Comprehensive nuclear materials 4 12 vanadium for nuclear systems

Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems Comprehensive nuclear materials 4 12 vanadium for nuclear systems

T Muroga

National Institute for Fusion Science, Oroshi, Toki, Gifu, Japan

ß 2012 Elsevier Ltd All rights reserved.

Abbreviations

DBTT Ductile–brittle transition temperature

dpa Displacement per atom

flibe Molten LiF-BeF2 salt mixture

GTA Gas tungsten arc

HFIR High Flux Isotope Reactor

HIP Hot isostatic pressing

IFMIF 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 devel-opment 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 demonstra-tion 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

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.9Without 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

pro-gram.10This 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, compat-ibility with breeder and coolants, high-temperature strength and low-temperature ductility during irradia-tion 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 dis-solve WF6or MoF6into Flibe for corrosion protec-tion of the wall surfaces by precipitaprotec-tion of W or Mo and reduction of the tritium inventory in vanadium alloys by enhancing reaction from T2to 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.12In this system, precipitates

of Wor Mo formed as a result of reaction fromT2to TF needs to be recovered from the flowing Flibe

vanadium alloys with the advantages and critical issues

4.12.3 Compositional Optimization Vanadium alloys potentially have low-induced acti-vation 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 main-taining 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 impu-rities, mostly oxygen However, excess Cr or Ti can Table 1 Breeding blanket concepts using vanadium alloys

Breeder and coolant

materials

Advantages Simple structure High TBR Small MHD pressure drop Critical issues MHD coating, T

recovery from Li

MHD coating, Li/Be compatibility,

T recovery from Li

REDOX control, recovery of W or

Mo, increase in TBR

D-T plasma

Reflector Neutron

Superconducting magnet

Vanadium alloy structures

Blanket

Coating with W,

Be, or C

Flowing liquid lithium

Shield

Figure 1 Illustration of self-cooled Li blanket with

V–4Cr–4Ti structural material.

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

brittle6as shown inFigure 2 With systematic efforts,

V–4Cr–4Ti has been regarded as the leading

candi-date 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

An especially high-purity V–4Cr–4Ti ingot pro-duced 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

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 remote-recycling limit after cooling and hence the remote-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 100-year cooling The dose rate of F82H and NIFS-HEAT-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 com-posites (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

during the breakdown process of NIFS-HEAT-2

950 ⬚C

950 ⬚C

1000 ⬚C

1000 ⬚C

1050 ⬚C

1100 ⬚C

1100 ⬚C

1150 ⬚C

1150 ⬚C

Annealing temperature

Cr + Ti (Wt %)

±20 ⬚C

–250

–200

–150

–100

–50

0

50

100

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.

10 –2

10 –5

10–4

10–3

10 –2

10 –1

100

10 1

10 2

103

104

10 5

Cooling time after shutdown (years)

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

Full-hands-on recycling

Full-remote recycling

Reduced activation ferritics (F82H)

SS316 for ITER (SS316LN-IG) Pure SiC/SiC

V–4Cr–4Ti (NIFS-HEAT)

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.

ingots.4Bands 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

pro-cess The precipitates aligned along the working

direc-tion 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

homo-genized distribution of the precipitates The thin

pre-cipitates were formed at973 K and disappeared at

1273–1373 K At 1373 K, new precipitates, which were

composed of V and C, were observed at grain

bound-aries They seem to be formed as a result of

redistri-bution of C induced by the dissolution of the thin

precipitates The impact of the inhomogeneous

micro-structure can influence the fracture properties.14

heat treatment temperature for three V–4Cr–4Ti

materials: NIFS-HEAT-1, NIFS-HEAT-2, and

US-DOE-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

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 treat-ment temperature of 1173–1273 K was suggested Plates, sheets, rods, and wires were fabricated mini-mizing 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

Ingot Hot forging

1423 K

Hot/cold roll

1373 K/RT

Heat treatment

973 K 1273 K 1373 K 1573 K Ti-rich blocky precipitates (with N, O, C)

Formation Elongation, band structure Dissolution

Ti–O–C thin precipitates Formation Coarsening Dissolution

V–C on GB

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.

120 140 160 180 200 220 240 260

200 400 600 800 1000 1200 1400 1600

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

V–4Cr–4Ti

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.

in Japan maintaining the impurity level, fine grain size,

and straight band precipitate distribution by

maintain-ing a constant reduction ratio between the

intermedi-ate heat treatments.16 The fine-scale electron beam

welding technology was enhanced as a result of the

efforts for fabricating the creep tubes, including

plug-ging 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.18In Russia,

fabrica-tion technology is in progress for construcfabrica-tion of a Test

Blanket Module (TBM) for ITER (International

Ther-monuclear Experimental Reactor).19

Joining of V–4Cr–4Ti by gas tungsten arc (GTA)

and laser welding methods was demonstrated GTA

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

controlled by combined use of plates of NIFS-HEAT-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

(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 irradia-tion effects on the weld joint are available at present

0 5 10 15

Test temperature (K)

US/US

320 K

EU= 13 J

NH1/NH1

188 K

US/HP

183 K

128 K NH1/HP

0 100 200 300

Oxygen in weld metal (wppm) NH1/HP

US/HP

NH1/NH1

US/US

DBTT = +60 K/100 wppm O

Plate/filler

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.

The welding results in complete dissolution of

Ti-CON 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

clus-ter density at the weld metals However, the overall

evaluation of the radiation effects remains to be

per-formed Especially, elimination of radiation-induced

degradation by applying appropriate conditions of

post-weld 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

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

NIFS-HEAT-2

4.12.5 Fundamental Study on

Impurity Effects

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

Research with model V–4Cr–4Ti alloys doped with O and N provided information on the partition-ing of O and N into the precipitates and matrix The density of the blocky precipitates and thin pre-cipitates increased with N and O levels, respectively

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

to1373 K On the other hand, O exists in the matrix, the blocky and the thin precipitates, and the partition-ing 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 impor-tant but that of O is crucial It is to be noted, however, that N contamination during the operation can influ-ence the properties of vanadium alloys seriously Fundamental information on the impurity dis-tribution and interaction with solutes and dislocations

is obtained by serrated flow in tensile deformation as shown inFigure 10 Temperature and stain rate depen-dence 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

Thermal creep is a potential factor which can deter-mine the upper operation limit of vanadium alloys

Crack

Intergranular fracture

W V–4Cr–4Ti

10 µm

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

Previously, uniaxial tensile creep tests and biaxial

pressurized creep tube tests were carried out in

vac-uum for evaluation of the creep deformation

charac-teristics Figure 11 shows summary of the creep

deformation rate as a function of applied stress.3In

this plot, the creep data were described by a

power-law equation24:

de=dt ¼ AðDGb=kTÞðs=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

Nitrogen level (wppm) 50

100 150 200 250 300

0 200 400 600 800 1000 1200 0 200 400 600 800 1000 1200

Oxygen level (wppm)

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

Pure V, as-melted Pure V, 1373 K

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.

26 t 1.9 t

2 d

8 d (mm)

Plates, sheets, wires, and rods

Laser weld joint

Thin pipes

W coating by plasma spraying

Creep tubes

6.6 t 4.0 t 1.0 t 0.5 t 0.25 t

f 4.57 ⫻ 0 25 t ⫻ 400 mm

f 10 ⫻ 0 5 t ⫻ 100 m m 20 mm

W coating

NIFS-HEAT-2

5 mm

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

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

NIFS-HEAT-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

surface hardening during exposure to Li Further investigation is necessary for understanding the envi-ronmental 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

4.12.7 Corrosion, Compatibility, and Hydrogen Effects

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 magnetohydro-dynamic effects (see also Chapter 4.21, Ceramic Coatings as Electrical Insulators in Fusion Blan-kets) 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 Espe-cially, the N level influences the corrosion in complex manners.28,29Figure 13shows a summary of the weight

10 -3

10 -5

10 -6

10 -7

10 -8

10 -9

10 -10

10 -11

10 -12

10 -2

s/G

Uniaxial tests

310 wppm O

Biaxial tests

699 wppm O

n = 4.3

n = 3.7

n = 0.84

n = 13

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.

0

10 -6

10 -7

10 -8

10 -9

10 -10

Creep strain (%)

In vacuum In lithium

50 MPa

70 MPa

90 Mpa

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.

1073 K

973 K

873 K

773 K

673 K

RT

Strain (%) 200

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

various temperatures.

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:Cr2:1 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 year 1

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

oxi-dizing environments is of interest both for the

perfor-mance 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

neces-sary 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 inFigure 14.34

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

The effects of oxygen level on hydrogen embrittle-ment have been investigated.Figure 15compares elon-gation versus hydrogen concentration for V–4Cr–4Ti

Ti

Cr

Ti:Cr = 2:1

50

40

30

20

10

V

I

II

–47.4 –8.2 –2.1

+7.5

+11

+0.4

+11.9 +6.7

+2.5

+1.2

–22.0 –19.0

–19.7 –2.4

–21.0 –26.4

–52.5 +5.8

Figure 13 The compositions of V–Ti–Cr alloys (wt%) with

increase (area I) and decrease (area II) of mass (g cm2) 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.

0 0.02 0.04 0.06 0.08 0.1

773

V–4Cr–4Ti

V–4Cr–4Ti–0.5Si

V–4Cr–4Ti–0.5Al V–4Cr–4Ti–0.5Y

Oxidation temperature (K)

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.

0 10

20

30 40 50

0 100 200 300 400 500 600 700

Natesan (BL-71 O:670 wppm)

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

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

Hydrogen concentration (wppm)

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.

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

radia-tion effects are considered to be loss of ductility at

lower temperature, embrittlement enhanced by

trans-mutant 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

long-term 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

inves-tigation, it is inferred that interaction of dislocations with

radiation-induced defect clusters, precipitates, or com-plexes of the two species is responsible If the precipi-tates, most likely Ti-CON, play the role in this process, reduction of impurities in the matrix can improve the properties.Figure 17compares 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.38The effects of inter-stitial impurities on the formation of dislocation loops and precipitates were investigated by ion irradiations

densities of loops and precipitates.39The 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.3The 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

Ttest= Tirr = 543 K

g = 011

100 nm

0

700 600

400 500

300 200 100 0

693 K

598 K

543 K

383 K

Normalized crosshead displacement (mm mm –1 )

Ttest~ Tirr

Load-elongation curves for V–4Cr–4Ti irradiated in HFBR to 0.5 dpa

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.