4.12

Vanadium for Nuclear Systems

T. Muroga
National Institute for Fusion Science, Oroshi, Toki, Gifu, Japan

ß 2012 Elsevier Ltd. All rights reserved.

4.12.1

Introduction

391

4.12.2
4.12.3
4.12.4
4.12.5
4.12.6
4.12.7
4.12.8
4.12.9
4.12.10
4.12.11
4.12.12
4.12.13
References

Vanadium Alloys for Fusion Reactors
Compositional Optimization

Fabrication Technology
Fundamental Study on Impurity Effects
Thermal Creep
Corrosion, Compatibility, and Hydrogen Effects
Radiation Effects
Tritium-Related Issues
Development of Advanced Alloys
Critical Issues
Vanadium Alloy Development for Fusion Blankets
Summary

391
392
393
396
396
398
400
401
402
403
403
404
405

Abbreviations
DBTT
dpa
flibe
GTA

HFIR
HIP
IFMIF

Ductile–brittle transition temperature
Displacement per atom
Molten LiF-BeF2 salt mixture
Gas tungsten arc
High Flux Isotope Reactor
Hot isostatic pressing
International Fusion Materials Irradiation
Facility
IP
Imaging plate
ITER
International Thermonuclear
Experimental Reactor
LMFBR Liquid Metal Fast Breeder Reactor
MA
Mechanical alloying
PWHT Postweld heat treatment
RAFM Reduced activation ferritic/martensitic
REDOX Reduction–oxidation reaction
TBM
Test Blanket Module
TBR
Tritium breeding ratio
TEM
Transmission electron microscope


4.12.1 Introduction
Vanadium alloys were candidates for cladding
materials of Liquid Metal Fast Breeder Reactors
(LMFBR) in the 1970s.1 However, the development
was suspended mainly because of an unresolved issue

of corrosion with liquid sodium. Vanadium alloys
attracted attention in the 1980s again for use in fusion
reactors because of their ‘low activation’ properties.
At present, vanadium alloys are considered as one
of the three promising candidate low activation
structural materials for fusion reactors with reduced
activation ferritic/martensitic (RAFM) steels and
SiC/SiC composites. Overviews of vanadium alloys
for fusion reactor applications are available in the
recent proceedings papers of ICFRM (International
Conference on Fusion Reactor Materials).2–6 This
chapter highlights the recent progress in the development of vanadium alloys mainly for application
in fusion nuclear systems.

4.12.2 Vanadium Alloys for Fusion
Reactors
Various tritium breeding fusion blanket concepts
have been studied with different combinations of
structural materials, tritium breeding materials, and
cooling materials. Vanadium alloys have been used in
most cases with liquid lithium as the breeding and
cooling materials (self-cooled V/Li blankets) for
advanced concepts of DEMO (fusion demonstration power plant) and commercial fusion reactors.7,8
Because of high atomic density of Li atoms in liquid

Li relative to Li-ceramics, Li–Pb, and molten-salt
391


392

Vanadium for Nuclear Systems

Flibe, V/Li systems can obtain high tritium breeding
ratio (TBR) without using the neutron multiplier Be.
A neutronics calculation showed that ‘tritium self
sufficiency’ can be satisfied without Be both in
Tokamak and Helical reactor systems.9 Without the
necessity of using beryllium as a neutron multiplier,
the replacement frequency of the blanket will be
reduced because the blanket system is free from
the periodic replacement due to the lifetime of Be,
which can lead to enhanced plant efficiency.
V/Li blankets can be designed with a simple
structure as schematically shown in Figure 1. The
blanket is composed of Li cooling channels made of
vanadium alloys, reflectors, and a shielding area,
which is in contrast to more complex solid breeder
blankets that need a solid breeder zone, a neutron
multiplier beryllium zone, cooling channels using gas
or water, and tritium recovery gas flow channels in
addition to reflectors and shielding.
A self-cooled Li blanket using neutron multiplier
beryllium was also designed in the Russian program.10 This concept can downsize the blanket area
because of efficient tritium generation per zone.

However, the blanket structure must be more

complex than V/Li and new issues need to be solved
such as Li/Be compatibility.
General requirements for structural materials of
fusion blankets include dimensional stability, compatibility with breeder and coolants, high-temperature
strength and low-temperature ductility during irradiation. For vanadium alloys, issues concerning industrial
maturity such as developing large-scale manufacturing
technology need to be resolved.
Vanadium alloys could be a candidate structural
material for molten-salt Flibe (LiF–BeF2) blankets.
For this application, a concept was proposed to dissolve WF6 or MoF6 into Flibe for corrosion protection of the wall surfaces by precipitation of W or Mo
and reduction of the tritium inventory in vanadium
alloys by enhancing reaction from T2 to TF, which is
more highly soluble in Flibe than T2.11 The TBR of
Flibe/V blankets may be marginal, but the neutron
shielding capability for the superconductor magnet
systems may be superior relative to V/Li according to
neutronics investigation.12 In this system, precipitates
of Wor Mo formed as a result of reaction from T2 to TF
needs to be recovered from the flowing Flibe.
Table 1 summarizes the blanket concepts using
vanadium alloys with the advantages and critical
issues.

Flowing liquid lithium

4.12.3 Compositional Optimization

Superconducting

magnet

Neutron
D-T plasma

Shield

Coating
with W,
Be, or C

Reflector
Vanadium
alloy
structures

Blanket
Figure 1 Illustration of self-cooled Li blanket with
V–4Cr–4Ti structural material.

Table 1

Vanadium alloys potentially have low-induced activation characteristics, high-temperature strength,
and high thermal stress factors. For the optimization
of the composition, both major alloying elements and
minor impurities need to be controlled. For maintaining the low activation properties, use of Nb and
Mo, which used to be the candidate alloying elements
for application to LMFBR, need to be avoided.
Cr was known to increase the strength of vanadium
at high temperature and Ti was known to enhance

ductility of vanadium by absorbing interstitial impurities, mostly oxygen. However, excess Cr or Ti can

Breeding blanket concepts using vanadium alloys

Concept

V/Li

V/Be/Li

V/Flibe

Breeder and coolant
materials
Use of neutron multiplier Be
Advantages
Critical issues

Liquid Li

Liquid Li

Molten-salt Flibe

No
Simple structure
MHD coating, T
recovery from Li

Yes

High TBR
MHD coating, Li/Be compatibility,
T recovery from Li

No
Small MHD pressure drop
REDOX control, recovery of W or
Mo, increase in TBR


Vanadium for Nuclear Systems

lead to loss of ductility. Hence, optimization of Cr and
Ti levels for V–xCr–yTi has been investigated. It was
known that with x þ y > 10%, the alloys became
brittle6 as shown in Figure 2. With systematic efforts,
V–4Cr–4Ti has been regarded as the leading candidate. For low activation purposes, the level of Nb, Mo,
Ag, and Al needs to be strictly controlled.
Large and medium heats of V–4Cr–4Ti have
been made in the United States, Japan, and Russia.

100
1150 ЊC

50

±20 ЊC

DBTT (ЊC)


0
1100 ЊC

950 ЊC

–50
1000 ЊC

–100

Annealing
temperature
950 ЊC
1000 ЊC
1050 ЊC
1100 ЊC
1150 ЊC

–150
–200
–250

5

10

15
Cr + Ti (Wt %)

20


25

Figure 2 DBTT as a function of Cr þ Ti (wt%) of V–Cr–Ti
alloy for various annealing temperatures. Reproduced
from Zinkle, S. J.; Matsui, H.; Smith, D. L.; Rowcliffe, A. L.;
van Osch, E.; Abe, K.; Kazakov, V. A. J. Nucl. Mater. 1998,
258–263, 205–214, with permission from Elsevier.

393

An especially high-purity V–4Cr–4Ti ingot produced by the National Institute for Fusion Science
(NIFS) in collaboration with Japanese Universities
(NIFS-HEAT-1 and 2) showed superior properties in
manufacturing due to their reduced level of oxygen
impurities.4
Figure 3 compares the contact dose rate after use
in the first wall of a fusion commercial reactor for four
reference alloys. The full-remote and full-hands-on
recycle limits are shown to indicate the guideline for
recycling and reuse.13 SS316LN-IG (the reference
ITER structural material) will not reach the remoterecycling limit after cooling and hence the recycling is
not feasible. F82H (reference RAFM steel) and
NIFS-HEAT-2 behave similarly, but NIFS-HEAT-2
shows significantly lower dose rate before the 100year cooling. The dose rate of F82H and NIFSHEAT-2 reached a level almost two orders lower
than the remote-recycle limit by cooling for 100 and
50 years, respectively. The dose rate of SiC/SiC composites (assumed to be free from impurities because of
lack of reference composition) is much lower at
<1 year cooling, but slightly higher at >100 year
cooling relative to F82H and NIFS-HEAT-2.


4.12.4 Fabrication Technology
Figure 4 summarizes the microstructural evolution
during the breakdown process of NIFS-HEAT-2

105
104

Contact dose rate (Sv h–1)

103

Reduced
activation
ferritics (F82H)

102
101
100
10–1

V–4Cr–4Ti (NIFS-HEAT)

FFHR Li blanket
first wall
neutron 1.5 MW m–2
operation

SS316 for ITER
(SS316LN-IG)


Pure SiC/SiC

Full-remote
recycling

10–2
10–3
Full-hands-on
recycling

10–4
10–5
10–2

10–1

100
101
102
103
Cooling time after shutdown (years)

104

Figure 3 Contact dose after use in first wall of a fusion commercial reactor for four reference alloys. SS316LN-IG: the
reference ITER structural material F82H: reference reduced activation ferritic/martensitic steel NIFS-HEAT-2: reference
V–4Cr–4Ti alloy SiC/SiC: assumed to be impurity-free.



394

Vanadium for Nuclear Systems

Ingot

Hot forging
1423 K
Formation

Heat treatment
Hot/cold roll
1373 K/RT
973 K
1273 K
1373 K
1573 K
Ti-rich blocky precipitates (with N, O, C)
Elongation, band structure
Dissolution
Ti–O–C thin precipitates
Formation Coarsening Dissolution

V–C on GB

50 mm

50 mm

25 mm


1 mm

1 mm

1 mm

50 mm

Figure 4 Microstructural evolution during the breakdown process of V–4Cr–4Ti ingots. Reproduced from Muroga, T.;
Nagasaka, T.; Abe, K.; Chernov, V. M.; Matsui, H.; Smith, D. L.; Xu, Z. Y.; Zinkle, S. J. J. Nucl. Mater. 2002, 307–311, 547–554.

260
240
Vickers hardness (Hv)

ingots.4 Bands of small grains aligned along the rolling
direction were observed at the annealing temperature
below 1223 K. The grains became homogeneous at
$1223 K. The examination showed that optimization
of size and distribution of Ti-CON precipitates are
crucial for good mechanical properties of the V–4Cr–
4Ti products. Two types of precipitates were observed,
that is, the blocky and the thin precipitates. The blocky
precipitates formed during the initial fabrication process. The precipitates aligned along the working direction during the forging and the rolling processes
forming band structures, and were stable to 1373 K.
Since clustered structures of the precipitates result in
low impact properties, rolling to high reduction ratio
is necessary for making a thin band structure or homogenized distribution of the precipitates. The thin precipitates were formed at $973 K and disappeared at
1273–1373 K. At 1373 K, new precipitates, which were

composed of V and C, were observed at grain boundaries. They seem to be formed as a result of redistribution of C induced by the dissolution of the thin
precipitates. The impact of the inhomogeneous microstructure can influence the fracture properties.14
Figure 5 shows the hardness as a function of final
heat treatment temperature for three V–4Cr–4Ti
materials: NIFS-HEAT-1, NIFS-HEAT-2, and USDOE-832665 (US reference alloy).15 The hardness
has a minimum at 1073–1273 K, which corresponds
to the temperature range where formation of the thin
precipitates is maximized. With the heat treatment
higher than this temperature range, the hardness
increases and the ductility decreases because the

NIFS-HEAT-1
NIFS-HEAT-2
US-DOE 832665

220
200
180
160
140
V–4Cr–4Ti
120
200 400 600 800 1000 1200 1400 1600
Annealing temperature (K)

Figure 5 Vickers hardness as a function of annealing
temperature for NIFS-HEAT-1, NIFS-HEAT-2, and US-DOE
832665. Reproduced from Heo, N. J.; Nagasaka, T.;
Muroga, T. J. Nucl. Mater. 2004, 325, 53–60.


precipitates dissolve enhancing the level of C, N, and
O in the matrix. Based on the evaluation of various
properties in addition to the hardness as a function
of heat treatment conditions, the optimum heat treatment temperature of 1173–1273 K was suggested.
Plates, sheets, rods, and wires were fabricated minimizing the impurity pickup and maintaining grain
and precipitate sizes in Japanese, US, and Russian
programs. Thin pipes, including those of pressurized
creep tube specimens, were also successfully fabricated


Vanadium for Nuclear Systems

in Japan maintaining the impurity level, fine grain size,
and straight band precipitate distribution by maintaining a constant reduction ratio between the intermediate heat treatments.16 The fine-scale electron beam
welding technology was enhanced as a result of the
efforts for fabricating the creep tubes, including plugging of end caps.17 In the United States, optimum
vacuum level was found for eliminating the oxygen
pick-up during intermediate annealing to fabricate
thin-walled tubing of V–4Cr–4Ti.18 In Russia, fabrication technology is in progress for construction of a Test
Blanket Module (TBM) for ITER (International Thermonuclear Experimental Reactor).19
Joining of V–4Cr–4Ti by gas tungsten arc (GTA)
and laser welding methods was demonstrated. GTA

395

is a suitable technique for joining large structural
components. GTA welding technology for vanadium
alloys provided a significant progress by improving
the atmospheric control. The results are summarized
in Figure 6. Oxygen level in the weld metal was

controlled by combined use of plates of NIFSHEAT-1 (181 wppm O) or US-8332665 (310 wppm
O) and filler wire of NIFS-HEAT-1, US-8332665, or a
high-purity model alloy (36 wppm O). As demonstrated
in Figure 6, ductile–brittle transition temperature
(DBTT) of the joint and the oxygen level in the
weld metal had a clear positive relation. This motivated
further purification of the alloys for improvement of
the weld properties.20 Only limited data on irradiation effects on the weld joint are available at present.

15

Absorbed energy (J)

EU = 13 J

10

Plate/filler
NH1/HP
128 K

US/HP
183 K

NH1/NH1

US/US

188 K


320 K

5

0
50

100

150

200
250
Test temperature (K)

350

400

US/US

300

DBTT (K)

300

US/HP

200


NH1/NH1

100

NH1/HP
DBTT = +60 K/100 wppm O

0

0

50

100

150

200

250

300

350

400

Oxygen in weld metal (wppm)
Figure 6 Upper: Absorbed energy of Charpy impact tests of V–4Cr–4Ti weld joints as a function of test temperature for

various combinations of plates and fillers. Lower: DBTT of V–4Cr–4Ti weld joints as a function of oxygen level in the weld metal.
NH1, NIFS-HEAT-2 (O: 181 wppm); US, US-DOE 832665 (O: 310 wppm); HP, high-purity model V–4Cr–4Ti alloy
(O: 36 wppm). Reproduced from Nagasaka, T.; Grossbeck, M. L.; Muroga T.; King, J. F. Fusion Technol. 2001, 39, 664–668.


396

Vanadium for Nuclear Systems

The welding results in complete dissolution of TiCON precipitates and thus results in significant
increase in the level of C, O, and N in the matrix. In
such conditions, radiation could cause embrittlement.
Some TEM observations showed enhanced defect cluster density at the weld metals. However, the overall
evaluation of the radiation effects remains to be performed. Especially, elimination of radiation-induced
degradation byapplying appropriate conditions of postweld heat treatment (PWHT) is the key issue.
For the use of vanadium alloys as the blanket of
fusion reactors, the plasma-facing surfaces need to be
protected by armor materials such as W layers. Limited
efforts are, however, available for developing the
coating technology. A low pressure plasma-spraying
method was used for coating W on V–4Cr–4Ti for use
at the plasma-facing surfaces. The major issue for the
fabrication is the degradation of the vanadium alloy
substrates by oxidation during the coating processes.
Figure 7 shows the result of bending tests of the coated
samples. The crack was initiated within the W layer
propagating parallel to the interface and followed by
cracking across the interface. Thus, in this case, the
quality of W coating layer is the issue rather than the
property of the V–4Cr–4Ti substrate or the interface.

Hardening of substrate V–4Cr–4Ti by the coating
occurred but was shown to be in acceptable range.21
Figure 8 is a collection of the products from
NIFS-HEAT-2.

Research with model V–4Cr–4Ti alloys doped
with O and N provided information on the partitioning of O and N into the precipitates and matrix.
The density of the blocky precipitates and thin precipitates increased with N and O levels, respectively.
Figure 9 shows hardness as a function of N and
O levels in V–4Cr–4Ti after melting and annealing
at 1373 K for 1 h.22 Hardness after annealing at
1373 K, where only the blocky precipitates were
observed in the matrix, increased to a certain extent
with O level ($4.5 Hv/100 wppm O), but only very
weakly with N level ($0.9 Hv/100 wppm N). These
data suggest that, after the annealing, most of the
N is included in the blocky precipitates and stable
to $1373 K. On the other hand, O exists in the matrix,
the blocky and the thin precipitates, and the partitioning changes with the heat treatment. Thus, for the
purpose of the property control of V–4Cr–4Ti, the
level of N before the heat treatment is not so important but that of O is crucial. It is to be noted, however,
that N contamination during the operation can influence the properties of vanadium alloys seriously.
Fundamental information on the impurity distribution and interaction with solutes and dislocations
is obtained by serrated flow in tensile deformation as
shown in Figure 10. Temperature and stain rate dependence of the flow showed that the serrated flow above
673 K is related to C and O and above 773 K to N. Small
serration height at 673 K for NIFS-HEAT-1 (1–3 MPa)
relative to that of US-832665 ($9 MPa) was observed
and attributed to the difference in O level.23


4.12.5 Fundamental Study on
Impurity Effects
4.12.6 Thermal Creep
Effects of C, O, and N on the property of vanadium are
a long-standing research subject. However, research
into the effects of C, O, and N on V–4Cr–4Ti is limited.

Thermal creep is a potential factor which can determine the upper operation limit of vanadium alloys.

Crack
W
V–4Cr–4Ti
Intergranular fracture
500 µm

50 µm

10 µm
Figure 7 Cross-section of W coating on V–4Cr–4Ti after bending tests. Fracture started in the W coating layer.


Vanadium for Nuclear Systems

397

(mm)
f 4.57 ϫ 0. 25 t ϫ 400 mm
6.6 t
0.5 t


26 t
1.9 t

1.0 t

4.0 t
0.25 t

2d
8d
f 10 ϫ 0. 5 t ϫ 100 m m
Plates, sheets, wires, and rods

Thin pipes

20 mm
Creep tubes

W coating

NIFS-HEAT-2
(V–4Cr–4Ti)

0.5 mm

W coating by plasma spraying

Laser weld joint

5 mm


Figure 8 Collection of the V–4Cr–4Ti products manufactured by the Japanese program.

V–4Cr–4Ti, as-melted
V–4Cr–4Ti, 1373 K
Pure V, as-melted
Pure V, 1373 K

Vickers hardness (Hv)

300

250

200

150

100

50

0

200 400 600 800 1000 1200 0
Oxygen level (wppm)

200 400 600 800 1000 1200
Nitrogen level (wppm)


Figure 9 Vicker’s hardness as a function of O and N levels for V–4Cr–4Ti after melting and annealing at 1373 K for 1 h.
Reproduced from Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl. Mater. 2002, 307–311, 620–624.

Previously, uniaxial tensile creep tests and biaxial
pressurized creep tube tests were carried out in vacuum for evaluation of the creep deformation characteristics. Figure 11 shows summary of the creep
deformation rate as a function of applied stress.3 In
this plot, the creep data were described by a powerlaw equation24:
de=dt ¼ AðDGb=kT Þðs=GÞn

where de/dt is the creep rate, s is the applied stress,
D is the lattice diffusion coefficient, G is the shear
modulus, b is the Burgers vector, k is the Boltzmann
constant, T is the absolute temperature, and A is a
constant. The exponent of the function (n) changed
from <1 to >10 with the increase in the stress.
A new apparatus for biaxial creep testing in
Li provided opportunities for examining creep


398

Vanadium for Nuclear Systems

Stress (MPa)

200

0

10


30

20
Strain (%)

Figure 10 Tensile deformation curves of V–4Cr–4Ti at
various temperatures.

10-5
Uniaxial tests
310 wppm O

10-6

(de/dt)kT/DGb

n = 3.7

10-9
10-10

n = 4.3
n = 0.84

10-11
10-12

10-7


10-8

In vacuum
50 MPa
70 MPa
90 Mpa

10-9

10-10

0

2

4
6
8
Creep strain (%)

In lithium

10

12

Figure 12 Creep strain rate as a function of creep strain
for the same batch of NIFS-HEAT-2 creep tubes in vacuum
and Li environments. Modified from Li, M.; Nagasaka, T.;
Hoelzer, D. T.; et al. J. Nucl. Mater. 2007, 367–370, 788–793;

Fukumoto, K.; Nagasaka, T.; Muroga, T.; Nita, N.; Matsui, H.
J. Nucl. Mater. 2007, 367–370, 834–838.

n = 13

10-7
10-8

Creep strain rate (1 s–1)

10-6
1073 K
973 K
873 K
773 K
673 K
RT

surface hardening during exposure to Li. Further
investigation is necessary for understanding the environmental effects on impurity redistribution and creep
performance.
Microstructural observations of the creep tube
specimens tested at 1123 K showed free dislocations
and dislocation cell at 100 and 150 MPa, respectively.
This change of dislocation structure can cause the
change in power-law creep behavior.27

Biaxial tests
699 wppm O
10-3


10-2

4.12.7 Corrosion, Compatibility, and
Hydrogen Effects

s/G
Figure 11 Thermal creep deformation rate of V–4Cr–4Ti
as a function of applied stress for uniaxial and biaxial tests.
The definition of the terms and the function from which
n is extracted are indicated in the text. Reproduced from
Kurtz, R. J.; Abe, K.; Chernov, V. M.; Hoelzer, D. T.;
Matsui, H.; Muroga, T.; Odette, G. R. J. Nucl. Mater.
2004, 329–333, 47–55.

deformation in Li with that in vacuum.25 However,
the correlation of creep data is subject to the alloy
heat and manufacturing processes as well as test
methods and environments. Figure 12 shows the
comparison of the NIFS-HEAT-2 creep strain rate
versus creep strain data for tests in vacuum and Li
environments at 1073 K, for the same batch of NIFSHEAT-2 creep tubes.25,26 The figure clearly shows
reduced strain rate in Li environments. A possible
factor could be N pick-up from Li and the resulting

In a Li/V blanket, it is believed that the interior of the
wall needs to be coated with insulator ceramics for
mitigating the pressure drop caused by magnetohydrodynamic effects (see also Chapter 4.21, Ceramic
Coatings as Electrical Insulators in Fusion Blankets). Corrosion of vanadium alloys in liquid Li
might not be a concern if the entire inner wall is

covered with the insulating ceramic coating. However,
since the idea to cover the insulator ceramic coating
again with a thin vanadium or vanadium alloy layer was
presented for the purpose of preventing liquid lithium
from intruding into the cracks in the ceramics coating,
the corrosion of vanadium alloys in liquid lithium again
attracted attention. It is known that the corrosion of
vanadium alloys in liquid lithium is highly dependent
on the alloy composition and lithium chemistry. Especially, the N level influences the corrosion in complex
manners.28,29 Figure 13 shows a summary of the weight


Vanadium for Nuclear Systems

0.1

Ti
50

I

30 +11.9
20 +11

Ti:Cr = 2:1

+2.5

+6.7


+5.8

Weight gain (mg mm–2)

0.08
40

+1.2

–19.0

–22.0

V–4Cr–4Ti
V–4Cr–4Ti–0.5Si
V–4Cr–4Ti–0.5Al
V–4Cr–4Ti–0.5Y

0.06

0.04

II

+0.4

10

–8.2


–2.1

–19.7
–21.0

20

–47.4

0

–52.5
–26.4

30

Cr
40

50

Figure 13 The compositions of V–Ti–Cr alloys (wt%) with
increase (area I) and decrease (area II) of mass (g cmÀ2) after
holding of samples in Li at 973 K, 500 h. Reproduced from
Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P.
J. Nucl. Mater. 2000, 283–287, 1282–1286.

gain and loss in V–xCr–yTi systems in Li.30 High Ti
alloys showed a weight increase by forming a TiN layer
and high Cr alloys exhibited a weight loss as a result of

the dissolution of Cr–N complexes. As the boundary of
the two contradictory changes, Ti:Cr$2:1 was
observed.
Recently, a corrosion test using monometallic
thermal convection Li loop made of V–4Cr–4Ti
was conducted at 973 K for 2355 h. Because of the
temperature gradient, weight loss and weight gain of
V–4Cr–4Ti samples occurred at the hot leg and cold
leg, respectively. However, the loss rate corresponded
to only <1 mm yearÀ1 and the degradation of the
mechanical properties were shown to be small.31
V–4Cr–4Ti alloys have been developed mainly
for use in Li environments, which are extremely
reducing conditions. For the use of vanadium alloys in
oxidizing conditions, a different alloy optimization may
be necessary. The corrosion of vanadium alloys in oxidizing environments is of interest both for the performance of the pipe exterior out of the breeding blanket
and application in non-Li coolant systems such as gas
and water systems. Oxidation kinetics of vanadium
alloys were studied and showed either parabolic or
linear kinetics.32,33 As the surface oxide layer is not
formed or, if formed, not protective to the internal
oxidation, alloying with other oxide-formers is necessary for improvement. The addition of Si, Al, or Y was
shown to significantly suppress the weight gain during
exposure to air above 873 K as shown in Figure 14.34

773

873
973
Oxidation temperature (K)


1023

Figure 14 Weight gain of V–4Cr–4Ti with Si, Al, and
Y exposed to air for 1 h. At 1023 K, the weight gain was not
measured for V–4Cr–4Ti because the surface oxidized layer
melted. Reproduced from Fujiwara, M.; Natesan, K.;
Satou, M.; Hasegawa, A.; Abe, K. J. Nucl. Mater. 2002,
307–311, 601–604.

Natesan (BL-71 O:670 wppm)

50

DiStefano (US-832665 O:310 wppm)
DiStefano (Preoxidized US-832665 O:800 wppm)
Chen (SWIP-Heat O:900 wppm)

40
Total elongation (%)

+7.5

–2.4

V

Melted

0.02

10

399

Chen (NIFS-HEAT-2 O:158 wppm)

30

20

10

0

0

100 200 300 400 500 600
Hydrogen concentration (wppm)

700

Figure 15 Total elongation as a function of hydrogen
concentration for V–4Cr–4Ti alloys with different O levels.
Modified from DiStefano, J. R.; Pint, B. A.; DeVan, J. H.
J. Nucl. Mater. 2000, 283–287, 841–846; Chen, J. M.;
Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y. J. Nucl. Mater.
2004, 325, 79–86.

However, the addition of these elements was not effective in suppressing corrosion in water. Increase in Cr
level was shown to be effective, instead.

The effects of oxygen level on hydrogen embrittlement have been investigated. Figure 15 compares elongation versus hydrogen concentration for V–4Cr–4Ti


400

Vanadium for Nuclear Systems

alloys with various O levels. The loss of ductility by
hydrogen charging was shown to be enhanced
by impurity oxygen.35,36

4.12.8 Radiation Effects
A fair amount of data is available for radiation
response of vanadium alloys partly because they
were candidates of cladding materials of LMFBR.
For example, void swelling is known to be quite
small if the alloy contains Ti. However, data are
limited for V–4Cr–4Ti because this composition
was decided as the reference one for fusion only
recently. For this alloy, the feasibility issues of radiation effects are considered to be loss of ductility at
lower temperature, embrittlement enhanced by transmutant helium at high temperature, and irradiation
creep at intermediate to high temperature.
The mechanism of the loss of uniform elongation
of vanadium alloys at relatively low temperature
(<673 K) and low dose ($0.1 dpa) has been a longterm research subject. Microstructural observation
after tensile tests showed that radiation-induced
defect clusters were lost in layer structures and the
defect-free zones were accompanied by dislocation
channels as shown in Figure 16.37 This fact implies
flow localization during deformation. Although the

mechanism of the flow localization needs further investigation, it is inferred that interaction of dislocations with

radiation-induced defect clusters, precipitates, or complexes of the two species is responsible. If the precipitates, most likely Ti-CON, play the role in this process,
reduction of impurities in the matrix can improve the
properties. Figure 17 compares the uniform elongation
after irradiation for V–(4–5)Cr–(4–5)Ti alloys and those
with doping of Al, Si, and Y. The significant increase in
uniform elongation by the addition of Al, Si, and Y,
which are known as getters of interstitial impurities
such as O, N, and C in the matrix, suggests that the
reduction of the interstitial impurities in solution
enhances the radiation resistance.38 The effects of interstitial impurities on the formation of dislocation loops
and precipitates were investigated by ion irradiations.
Figure 18 shows temperature dependence of the
densities of loops and precipitates.39 The loop density
was not influenced by O level, but the precipitate
density increased with O level below 973 K.
Helium embrittlement is a critical issue, which is
thought to determine the upper temperature limit
for vanadium alloys. Past experimental evaluations
of the helium effects involved uncertainties because
controlled generation of helium during irradiation in
a similar manner to that in fusion condition has been
quite difficult. As a result, the past evaluation of the
helium effects varied from weak to very strong.3 The
Dynamic Helium Charging Experiment (DHCE)
using fission reactors40 is one of the few potential
neutron irradiation experiments with controlled
variation of He/dpa ratio including typical fusion


Load-elongation curves for V–4Cr–4Ti
irradiated in HFBR to 0.5 dpa
Engineering stress (MPa)

700
600
500

383 K

543 K

Ttest ~ Tirr
598 K
693 K

400
300
200
100
0
0

g = 011

0.05
0.1
0.15
0.2
0.25

0.3
Normalized crosshead displacement (mm mm–1)
Ttest = Tirr = 543 K

100 nm
Figure 16 Tensile test curves for V–4Cr–4Ti irradiated in HFBR to 0.5 dpa and microstructure after the tensile test.
Reproduced from Rice, P. M.; Zinkle, S. J. J. Nucl. Mater. 1998, 258–263, 1414–1419.


Vanadium for Nuclear Systems
Annealing conditions: 900–1125 ЊC, 3.6 or 7.2 ks

Uniform elongation (%)

12
10
8

ATR, 0.7–4.7 dpa
FFTF, EBR-II, 10–54 dpa
Loomis et al.,51 FFTF, 13–33dpa
Zinkle et al.,52 EBR-II, 4 dpa
Tsai et al.,53 ATR, 4.1–4.3 dpa
Tsai et al.,54 BOR-60, 17–19 dpa
Snead et al.,56 HFBR, 0.5 dpa
Chung et al.,55 HFIR, 10 dpa

Creep strain (%)

V–4.8Ti–4.0Cr–Si, Al, Y

V–3.8Ti–5.9Cr–Si, Al, Y
V–(4–5)Cr–(4–5)Ti

Test temperature–Irradiation temperature

US-832665 698 K-Li (HFIR)
NIFS-HEAT-2 698 K-Li (HFIR)
NIFS-HEAT-2 731 K-Na (JOYO)

0.2

0.1

6
0
4
2
0

0

100

200
300
400
500
Irradiation temperature (ЊC)

600


700

Figure 17 Uniform elongations as a function of irradiation
temperature for V–(4–5)Cr–(4–5)Ti alloys and those with
addition of Si, Al, and Y. Reproduced from Satou, M.;
Chuto T.; Abe, K. J. Nucl. Mater. 2000, 283–287, 367–371.

1025
O:389, N:14 wppm (loop)
O:28, N:27 wppm (loop)
O:389, N:14 wppm (precipitates)
O.28, N:27 wppm (precipitates)

1024
Defect density (m–3)

401

50

150
100
Applied stress (MPa)

200

Figure 19 Creep strain as a function of applied stress for
V–4Cr–4Ti (US-832665 and NIFS-HEAT-2) irradiated in Li
(HFIR) and Na (JOYO) environments. The creep strain was

normalized as that at two displacements per atom.
Reproduced from Fukumoto, K.; Narui, M.; Matsui, H.; et al.
J. Nucl. Mater. 2009, 386–388, 575–578.

environments. The data also compare the performances of US and Japanese reference alloys.41 It
was found that the creep strain rate exhibited a linear
relationship with the effective stress up to 150 MPa at
$700 K and the differences with the environments
and the heats are small.

4.12.9 Tritium-Related Issues

1023

1022

1021

1020
400

0

500

600

700 800 900
Temperature (K)


1000 1100

Figure 18 Densities of dislocation loops and precipitates
as a function of irradiation temperature for two V–4Cr–4Ti alloys
with different O and N levels (0.1 dpa by Cu ion irradiation).
Reproduced from Watanabe, H.; Suda, M.; Muroga, T.;
Yoshida, N. J. Nucl. Mater. 2002, 307–311, 408–411.

conditions. DHCE is highly anticipated as a potential
method to extend our understanding of the helium
effects. However, for conclusive evaluation, a 14 MeV
neutron source is certainly necessary.
The irradiation creep tests have made progress
recently, partly because of the progress in fabricating
high quality pressurized creep tube specimens with
reduced impurity levels. Figure 19 shows the normalized creep strain as a function of applied stress
by irradiation in HFIR and JOYO in Li and Na

In the blanket, the tritium inventory is not considered
to be the issue once liquid Li is used as the breeding
and cooling materials owing to the high hydrogen
solubility of Li. The behavior of hydrogen and its
isotopes in vanadium alloys is a concern for tritium
retention in the first wall. Deuterium retention
of V–4Cr–4Ti was investigated by deuterium ion
implantation followed by thermal desorption, in
comparison with other candidate first wall materials.
The study showed that the retained amount at
380 K was one and two orders of magnitude larger
than graphite and tungsten, respectively. For the irradiation at 773 K, the retained amount was almost the

same as that of graphite and one order larger than
tungsten.42 Surface composition was also known to
influence the hydrogen transport. For example, the
rate of absorption was highly influenced by prior heat
treatment, inducing Ti surface segregation.43
Recent progress in detecting tritium by means of
imaging plate (IP) enhanced the understanding of the
tritium behavior in vanadium alloys. Figure 20 compares IP images of cold rolled V–4Cr–4Ti and pure
V after tritium charging. Tritium is preferentially
absorbed in Ti-rich precipitates that have a band
structure to the rolling direction.44


402

Vanadium for Nuclear Systems

Pure Vanadium

Rolling direction

V–4Cr–4Ti (NIFS-HEAT-2)

Low

High

2 mm

Low


High

2 mm

Figure 20 Distribution of tritium measured by imaging plates for cold-rolled V–4Cr–4Ti and pure vanadium. The tritium
was charged by gas absorption. After Homma, H.; Hatano, Y.; Daifuku, H.; et al. J. Nucl. Mater. 2007, 367–370, 887–891.

4.12.10 Development of Advanced
Alloys
The performance of structural materials can strongly
influence the blanket design. Especially, the operation temperature window and expected lifetime are
the key parameters. Increase in the upper operation
temperature limit can enhance the blanket operation
temperature and thus plant efficiency. Therefore,
enhancing the high-temperature strength is the key
issue for improving the performance of the blanket
and thus the attractiveness of the fusion power systems. For this purpose, efforts have been made to
develop advanced vanadium alloys with potential
use at higher temperature.
One of the relatively simple ways to enhance the
strength of the alloy is to change the thermal and
mechanical treatment of the alloys. Especially, formation of a high density of precipitates can
strengthen the alloy. Figure 21 shows microstructure
and hardness of V–4Cr–4Ti as a function of the
temperature of reheating for 1 h after annealing at
1373 K for 1 h. The annealing at 1373 K dissolves
most of the thin precipitates and the reheating can
form new precipitates. By choosing an appropriate
reheating temperature (873–973 K), the materials can

be strengthened by the high density of fine precipitates.
However, the strengthening by this treatment will be
lost at >973 K because of the coarsening of the precipitates. To prevent the coarsening, cold work was

applied to the specimens. Figure 22 shows the minimum creep rate for standard V–4Cr–4Ti and solution
annealed, aged, and cold-worked V–4Cr–4Ti. Suppression of the creep rate occurred at 1073 K but
only with relatively high stresses.45 Microstructural
analysis showed that the suppressive role of coldwork-induced dislocations was lost during the creep
deformation by the change in the nature of the dislocations from sessile ah100i type to gliding a/2h111i
type.46 Further efforts are being made, for example,
to cold-work followed by aging (strain-aging-induced
strengthening).
High-temperature strength of V–Cr–Ti alloys can
be enhanced by increasing the Cr level. However,
high Cr alloys have low ductility and fabricability
issues. Recent detailed survey in V–xCr–4Ti alloys
showed that the strength at high temperature
increases with a small change in the DBTT with the
Cr level at $7%.47
High-strength vanadium alloys were made by addition of Y, O, and N to vanadium followed by mechanical alloying (MA) and hot isostatic pressing (HIP). The
addition of Y, O, and N was intended to enhance
mechanical properties by dispersion of Y2O3 and YN
and scavenging O and N from the matrix. Alloys produced by optimization of the processes had small grains
and homogeneously dispersed particles and showed
higher tensile strength than those of NIFS-HEATs
with moderate uniform elongation, both at room temperature and 1073 K as shown in Figure 23.48 Fine


Vanadium for Nuclear Systems


450
Larger grain
Yield stress (MPa)

873 K

403

Perfectly dissolved
1373 K

400

Solution

350
Precipitation
300

200 nm
250

As solution 873 973 1073 1173 1273 1373
heat treated Reheat temperature (K)

973 K

1073 K

1273 K


1173 K

Start to dissolve
Coarse precipitates
(low density)

Fine precipitates

Figure 21 Hardness and microstructure of V–4Cr–4Ti as a function of reheating temperature for 1 h after annealing at
1373 K for 1 h.

4.12.11 Critical Issues

Minimum creep rate (s–1)

100–4

NIFS-HEAT-2
10–5

800 ЊC
10–6

750 ЊC

10–7

700 ЊC


10–8

10–9
100

120

140

160

SAACW
STD
180

200 220 240 260 280 300

Stress (MPa)
Figure 22 Minimum creep rates as a function of applied
stress for V–4Cr–4Ti with standard heat treatment
(1273 K for 1 h: STD) and precipitate-hardening heat
treatment (1373 K for 1 h, 873 K for 20 h, and cold rolled:
SAACW). Reproduced from Chen, J. M.; Nagasaka, T.;
Muroga, T.; Qiu, S. Y.; Li, C.; Nita, N. J. Nucl. Mater. 2008,
374, 298–303.

grain and oxide dispersion increased high-temperature
strength and inhibited formation of interstitial loops
in the matrix by neutron irradiation because of the
enhanced defect sinks. Thus, mechanically alloyed

vanadium alloys have the potential to extend both
low- and high-temperature operation limits.
Other efforts to improve high-temperature
strength of vanadium alloys include strengthening
by internal oxidation.49

With the recent progress in the fabrication technology,
the number of critical issues for the development of
vanadium alloys for fusion reactors has been reduced.
The remaining critical issues are thermal and irradiation creep, transmutant helium effects on high temperature mechanical properties, and radiation effects on
fracture properties. The effect of helium, particularly,
is still uncertain and can be evaluated precisely only
with the use of 14 MeV neutrons. This fact highly
motivates the construction of a 14 MeV neutron source.
With the progress of the properties of vanadium
alloys, the blanket concepts using the alloy become
more attractive. Extension of the operation temperature window and lifetime of vanadium alloys
contribute to the improvement of the quality of the
blanket. Therefore, exploration of advanced vanadium alloys from the current reference alloy is a
valuable challenge for enhancing the expected performance, and then attractiveness, of fusion reactors.

4.12.12 Vanadium Alloy
Development for Fusion Blankets
In the fusion materials development strategy, the
candidate structural materials are categorized into
reference and advanced materials. As the reference


404


Vanadium for Nuclear Systems

800

TTest = 298 K

V–1.0 vol% Y2O3–0.7 vol% YN
NIFS-HEAT-1

Stress (MPa)

600

400

TTest = 1073 K

200

0

Strain (%)

20%

400 nm

Figure 23 Tensile test curves and microstructure of V–Y2O3–YN produced by mechanical alloying (MA) in comparison with
V–4Cr–4Ti (NIFS-HEAT-1). After Kuwabara, T.; Kurishita, H.; Hasegawa, M. J. Nucl. Mater. 2000, 283–287, 611.


Flexible support

Be
Li
WC

First wall

Figure 24 A layout of the V/Be/Li test blanket module for
International Thermonuclear Experimental Reactor
proposed by Russia. After Kirillov, I. R.; Shatalov, G. E.;
Strebkov, Y. S.; the RF TBM Team. Fusion Eng. Des. 2006,
81, 425–432.

the advanced materials, which will contribute to
increasing attractiveness of the fusion system in
terms of cost of electricity and environmental benignness. It is recognized that the development of the
advanced materials must also be enhanced now due
to the long lead time necessary for their development.
It should also be noted that vanadium alloys are the
only nonferromagnetic and ductile materials of the
three candidates. If the impact of the ferromagnetism
of the RAFM on plasma operation should be unavoidable and the brittleness of SiC/SiC should be determined unaccepted by design studies, vanadium alloys
could be the only candidate of low activation structural
materials for fusion reactors.
As shown in the summary of critical issues, a
14 MeV neutron source is highly necessary for the
qualification of vanadium alloys. IFMIF (International Fusion Materials Irradiation Test Facility, a
14 MeV neutron source) is under design and is recognized to be essential for developing structural materials
for fusion reactors. The TBM to be installed in ITER

is also considered to be an important milestone for
technological integration. Figure 24 shows the design
of the V/Li TBM in ITER proposed by Russia.50 The
development of vanadium alloys is planned to proceed with IFMIF for qualification of the alloy and
ITER-TBM for technology integration, in addition
to fundamental studies using fission reactors, etc.

4.12.13 Summary
materials, RAFM steels were selected because they
have the most matured industrial infrastructure.
Development of the reference materials is crucial
for the realization of DEMO (fusion demonstration
power plant) in a timely manner. On the other hand,
vanadium alloys and SiC/SiC were nominated as

As to the application in nuclear systems, vanadium
alloys were once candidate cladding materials for
LMFBR, but, at present, are considered mostly
as candidate low activation structural materials for
fusion reactors.


Vanadium for Nuclear Systems

Vanadium alloys, with the reference composition
of V–4Cr–4Ti, are one of the three candidate
low activation structural materials with RAFM and
SiC/SiC. They are the only nonferromagnetic and
ductile materials of the three candidates and thus are
promising for advanced structural materials of fusion

reactors. The self-cooled liquid lithium blanket using
structural materials of vanadium alloys is an attractive
concept because of the high heat transfer capability,
high-temperature operation, simple structure, high tritium breeding capability, and low tritium leakage.
Recent progress, especially in the fabrication
technologies, has successfully reduced the number
of critical issues enhancing the feasibility of the alloys
for fusion application. Major remaining issues of
vanadium alloys are thermal and irradiation creep,
transmutant helium effects on mechanical properties,
and radiation effects on fracture properties. For conclusive characterization of the helium effects, the use
of IFMIF is essential.
Efforts are also being made to develop advanced
vanadium alloys to extend the temperature window
and lifetime of vanadium alloys in fusion reactor
environments.

17.
18.
19.
20.
21.
22.
23.
24.

25.
26.
27.
28.

29.
30.
31.
32.
33.
34.
35.

References
1.
2.
3.
4.
5.
6.
7.
8.
9.
10.
11.
12.
13.
14.
15.
16.

Suzuki, T.; Noda, T.; Iwao, N.; Kainuma, T.; Watanabe, R.
J. Nucl. Mater. 1976, 62, 205–212.
Kurtz, R. J.; Abe, K.; Chernov, V. M.; et al. J. Nucl. Mater.
2000, 283–287, 70–78.

Kurtz, R. J.; Abe, K.; Chernov, V. M.; et al. J. Nucl. Mater.
2004, 329–333, 47–55.
Muroga, T.; Nagasaka, T.; Abe, K.; et al. J. Nucl. Mater.
2002, 307–311, 547–554.
Muroga, T.; Chen, J. M.; Chernov, V. M.; et al. J. Nucl.
Mater. 2007, 367–370, 780–787.
Zinkle, S. J.; Matsui, H.; Smith, D. L.; et al. J. Nucl. Mater.
1998, 258–263, 205–214.
Najmabadi, F.; et al. Fusion Eng. Des. 1997, 38, 3–25.
Sokolov, Y. A.; et al. Fusion Eng. Des. 1998, 41, 525–529.
Tanaka, T.; Muroga, T.; Sagara, A. Fusion Sci. Technol.
2005, 47, 530–534.
Kolbasov, B. N.; Belyakov, V. A.; Bondarchuk, E. N.; et al.
Fusion Eng. Des. 2008, 83, 870–876.
Muroga, T.; Tanaka, T.; Li, Z.; Sagara, A.; Sze, D. K. Fusion
Sci. Technol. 2007, 52, 682–686.
Sze, D. K.; McKarthy, K. A.; Sawan, M. E.; Tillack, M. S.;
Ying, A. Y.; Zinkle, S. J. Fusion Technol. 2001, 39, 746–750.
Dolan, T. J.; Butterworth, G. J. Fusion Technol. 1994, 26,
1014–1018.
Donahue, E. G.; Odette, G. R.; Lucas, G. E. J. Nucl. Mater.
2000, 283–287, 518–522.
Heo, N. J.; Nagasaka, T.; Muroga, T. J. Nucl. Mater. 2004,
325, 53–60.
Nagasaka, T.; Muroga, T.; Iikubo, T. Fusion Sci. Technol.
2003, 44, 465–469.

36.
37
38.

39.
40.
41.
42.
43.
44.
45.
46.
47.

48.
49.
50.
51.

405

Fukumoto, K.; Matsui, H.; Narui, M.; Nagasaka, T.;
Muroga, T. J. Nucl. Mater. 2004, 335, 103–107.
Rowcliffe, A. F.; Hoelzer, D. T.; Kurtz, R. J.; Young, M.
J. Nucl. Mater. 2007, 367–370, 839–843.
Chernov, V. M.; et al. Nucl. Fusion 2007, 47, 839–848.
Nagasaka, T.; Grossbeck, M. L.; Muroga, T.; King, J. F.
Fusion Technol. 2001, 39, 664–668.
Nagasaka, T.; Muroga, T.; Noda, N.; Kawamura, M.;
Ise, H. Fusion Sci. Technol. 2005, 47, 876–880.
Heo, N. J.; Nagasaka, T.; Muroga, T.; Matsui, H. J. Nucl.
Mater. 2002, 307–311, 620–624.
Fukumoto, K.; Yamamoto, T.; Nakao, N.; Takahashi, S.;
Matsui, H. J. Nucl. Mater. 2002, 307–311, 610–614.

Li, M.; Hoelzer, D. T.; Grossbeck, M. L.; Rowcliffe, A. F.;
Zinkle, S. J.; Kurtz, R. J. J. Nucl. Mater. 2009, 386–388,
618–621.
Li, M.; Nagasaka, T.; Hoelzer, D. T.; et al. J. Nucl. Mater.
2007, 367–370, 788–793.
Fukumoto, K.; Nagasaka, T.; Muroga, T.; Nita, N.;
Matsui, H. J. Nucl. Mater. 2007, 367–370, 834–838.
Gelles, D. S.; Toloczko, M. B.; Kurtz, R. J. J. Nucl. Mater.
2007, 367–370, 869–875.
Chopra, O. K.; Smith, D. L. J. Nucl. Mater. 1988,
155–157, 683.
Evtikhin, V. A.; Lyublinski, I. E.; Vertkov, A. V. J. Nucl.
Mater. 1998, 258–263, 1487–1491.
Eliseeva, O. I.; Fedirko, V. N.; Chernov, V. M.; Zavialsky, L. P.
J. Nucl. Mater. 2000, 283–287, 1282–1286.
Pint, B. A.; Pawel, S. J.; Howell, M.; et al. J. Nucl. Mater.
2009, 386–388, 712–715.
DiStefano, J. R.; DeVan, J. H. J. Nucl. Mater. 1997, 249,
150–158.
Natesan, K.; Uz, M. Fusion Eng. Des. 2000, 51–52, 145–152.
Fujiwara, M.; Natesan, K.; Satou, M.; Hasegawa, A.;
Abe, K. J. Nucl. Mater. 2002, 307–311, 601–604.
Chen, J. M.; Muroga, T.; Qiu, S.; Xu, Y.; Den, Y.; Xu, Z. Y.
J. Nucl. Mater. 2004, 325, 79–86.
DiStefano, J. R.; Pint, B. A.; DeVan, J. H. J. Nucl. Mater.
2000, 283–287, 841–846.
Rice, P. M.; Zinkle, S. J. J. Nucl. Mater. 258–263, 1414–1419.
Satou, M.; Chuto, T.; Abe, K. J. Nucl. Mater. 2000,
283–287, 367–371.
Watanabe, H.; Suda, M.; Muroga, T.; Yoshida, N. J. Nucl.

Mater. 2002, 307–311, 408–411.
Smith, D. L.; Matsui, H.; Greenwood, L.; Loomis, B. A.
J. Nucl. Mater. 1988, 155–157, 1359–1363.
Fukumoto, K.; Narui, M.; Matsui, H.; et al. J. Nucl. Mater.
2009, 386–388, 575–578.
Yamauchi, Y.; Yamada, T.; Hirohata, Y.; Hino, T.;
Muroga, T. J. Nucl. Mater. 2004, 329–333, 397–400.
Hayakawa, R.; Hatano, Y.; Fukumoto, K.; Matsui, H.;
Watanabe, K. J. Nucl. Mater. 2004, 329–333, 411–415.
Homma, H.; Hatano, Y.; Daifuku, H.; et al. J. Nucl. Mater.
2007, 367–370, 887–891.
Chen, J. M.; Nagasaka, T.; Muroga, T.; Qiu, S. Y.; Li, C.;
Nita, N. J. Nucl. Mater. 2008, 374, 298–303.
Muroga, T.; Nagasaka, T.; Chen, J. M.; Li, Y. F.;
Watanabe, H. J. Nucl. Mater. 2009, 386–388, 606–609.
Sakai, K.; Satou, M.; Fujiwara, M.; Takahashi, K.;
Hasegawa, A.; Abe, K. J. Nucl. Mater. 2004, 329–333,
457–461.
Kuwabara, T.; Kurishita, H.; Hasegawa, M. J. Nucl. Mater.
2000, 283–287, 611.
Tyumentsev, A. N.; Korotaev, A. D.; Pinzhin, Y. P.; et al.
J. Nucl. Mater. 2007, 367–370, 853–857.
Kirillov, I. R.; Shatalov, G. E.; Strebkov, Y. S. The RF TBM
Team. Fusion Eng. Des. 2006, 81, 425–432.
Loomis, B. A.; Nowicki, L. J.; Smith, D. L. J. Nucl. Mater.
1994, 212–215, 790.


406


Vanadium for Nuclear Systems

52. Zinkle, S. J.; Alexander, D. J.; Robertson, J. P.; et al.
Eatherly. Fusion Mater. 1996, 73. DOE/ER-0313/21.
53. Tsai, H.; Nowicki, L. J.; Billone, M. C.; Chung, H. M.
Smith, D. L. Fusion Mater. 1997, 70. DOE/ER-0313/23.
54. Tsai, H.; Gazda, J.; Nowicki, L. J.; Billone, M. C.; Smith, D. L.
Fusion Mater. 1998, 15. DOE/ER-0313/24.

55.
56.

Chung, H. M.; Nowicki, L. J.; Smith, D. L. Fusion Mater.
1997, 29. DOE/ER-0313/22.
Snead, L. L.; Zinkle, S. J.; Alexander, D. J.; Rowcliffe, A. F.;
Robertson, J. P.; Eatherly, W. S. DOE/ER-0313/23 (Dec.
31, 1997) 81. Available at />fusionreactor/dec97.shtml