4.5%, respectively.
Irradiation at 913 K improves uniform and total elongation values only slightly to $2.5% and 8%. These
values represent more than a 50% reduction in ductility over the unirradiated values. No known irradiated
property data for T-111 exist for temperatures above

913 K. As tensile strengths of both Ta–10W and T-111
exceed 1000 MPa at temperatures below 1000 K19,56,59
and are well above the stresses that produce brittle
behavior in vanadium alloys for which more data are
available, it is likely that these Ta alloys are embrittled
under these conditions.3 Further expansion of the

irradiated materials database including fracture toughness data for Ta and Ta alloys irradiated near and
above 1000 K is much needed to ascertain the upper
temperature limitations. However, based on this preliminary data, temperatures below1000 K may need to
be avoided for Ta and Ta-base alloys.
Low-fluence irradiations to 1.2 Â 1015 n cmÀ2 at
room temperature and 623 K have been performed
to evaluate the performance of T-111 and Ta–10W
for use in radioisotope power applications.60 These
low-dose irradiations produced little change in the
tensile properties of the two alloys. Some variations
in the total elongation were observed in T-111, which
may be related to the distribution and make-up of the
Hf-rich compounds in the material as well as the
effects of radiation. Thermal stability of T-111 can
be an issue, as a brittle behavior following 1100 h
aging at 1398 K has been reported,41 due to precipitation of Hf-rich compounds along grain boundaries.
It is not known how the combination of long-term
thermal aging under irradiation affects the structure–
property relationships or how the detrimental precipitation of the interstitial elements with Hf can be
controlled.


194

Radiation Effects in Refractory Metals and Alloys

4.06.4 Molybdenum and Mo-Base
Alloys
4.06.4.1 Introduction and History of Mo and
Mo Alloys

Molybdenum and its alloys are the perennial candidates for refractory metal alloy use in irradiation
environments, due in part to their high melting
temperature (2896 K), good thermal properties,
high-temperature strength, and lower induced
radioactivity (as compared to tantalum). The density
of molybdenum (10.28 g cmÀ3) is also significantly
lower than that of Ta and W, though greater than Nb.
But like other refractory metal alloys, Mo can present difficulties in fabrication, low-temperature ductility, and low-temperature embrittlement from
radiation damage. The TZM (Mo–0.5%Ti–0.1%
Zr) and Mo–Re alloys were examined as part of the
SP-100 and JIMO/Prometheus space reactor programs, respectively, and offer additional benefits of
improved high-temperature strength over the pure
metal.5,19 Molybdenum and its alloys have also been
examined for plasma facing and diverter components
in fusion reactor designs due to the relatively low
sputter yield, high thermal conductivity, and thermal
compatibility with other structural materials.5–27,29–63
In addition, because of these benefits, Mo has also been
examined for use as a grazing incident metal mirror in
fusion diagnostic port designs.64,65
As in all other refractory metals, the mechanical
properties are influenced by impurity concentrations, particularly through grain boundary weakening. However, improvements in Mo ductility are
achievable through grain refinement, impurity control, and the addition of Re or reactive elements such
as Ti and Zr. An upper limit to the acceptable level of
C was also found to improve grain boundary strength.
Low-carbon arc-cast molybdenum (LCAC-Mo) is
one such example, in which oxygen impurities are
reduced to tens of ppm, nitrogen to <10 ppm, and
carbon to <100 ppm.66 Higher levels of C will result
in reduced fracture toughness, unless additional reactive alloy additions are present in the alloy. The TZM

alloy also incorporates a small level of carbon to
produce Ti- and Zr-carbide strengthening.
Improvements in ductility and toughness through
the ‘rhenium effect’ have been observed in Mo for
some time,67–69 and generally occurs when Group
VIa metals are alloyed with elements from Group
VIIa and VIIIa metals.70,71 Explanations for this phenomenon range from enhanced mechanical twinning,
reduced resistance to dislocation glide, reduction of

oxygen at grain boundaries, and increased interstitial
oxygen solubility.67,71–75 Critical evaluation76,77 of
the initial work that had suggested a maximum tensile ductility near 11–13 wt% Re78,79 was found to be
inconclusive because of inadequate control of O and
C impurity levels in the earlier studies. Higher concentration alloys with 40–50 wt% Re have also been
examined for use in the radiation environments.
Alloys with Re concentrations up to 41–42% are
single-phase solid-solution a-Mo, while those at
higher levels incorporate the s-Re2Mo phase. Commercially available alloys include Mo–41Re and
Mo–47.5Re (sometimes referred to as Mo–50Re).
Recently, introduction of oxide dispersion strengthened (ODS)-Mo through the incorporation of lanthanum oxide particles has been examined.80–82 These
alloys show great resistance to recrystallization and
high-temperature deformation while maintaining low
ductile-to-brittle transition temperatures (DBTT)
partly because of their refined grain structure.83–85
The radiation effects database for Mo and its
alloys is limited to scattered scoping examinations,
which show little overlap in the experimental variables such as material purity, alloying level, material
thermomechanical history, irradiation conditions, and
postirradiation test conditions. Where available, information on the physical and mechanical property
changes to LCAC-Mo, TZM, Mo–Re alloys, and

ODS-Mo will be reviewed.
4.06.4.2 Irradiation-Induced Swelling and
Physical Property Changes in Mo and
Mo-Base Alloys
Two earlier reviews of the irradiation-induced properties of Mo and TZM have been presented as part of
the UWMAK-III fusion reactor study86 and for the
SP-100 space nuclear power program.19 Much of the
known swelling data on irradiated Mo is contained in
these reviews, with the majority of data for irradiations <10 dpa and temperatures below 1073 K. The
swelling data available are considerably scattered,
with little coherence to examinations on the swelling
as a function of temperature or dose.
Swelling in Mo is expected to begin around 573–
673 K and continue to temperatures near 1573 K.19
Maximum swelling in pure Mo remains below 4%
for fluences up to 1 Â 1023 n cmÀ2 (E > 0.1 MeV),
$50 dpa, with peak swelling at irradiation temperatures near 900 K. Attempts at consolidating the
reported swelling data as a function of irradiation
temperature through normalizing the fluences proved


Radiation Effects in Refractory Metals and Alloys

195

at 1176 K.91 Swelling is expected to reach a maximum
of 3–4% on the development of the void lattice structure, based on an attainment of an equilibrium ratio of
void diameter to void superlattice parameter.92 At
temperatures >1423 K, void lattice formation is no
longer observed, leading to the high values of swelling

observed in the material ion irradiated to high doses.88
The onset of void growth in neutron-irradiated
material appears to be accelerated in cold-worked
materials compared to annealed materials, reaching
a maximum in swelling at doses near 40 dpa for temperatures below 873 K and 20 dpa at higher temperatures.89 At higher doses, swelling decreases through
void shrinkage, with swelling values approaching
those of annealed materials. Void shrinkage has also
been reported by Bentley et al.93 and Evans53 to occur
because of changes in the void sink bias89 presumably due to the segregation of transmuted species
at the void surfaces, making them more attractive
for interstitials.
Irradiation-induced swelling in TZM has been
reported53,94–97 and generally shows similar temperature dependence as the pure metal. The fluence and
temperature dependence of swelling of TZM was
examined by Powell et al.95 and Gelles et al.,94 with

to be inaccurate in determining the upper bound
limit for maximum swelling.19 The swelling data collected from numerous sources,53,87–89 including those
contained in the review work of Brimhall et al.90
for irradiated Mo as a function of dose and irradiation temperature, are provided in Figure 12. Void
swelling was found to be <1% in Mo irradiated to
8 Â 1022 n cmÀ2, E > 0.1 MeV at temperatures between
673 and 1173 K by Evans.53 Void swelling studied by
Stubbins et al.88 in 3.1 MeV 51V+ ion-irradiated Mo
between 1173 and 1393 K up to 50 dpa remained
below 4%, while irradiations between 1523 and
1780 K were near 10%.
Void ordering has been observed in both neutronirradiated28,89,91 and ion-irradiated Mo88 at temperatures between 700 and 1373 K. Garner and Stubbins89
examined the irradiation and material conditions
that contribute to void ordering. Irradiation temperatures near 700 K delineate the lower boundary temperature for void lattice formation at irradiations

above 20 dpa. At lower doses, void lattice formation
was not observed. The void superlattice constant, measured as the distance between void centers along the
<100> direction in the material, is found to increase
with temperature from $2.4 nm at 700 K to 4.5 nm

4
Garner and Stubbins89 (neutron)
Stubbins et al.88 (ion)
Evans53 (neutron)
Lee et al.87 (neutron)
Brimhall et al.90 (neutron)
Brimhall et al.90 (ion)

DV/ V (%)

3

2

1
100
0

10
400
800
ion
tem 1200
per
atu

re

diat

0.1
(K)

1600

dp
a

1

Irra

0.01

Figure 12 Irradiation-induced swelling (DV/V ) as a function of irradiation temperature and displacement damage (dpa)
for pure Mo. The irradiation source is marked in the key. Reproduced from Lee, F.; Matolich, J.; Moteff, J. J. Nucl. Mater.
1976, 62, 115–117; Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41; Stubbis, J. F.; Moteff, J.; Taylor, A. J. Nucl. Mater. 1981, 101,
64–77; Garner, F. A.; Stubbins, J. F. J. Nucl. Mater. 1994, 212–215, 1298–1302; Brimhall, J. L.; Simonen, E. P.; Kissinger,
H. E. J. Nucl. Mater. 1973, 48, 339–350.


196

Radiation Effects in Refractory Metals and Alloys

results from the latter shown in Figure 13. Peak

swelling in TZM following irradiation to 1.78 Â 1023
n cmÀ2 and 873 K remained below 4%, though the
data are limited to irradiation temperatures below
923 K. Only limited data are available on direct comparisons between TZM and pure Mo, with Bentley
and Wiffen96 reporting 1% swelling in Mo–0.5%Ti
and TZM alloys and 0.6% swelling in pure Mo under
the same irradiation conditions. Similarly, 4%
swelling was observed in TZM and 3% in pure Mo
following irradiation to 5.4 Â 1022 n cmÀ2 at 923 K.97
In examining Mo and TZM of different preirradiated material conditions, Evans53 observed equal
or greater swelling in TZM compared to Mo following irradiation to 3.5 Â 1022 n cmÀ2 (E > 0.1 MeV) at
823 and 873 K. However, in the materials irradiated
at 723 K for the same fluence, the TZM alloy showed
lower swelling, except in the carburized condition.
The Ti and Zr atoms not tied up as carbides are
assumed to have played a role in reducing void size
in the material at the lower temperature.
There is little information on the swelling behavior of Mo–Re alloys. Measured swelling of 0.44% in
Mo–50Re irradiated to 5.3 Â 1022 n cmÀ2 (E > 0.1
MeV) at temperatures which rose during irradiation
from 1128 to 1329 K was reported.26 For irradiated
Mo–Re alloys, radiation-induced segregation (RIS)
and transmutation can lead to precipitation of

equilibrium or nonequilibrium phases, which can be
detrimental to mechanical properties. This is examined
in the next section.
Electrical resistivity changes to 5.4 dpa irradiated
Mo at 733 K were examined by Zakharova et al.98
using single crystal samples. Increases in resistivity

of 10–14% and 92–110% were measured at postirradiation test temperatures of 298 and 77 K, respectively. The largest resistivity changes were measured
in the [100] direction. A residual 10% increase in
resistivity was measured following annealing above
0.6 Tm associated with the accumulation of transmuted radionuclides.
The changes in electrical resistivity of LCAC-Mo
over a 353–1373 K irradiation temperature range up
to 3.3 dpa were examined by Li et al.99 and Cockeram
et al.,100 with the latter examining the recovery of
resistivity following isochronal anneals. The room
temperature resistivity for 353 K irradiated LCACMo rapidly increases between 0.01 and 0.1 dpa
saturating near 0.2 dpa for an $42% increase over
the unirradiated value.99 Increases in room temperature resistivity of 10–12% were reported following
0.5–1.2 dpa irradiation at 543 K, and 3.3–5.3% after
1.4–2.4 dpa at 878 K. At irradiation temperatures
!1208 K, little (<3%) to no net increase in resistivity
was observed for irradiations up to 3.3 dpa. This is
reflected in the higher mobility of vacancies and

4.0
3.5

DV/V (%)

3.0

Gelles et al.94 (11-20 dpa)
Gelles et al.94 (36-53 dpa)
Gelles et al.94 (47-65 dpa)
Evans53


2.5
2.0
1.5
1.0
0.5
80

0.0
600

60
700

40

800
Ti

rr

(K)

900

20
1000

a

dp


1100
Figure 13 The swelling dependence on temperature and fluence for neutron-irradiated TZM. Adapted from Gelles, D. S.;
Peterson, D. T.; Bates, J. F. J. Nucl. Mater. 1981, 103–104, 1141–1146; Evans, J. H. J. Nucl. Mater. 1980, 88, 31–41.


Radiation Effects in Refractory Metals and Alloys

interstitials formed during irradiation to diffuse to
sinks where annihilation occurs, reducing the electrical scattering effects that these defects have at lower
irradiation temperatures. The small increases measured at the higher irradiation temperatures were primarily due to transmutation products. As is shown in
the next section, the changes in electrical resistivity
with increasing irradiation temperatures also correlate
with changes in measured hardness, though at a greater
level of sensitivity. This is controlled by microstructural changes, as the small dislocation loops and voids
of high distribution density appearing at the lower
irradiation temperatures coarsen into larger and
fewer defects that have less interaction with deformation dislocations.
4.06.4.3 Mechanical Properties of
Irradiated Mo and Mo Alloys
The mechanical property performance of pure Mo is
strongly controlled by the grain size, oxygen, nitrogen,
and carbon concentrations as well as alloy additions.
This is true for unirradiated as well as irradiated
properties.71,76,83,84 The sensitivity to embrittlement
at irradiation temperatures 873 K can be mitigated
through a reduction in oxygen and nitrogen while
keeping the carbon-to-oxygen ratio high to reduce
the segregation of oxygen and nitrogen to the grain
boundaries. A reduction in the grain size can further

increase the number of sinks and reduce the mean
distance that irradiation-induced defects must travel
at temperatures at which mobility is limited.
Irradiated mechanical properties of wrought
LCAC-Mo in both the recrystallized and stressrelieved conditions have been examined over several
decades.81,82,84–86,99–106 In general, LCAC-Mo undergoes significant increases in tensile strength
through radiation hardening 873 K, which produces
reductions in ductility and high DBTT values.
A summary of tensile properties as a function of irradiation temperature and dose is shown in Figure 14.
Irradiated stress-relieved LCAC-Mo shows less radiation embrittlement than as-crystallized materials at
temperatures <1208 K.100 However, at higher irradiation temperatures, recrystallization of the stressrelieved material occurs, leading to large changes in
the microstructure and less desirable properties.
Increases in hardness of 56% and 112% for stressrelieved and recrystallized LCAC-Mo, respectively,
are reported following irradiation to 1.2 dpa at
543 K.100 The increase in hardness decreases slightly
following 878 K irradiations to 2.4 dpa, but returns to

197

values near those for the unirradiated material for
irradiation at temperatures !1208 K. Materials irradiated in the stress-relieved condition at temperatures !1173 K can result in softening compared to
unirradiated materials because of recrystallization.107
A comparison of the changes in irradiated material
hardness as a function of dose and temperature is
plotted in Figure 15(a) for LCAC-Mo, TZM, and
ODS-Mo. The two latter alloys are discussed in
detail later. In the three materials investigated by
Cockeram et al.,107 the largest increase in hardness
was measured for irradiations at 873 K, counter to
what is observed in tensile properties for irradiation

between 573 and 873 K.81,85 However, tensile failure
in the lower temperature irradiated samples generally
occurs before the samples yield because of the elevation of the flow stress above the fracture stress for these
test conditions. Recovery of the hardness, measured
through 1 h anneals at increasing temperatures, is plotted in Figure 15(b) for the $573 K irradiated material.
The start of recovery is near 800 Kwith full recovery of
the LCAC-Mo hardness by 1253 K. For LCAC-Mo
irradiated at a higher temperature of 873 K, recovery
of hardness begins near 1163 K and is completed
near 1463 K.
Substantial increases in the DBTT to values
>773 K are observed for irradiated LCAC-Mo over
a range of fluences for irradiation temperatures
873 K.26,82,85,103,105,108 A summary of DBTT values
for LCAC-Mo is presented in Figure 16, along with
data from high-purity grain-refined LCAC-Mo,
TZM, and ODS-Mo, which is discussed next. In
general, recovery in the DBTT of LCAC-Mo is not
observed until irradiation temperatures above 873–
973 K, depending on material conditions. A reduced
sensitivity to low-temperature irradiation embrittlement of LCAC-Mo is observed in materials with
reduced levels of impurities. A high-purity form of
LCAC-Mo (HP-LCAC-Mo) was developed through
the use of 1873 K hydrogen atmosphere annealing of
LCAC-Mo plates prior to further arc casting, extrusion, and rolling into sheet stock.82 Levels of oxygen
and nitrogen were <4 wppm each, with carbon at
20 wppm. Average grain diameters of 1.3 and 452 mm
lengths were produced, and represent a considerable
change in aspect ratio compared to LCAC-Mo values
of 4–5 mm diameter and 78–172 mm lengths.100 The

DBTT of 573 K irradiated HP-LCAC-Mo showed
no increase over the unirradiated value (123 K) for
irradiations up to 0.11 dpa, and increased to 723 K by
1.29 dpa. For 873 K irradiations, the DBTT remained
below 223 K up to 1.29 dpa.


198

Radiation Effects in Refractory Metals and Alloys

2000
HP-LCAC (0.11-1.29 dpa, 573 K)

LCAC (unirr.)

LCAC (0.6 dpa, 573 K)

Yield stress (MPa)

1500

LCAC (13.1 dpa, 833-1057 K)
LCAC (12.3 dpa, 567 K)

HP-LCAC (unirr.)

EB (0.23 dpa, 873 K)

1000


LCAC (0.9-1.4 dpa, 878 K)

500
LCAC (3.3 dpa, 1373 K)

25

LCAC (3.3 dpa, 1373 K)

Total elongation (%)

20

LCAC (unirr.)

15

EB (0.23 dpa, 873 K)

10

LCAC (0.9 -1.4 dpa, 878 K)

HP-LCAC (unirr.)

5

HP-LCAC (0.11-1.29 dpa, 573 K)


LCAC (13.1 dpa, 833-1057 K)
LCAC (12.3 dpa, 567 K)

0
0

200

400

600

800

1000

1200

Test temperature (K)
Figure 14 Yield stress and total elongation as a function of test temperature for LCAC-Mo, HP-LCAC-Mo and electron
beam (EB) melted Mo. Irradiation dose and temperature (dpa, K) is marked. Unirradiated samples (unirr.) are also
presented. Reproduced from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008,
382, 1–23; Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Abe, K.; Takeuchi, T.; Kikuchi, M.;
Morozumi, S. J. Nucl. Mater. 1981, 99, 25–37.

The majority of the irradiated mechanical property database for TZM is limited to displacement
damage <5 dpa and irradiation temperatures
<1000 K, with much of the testing conducted at temperatures below that used in the irradiation. The database is also sparse, with little connectivity between

different experimental examinations. Furthermore,

significant differences are observed in the nonirradiated tensile values of TZM because of variations in
material processing leading to differences in grain size,
impurity level, distribution of the strengthening phase,
and differences in testing procedure. In general,


Radiation Effects in Refractory Metals and Alloys

400

LCAC-Mo (Tirr = 573 K)

LCAC-Mo (Tirr = 873 K)

LCAC-Mo (Tirr = 1173 K)

ODS-Mo (Tirr = 573 K)

ODS-Mo (Tirr = 873 K)

ODS-Mo (Tirr = 1173 K)

TZM (Tirr = 573 K)

TZM (Tirr = 873 K)

TZM (Tirr = 1173 K)

199


Change in hardness (MPa)

300
873 K

200
573 K

100

0
1173 K
-100
0.00

2.00

4.00

(a)

1700

Recovery of hardness

6.00
8.00
Displacement dose (dpa)

10.00


12.00

14.00

>1573 K

75-100%
1500

50-75%
25-50%

1 h anneal temperature (K)

0-25%
1300

No change

1100

900

700

500
Dose 0.6 dpa 12.3 dpa 1.4 dpa 13.1 dpa 12.3 dpa 3.9 dpa 13.1 dpa 12.3 dpa 3.9 dpa 13.1 dpa
Tirr 543 K 567 K 877 K 833 K 573 K 877 K 833 K 573 K 882 K 833 K
LCAC LCAC LCAC LCAC ODS

ODS
ODS
TZM
TZM
TZM
(b)

Figure 15 (a) Change in hardness as a function of neutron dose for LCAC-Mo, TZM, and ODS-Mo irradiated between
573 and 1173 K up to 13.1 dpa. (b) Recovery of hardness as a function of isochronal annealing temperature for material
irradiated $573 K. Adapted from Cockeram, B. V.; Smith, R. W.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2009, 393, 12–21.


200

Radiation Effects in Refractory Metals and Alloys

800

LCAC-Mo
Wiffen26 (13.3-23.4 dpa)
Chakin and Kazakov108 (2.1-3 dpa)
Hasegawa et al.106 (10.5-50 dpa)
Abe et al.104,105 (0.6-1.4 dpa)
Webster et al.148 (7.4-10 dpa)
Smith and Michel149 (5.3 dpa)
Scibetta et al.109 (0.15-0.19 dpa)
Cockeram et al.85 (12.3-13 dpa)
HP-LCAC-Mo
Cockeram et al.82 (1.29 dpa)
TZM

Cockeram et al.85 (12.3-13 dpa)
Wiffen26 (16 dpa)
ODS-Mo
Cockeram et al.85 (12.3-13.1 dpa)

700
600

DBTT (°C)

500
400
300
200
100
0
-100
0

100 200 300 400 500 600 700 800 900 1000 1100 1200 1300
Irradiation temperature (°C)

Cockeram et al.82 (1.29 dpa)

Figure 16 Summary of DBTT values as a function of neutron irradiation temperature for LCAC-Mo and TZM. Adapted
from Cockeram, B. V.; Smith, R. W.; Leonard, K. J.; Byun, T. S.; Snead, L. L. J. Nucl. Mater. 2008, 382, 1–23;
Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165.

displacement damage to the TZM alloy produces a
large increase in the yield strength of the material with

a corresponding drop in total elongation, with the
level of change increasing with dose and/or lower
irradiation temperatures. Some recovery of properties
begins to be observed at test conditions above 973 K for
materials irradiated at lower temperatures. A compilation of earlier and more recent tensile data for irradiated TZM is provided in Figure 17. Irradiation
1.2 dpa and temperatures >800 K showed little effective strengthening above the unirradiated values, for
tensile tests below the irradiation temperature, 85,109,110
though higher displacement doses resulted in a significant increase.81 Total elongation in material irradiated
<1000 K was limited, with uniform elongation values
<1%.111 Plastic instability following yielding was also
observed in irradiated TZM as well as LCAC-Mo. The
plastic instability stress, defined as the maximum true
stress in which plastic necking occurs, is strongly
dependent on test temperature but nearly independent
of fluence and irradiation temperature.81
Increases in the DBTT for irradiated TZM are
significant and do not diminish until irradiation temperatures slightly higher than that of LCAC-Mo,81
but comparisons may be difficult because of differences in materials and test methods. The DBTT value
for unirradiated TZM is $200 K and increases with
increasing irradiation damage. An increase of 230 K in
the DBTT over unirradiated values was reported for
TZM irradiated to 2–4.8 dpa at 644–661 K.111 DBTTs

of 750 K for Mo–0.5Ti irradiated to $16 dpa at
727 K26 and 1073 K for TZM irradiated to 12.3 dpa
at 573 K81,85 were also reported. A summary of DBTT
as a function of irradiation temperature for LCAC-Mo
and TZM is shown in Figure 16 for irradiation doses
<50 dpa.85 Excessive embrittlement is observed for
irradiations at <773 K. A reduction in the DBTT for

LCAC-Mo appears near 873 K, while the DBTT of
TZM remains high. Recovery of DBTT to near unirradiated values occurs for irradiations at >1073 K.81,85
The stage V recovery temperature for vacancy diffusion in Mo is $873 K; however, the kinetics for microstructural changes to occur are still relatively slow at
this temperature. Therefore, embrittlement issues can
be present until $1073 K.
Very little fracture toughness data exist for
irradiated TZM. For precracked compact tension
specimens irradiated to 0.29–0.35 dpa at 313–695 K,
a 4 MPa√m decrease in fracture toughness (15–
20 MPa√m at Trm, unirradiated84,109) was observed
up to the irradiation temperature.109 Kitsunai et al.112
examined the impact toughness of irradiated TZM
and alloys, incorporating 0.1–1 wt% TiC additions
to Mo. The Mo–TiC alloys showed dramatically
improved toughness levels over TZM with increasing
TiC concentration in samples irradiated to 0.08 dpa
between temperatures of 573 and 773 K. Shown
in Figure 18 is the shift in DBTT to lower temperatures for the irradiated TiC-containing Mo over
the TZM alloy. For the 1 wt% containing sample,


Radiation Effects in Refractory Metals and Alloys

1974 Kasakov et al.: unirradiated

201

2005 Cockeram et al./Byun et al. (unirradiated)

1974 Kasakov et al. (1.6 dpa, T irr = 823 K)


1974 Kasakov et al. (1.2 dpa, T irr = 1223 K)

2005 Cockeram et al. (3.9 dpa, T irr = 833 K)

2005 Cockeram et al. (12.3 dpa, T irr = 567 K)

2008 Byun et al. (3.9 to 13.1 dpa, T irr = 567 - 833 K)

2008 Byun et al. (3.9 to 13.1 dpa, T irr = 1057 - 1209 K)

1400

Yield strength (MPa)

1200

1000

800

600
Unirradiated range
400

200
200

400


600

(a)

800

1000

1200

1400

Test temperature (K)
2005 Cockeram et al./Byun et al. (unirradiated)
1974 Kasakov et al. (1.2 dpa, T irr = 1223 K)

1974 Kasakov et al.: unirradiated
1974 Kasakov et al. (1.6 dpa, T irr = 823 K)
2005 Cockeram et al. (3.9 dpa, T irr = 833 K)

2005 Cockeram et al. (12.3 dpa, T irr = 567 K)
1976 Steichen (4.79 dpa, T irr = 661 K)

1976 Steichen (2.12 dpa, T irr = 644 K)

20
18

Total elongation (%)


16
14
12
Unirradiated
range

10
8
6
4
2
0
200

(b)

300

400

500

600

700

800

900


1000

1100

1200

Test temperature (K)

Figure 17 Neutron-irradiated tensile data for TZM (a) yield stress versus test temperature and (b) total elongation versus
test temperature. Reproduced from Byun, T. S.; Li, M.; Cockeram, B. V.; Snead, L. L. J. Nucl. Mater. 2008, 376, 240–246;
Cockeram, B. V.; Smith, R. W.; Snead, L. L. J. Nucl. Mater. 2005, 346, 165; Kasakov, V. A.; Kolesnikov, A. N.;
Krassnoselov, V. A.; et al. Effect of neutron irradiation on properties of potential structural material for thermonuclear
reactors, USSR-US Exchange on CTR Materials, Nov 1974; Steichen, J. M. J. Nucl. Mater. 1976, 60, 13–19.

the DBTT remained unchanged with irradiation,
despite a $50% increase in Vickers microhardness.
More surprisingly, the Mo–1%TiC sample increased
in toughness following 0.8 dpa irradiation attributed
to grain boundary strengthening by (Ti, Mo)C
radiation-enhanced precipitates.
Developed to improve the low-temperature ductility and weld characteristics of unalloyed Mo, the
Mo–Re alloys have gained considerable attention
over the past decade for use in nuclear applications.
Single-phase solid solution a-Mo phase field extends

up to $42 wt% Re, above which the s-MoRe2
phase precipitates. At higher Re concentrations, the
w-MoRe3 phase is present. However, the exact phase
boundaries are not well delineated at temperatures
below 1773 K113,114 mainly because of the slow kinetics in phase development.115

The Mo–Re alloys show a hardening response to
irradiation stronger than that of the pure metal and
TZM following irradiation.116–118 The hardening
response of Mo–Re alloys ranging in composition
from 2 to 13 as well as 41 wt% Re following


202

Radiation Effects in Refractory Metals and Alloys

0.9
TZM
Total absorbed energy (J mm-3)

0.8

Mo-0.1%TiC
Mo-0.5%TiC

0.7

Mo-1%TiC
0.6
0.5
0.4
0.3
0.2
0.1
0

200

250

300
350
400
Test temperature (K)

450

500

Figure 18 Total absorbed energy versus test temperature
for TZM and 0.1–1.0 wt% TiC additions to Mo irradiated to
0.08 dpa at 573–773 K. Reproduced from Kitsunai, Y.;
Kurishita, H.; Narui, M.; Kayano, H.; Hiraoka, Y. J. Nucl.
Mater. 1996, 239, 253–260.

irradiation up to 20 dpa at temperatures between 681
and 1072 K was examined by Nemoto et al.116 A linear
increase in hardness with Re concentration was observed for the unirradiated controls as well as samples
irradiated at temperatures 874 K. For samples irradiated at 1072 K, little variation was observed with
increasing Re content, though hardness values
remained nearly double that of the unirradiated material. The dependence of hardness on irradiation temperature and fluence for Mo–5Re and Mo–41Re in
comparison with LCAC-Mo is presented in Figure 19.
The high degree of radiation hardening at temperatures <1000 K exhibited by the Mo–Re alloys
is further reflected in low ductility and reported
embrittlement. Tensile elongation values of <0.3%
were reported for Mo–5Re irradiated to 0.16 dpa and

tested at the irradiation temperature of 320 K,118
though higher total elongations of 8% were reported
for Mo–5Re fast reactor irradiated to 0.29 dpa and
tested near the irradiation temperature of $723 K.109
Unlike LCAC-Mo and TZM, the little deformation
occurring in Mo–Re is mostly uniform at temperatures below 1000 K,118,119 while a small degree of
work softening has been observed at higher temperatures.120 No evidence of dislocation channeling was
found during microstructural examination of tensile
tested Mo–5Re irradiated to 0.16 dpa at 373 K.118
For comparable irradiation fluences, dislocation
loop concentrations are approximately two times
higher than TZM and four to six times higher than
Mo, while dislocation loop diameters are smaller in

the Mo–5Re alloy.118 Void development begins to
appear in Mo–Re alloys above 623 K118 and shows a
slight increase in size with Re concentration (with
corresponding decrease in number density) up to
10 wt% with no further increase for the 41% Re
alloy.116 Irradiation-induced void swelling in Mo with
Re concentrations 13 wt% is 0.5–1.5% for $21 dpa
irradiated material at 681–1072 K, while swelling for
41 wt% Re was near 0.1%.116
Radiation-induced precipitation has been reported
by Nemoto et al.116 in Mo–(2–41)Re alloys irradiated
to 21 dpa between 681 and 1072 K, and by Edwards
et al.121 in Mo–41Re irradiated 28–96 dpa at 743–
1003 K. An initial formation of hcp-structured precipitates with a thin plate-like morphology consisting of
solid solution Re and Os was observed, appearing with
the {110}Mo//{0001}Re, <111>Mo//<2110>Re orientation relationship.121 The precipitation of these plates

on dislocation loops resulted in the high density of plates
observed, which dominates the microstructure. On further irradiation to higher doses or higher temperatures,
these plates develop and coarsen into the w-phase. This
nonequilibrium phase development in Re-lean alloys
was originally observed by Erck and Rehn122 in
Mo–(27–30)Re irradiated by 1.8 MeV He ions at 1023–
1348 K. The s-MoRe2 phase was also reported appearing in all the Mo–Re samples examined by Nemoto and
coworkers,116 but was suppressed in the stress-relieved
specimens compared to the recrystallized materials.
While a limited (<1%) amount of ductility was
reported in Mo–5Re alloys irradiated at temperatures <1000 K to 34 dpa,109,118,119 embrittlement of
Mo–(1–20)Re irradiated 723–1073 K in a fast reactor
up to 5 dpa and Mo–(13 and 47)Re irradiated 373–
673 K in a mixed spectrum reactor to 2 dpa has been
reported by Fabritsiev and Pokrovsky.62 The reduction in tensile strength, in many cases below the unirradiated values with no plasticity occurring in the
samples, was attributed to the hardening of the material by radiation-created defects along with RIS of
oxygen, nitrogen, and transmuted impurities to the
grain boundaries. The oxygen and nitrogen content in
the embrittled alloys was reported to be near 70 appm.
Irradiation hardening above 900 MPa was also
observed in Mo–41Re and Mo–47.5Re samples irradiated to 1.46 dpa at temperatures >1073 K.120 Failure
of Mo–41Re samples irradiated to 1.46 either prior to
yielding or after $5% elongation upon reaching
1600 MPa was observed at 1073 K (Tirr ¼ Ttest). Examples of the tensile curves for the two alloys in the
irradiated, 1100 h aged and as-annealed condition
tested at 1073 K is shown in Figure 20. Radiation


Radiation Effects in Refractory Metals and Alloys


203

1800
1600
Mo-41Re

Vickers hardness

1400
1200
1000

Mo-5Re

800
Unirrad.
600
400
LCAC-Mo

200
0
200

400

600
800
Irradiation temperature (K)


1000

1200

Mo (Nemoto et al.116 18-21 dpa)

Mo-5Re (Nemoto et al.116 18-21 dpa)

Mo-10Re (Nemoto et al.116 18–21 dpa)

Mo-41Re (Nemoto et al.116 18-21 dpa)

117 6.8-34 dpa)

Mo-5Re (Hasegawa et al.117 6.8-34 dpa)

Mo-5Re (Hasegawa et al.

117 6.8-34 dpa)

Mo-41Re (Hasegawa et al.

Figure 19 Vickers hardness as a function of neutron irradiation temperature and dose for LCAC-Mo, Mo–5Re, and
Mo–41Re alloys. Displacement damage levels are provided in the key. Reproduced from Nemoto, Y.; Hasegawa, A.; Satou,
M.; Abe, K.; Hiraoka, Y. J. Nucl. Mater. 2004, 324, 62–70; Hasegawa, A.; Ueda, K.; Satou, M.; Abe, K. J. Nucl. Mater. 1998,
258–263, 902–906.

2000
Mo-41Re


0.72 dpa

1.46 dpa

Mo-47.5Re

1500

1500 1.46 dpa
Stress (MPa)

0.72 dpa
0.72 dpa

1.46 dpa
1000

1.46 dpa

1000

0.72 dpa

500

500
1100 h
aged

1100 h

aged

Annealed
0

0

5

10

15
20
25
Strain (%)

30

35

0

5

10

15
20
25
Strain (%)


Annealed
30

35

40

Figure 20 Comparison of stress–strain curves for neutron irradiated, 1100 h and as-annealed Mo–41Re and Mo–47.5
Re samples at 1073 K (Tirr ¼ Ttest). Adapted from Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on
refractory metal alloys, ORNL/LTR/NR-PROM1/05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby,
J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater. 2007, 366, 388–406.


204

Radiation Effects in Refractory Metals and Alloys

Total elongation (%)

Ultimate tensile strength (MPa)

2500
Mo-41 (Busby)
Mo-Re (Fabrietsiev)
Mo-47.5Re (Busby)
Mo-5Re (Hasegawa)
Open symbols: unirrad. Closed symbols: irrad.

2000


Irrad. UTS
1500
Unirrad.
UTS

1000

500

Irrad. UTS
(brittle fracture)

15

10

5
Brittle fracture
0
600

700

800

900 1000 1100
Temperature (K)

1200


1300

1400

Figure 21 Tensile data comparisons of Mo–Re alloys detailing the upper limits for irradiation embrittlement. Adapted
from Busby, J. T.; Leonard, K. J.; Zinkle, S. J. Effects of neutron irradiation on refractory metal alloys, ORNL/LTR/NR-PROM1/
05-38; Oak Ridge National Laboratory: Oak Ridge, TN, Dec 2005; Busby, J. T.; Leonard, K. J.; Zinkle, S. J. J. Nucl. Mater.
2007, 366, 388–406.

hardening to levels over twice the as-annealed condition was observed for the alloys irradiated at 1223
and 1373 K; however, total elongation was between 4
and 12%. Analysis of the fractured surfaces of these
samples revealed intergranular failure, with the severity increasing with irradiation temperature. A comparison of mechanical property data of Mo–Re samples
from the sources discussed is shown in Figure 21.
The degree of RIS influencing the properties of
Mo–Re alloys varies with temperature, dose rate, and
total fluence. At temperatures <0.3 Tm, the recombination of vacancies and interstitials generated by
displacement damage dominates because of the limited defect mobility, and therefore RIS is not a factor.
At temperatures >0.5 Tm, a reduced driving force
for segregation occurs because of the high thermal
defect concentrations. At intermediate temperatures
($850–1430 K for Mo–Re), the radiation generated
point defects diffuse to defect sinks such as grain

boundaries or dislocations. Any preferential coupling
of vacancies or interstitial defects fluxes with solute
atoms, including transmuted species, will create
enrichment at the defect sinks. This is observed in
the nucleation of Re-rich phases in the microstructures of neutron-irradiated samples116,121 and the

degradation in mechanical properties and transition
to intergranular fracture in higher Re concentration
alloys.62,120 Further information on RIS can be found
in Chapter 1.18, Radiation-Induced Segregation.
Through modeling and experimental work, Erck
and Rehn123 showed that the degree of segregation
per dpa reaches a maximum for Mo–30 at.% Re
($45 wt% Re) near 1223 K and that for Mo–7 at.%
Re ($13 wt% Re) near 1473 K. While the Mo–5Re
alloys irradiated up to 20 dpa show some limited
ductility,109,118,119 the maximum irradiation temperatures were <1073 K and are therefore at or below
the lower temperature limit expected for RIS.


Radiation Effects in Refractory Metals and Alloys

1800
Tirr = 567 K, 12.3 dpa
Tirr = 882 K, 3.9 dpa
Tirr = 1143 K, 1.2 dpa
Tirr = 1143 K, 3.4 dpa
Unirradiated: longitudinal,
stress-relieved

1600
1400

Yield stress (MPa)

1200

1000
800
600
400
200

20

Total elongation (%)

The Mo–(41 and 47.5)Re alloys irradiated at 1073–
1373 K120 showed indications of RIS even at relatively
low damage levels, part of which may have been a
contribution of a thermal aging component, which in
the unirradiated as-aged Mo–41Re and Mo–47.5Re
showed increases in Re at the grain boundaries, leading to precipitation of s- and w-phases at the grain
boundaries in the 47.5Re containing alloy.115 Utilizing
Mo–Re alloys with a more moderate Re content may
improve the irradiation performance of these alloys,
especially when considering higher doses and/or longer irradiation times at temperatures at which thermal
precipitation effects may further compound RIS influences on mechanical properties.
Additional information on the fracture toughness
data for Mo–Re alloys is also needed. Preliminary
data by Scibetta and coworkers109 on precracked compact tension specimens of Mo–5Re showed reductions in fracture toughness from unirradiated values
of 17–23 MPa√m at room temperature and 623 K,
to $11 MPa√m for 0.35 dpa irradiated at 313 K, and
15 MPa√m for 0.29 dpa irradiated at 643 K. These
low irradiated values, at which no ductile crack growth
was observed in the specimens, are a concern.
Recent work examining the irradiated properties

of wrought, commercially available, ODS-Mo containing lanthanum oxide particles has shown promising
results.81,82,124 The fabrication methods produce a
microstructure consisting of elongated grains with
appreciable texturing and alignment of the oxide particles. The high degree of working associated with
fabrication produces a <2 mm grain size, which is
stabilized from growth by the ODS particles. Irradiation of the ODS-Mo up to 13.1 dpa at 567 K and
883–882 K produced an increase in yield strength
of 57–173%,124 while irradiation at 1143–1273 K
produced a 10–34% increase. The irradiated tensile
properties of ODS-Mo as a function of irradiation
temperature and dose from the work of Cockeram
and coworkers124 are shown in Figure 22. The
increases in radiation strength are comparable to the
higher limits for LCAC-Mo. The most striking result
of the ODS work is the improvement in the DBTT
for the irradiated samples.82,124 For 567 K irradiation to
12.3 dpa, the DBTT is $1073 K and is comparable
to that of LCAC-Mo and TZM (Figure 16). However, the DBTT for ODS-Mo irradiated to 13.1 dpa
at 833–882 K is $298 K, while that of LCAC-Mo is
573 and 973 K for TZM. For irradiation to 13.1 dpa at
1143–1209 K, the DBTT is $173 K, unchanged from
the nonirradiated material, while those of LCAC-Mo
and TZM are between 223–273 K.

205

15

10
Unirradiated

5

0
100

300

500

700
900
1100
Test temperature (K)

1300

1500

Figure 22 Yield stress and total elongation as a function
of test temperature for lanthanum oxide ODS-Mo neutron
irradiated up to 13 dpa at temperatures between 567 and
1209 K. Adapted from Cockeram, B. V.; Smith, R. W.;
Snead, L. L. J. Nucl. Mater. 2005, 346, 165–184.

The reduced susceptibility to irradiation embrittlement of ODS-Mo is in part due to the grain size
reduction and presence of the oxide particles. Reducing the distance of possible defect sinks such as grain
boundaries and offering additional sites such as the
oxide/matrix interface are particularly critical at lower
irradiation temperatures at which defect mobility is
limited. In addition, Cockeram and coworkers81,124

describe the fine but elongated grain structure as
enhancing the plain strain condition acting on each
plane of the lamina-shaped grains formed during the
fracture process, inducing larger plastic deformation
in irradiation-hardened material. This is also true
for the HP-LCAC-Mo containing the high aspect
ratio grains compared to other forms of LCAC-Mo
produced. Currently, no fracture data on the irradiated