Current Status of Reduced-Activation Ferritic-Martensitic Steels R&D for Fusion Energy

7/29/2019 Current Status of Reduced-Activation Ferritic-Martensitic Steels R&D for Fusion Energy

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Current Status of Reduced-Activation Ferritic/Martensitic Steels R&D

for Fusion Energy

Akihiko Kimura

Institute of Advanced Energy, Kyoto University, Uji 611-0011, Japan

Reduced-activation ferritic/martensitic (RAF/M) steels have been considered to be the prime candidate for the fusion blanket structural

material. The irradiation data obtained up to now indicates rather high feasibility of the steels for application to fusion reactors because of their

high resistance to degradation of material performance caused by both the irradiation-induced displacement damage and transmutation helium

atoms. The martensitic structure of RAF/M steels consists of a large number of lattice defects before the irradiation, which strongly retards the

formation of displacement damage through absorption and annihilation of the point defects generated by irradiation. Transmutation helium can

be also trapped at those defects in the martensitic structure so that the growth of helium bubbles at grain boundaries is suppressed. The major

properties of the steels are well within our knowledge, and processing technologies are mostly developed for fusion application. RAF/M steels

are now certainly ready to proceed to the next stage, that is, the construction of International Thermo-nuclear Experimental Reactor Test Blanket

Modules (ITER-TBM).

Oxide dispersion strengthening (ODS) steels have been developed for higher thermal efficiency of fusion power plants. Recent irradiation

experiments indicated that the steels were quite highly resistant to neutron irradiation embrittlement, showing hardening accompanied by no loss

of ductility. High-Cr ODS steels whose chromium concentration was in the range from 14 to 19 mass% showed high resistance to corrosion in

supercritical pressurized water. It is shown that the 14Cr-ODS steel is susceptible to neither hydrogen nor helium embrittlement.

A combined utilization of ODS steels with RAF/M steels will be effective to realize fusion power early at a reasonable thermal efficiency.

(Received November 22, 2004; Accepted January 27, 2005)

Keywords: reduced activation ferritic/martensitic steel, oxide dispersion strengthening steel, irradiation embrittlement, helium effects, fusion

blanket structural material

1. Introduction

Reduced-activation ferritic/martensitic (RAF/M) steels

have been considered to be the prime candidate for structural

material of the fusion power demonstration plant (DEMO)

reactor design and beyond,19) where neutron irradiationembrittlement and transmutation helium-induced embrittle-

ment are expected to be the critical issues for reactor

operation. Among the various candidates for fusion blanket

structural material, RAF/M steels have shown rather high

resistance to the degradation of material performance caused

by neutron irradiation15,1012) and/or helium implanta-

tion.15,1316) The superiority of RAF/M steels to the other

candidates in the material performance under a fusion

environment is considered to attribute to the high trapping

capacity of defects and helium atoms in the martensitic

structure that contains a large number of trapping sites, such

as dislocations, lath boundaries, precipitates and carbides,and prevent the defects and helium atoms from aggregating

into defect and/or helium clusters which cause a degradation

of structural material.16)

Japanese major candidate RAF/M steels, F82H17) and

JLF-1,18) have been developed from the 9Cr-1Mo heat

resisting steel by replacing several alloy elements to reduce

the residual radioactivity induced by fusion neutrons. The

chemical compositions and heat treatment conditions are

listed in Table 1.19,20) JLF-1 was developed by Japanese

universities, while F82H was developed by the Japan Atomic

Energy Research Institute and Japan Steel Engineering

Holding, Inc. The microstructures of both of the steels areof tempered martensite. Chromium levels were chosen to

minimize irradiation hardening.5) The chemical compositions

were determined to keep the mechanical properties at

unirradiated condition within the similar levels of those of

commercial 9Cr-1MoNbV steel without changing heat treat-

ment specifications.

As an International Energy Association (IEA) collabora-

tive work of surveillance test of material performance of

RAF/M steels, plates, tungsten-inert-gas (TIG) welded plates

and electron-beam (EB) welded ones were prepared from two

5-ton F82H ingots, and distributed to the participating

organizations in the world. Little difficulty was found inthe fabrication of RAF/M steels in an industrial large

quantity. A mock-up of blanket structure was successfully

fabricated using a hot-isostatic pressing (HIP) technique.8)

Distribution of the material with a good traceability

enabled to compare the data obtained by different organ-

izations. The IEA heat F82H has become a reference material

Table 1 Chemical compositions and heat treatment conditions of the Japanese RAF/M Steels.

mass% C Si Mn Cr V W N Ta

F82H 0.09 0.1 0.21 7.46 0.15 1.96 0.006 0.023

JLF-1 0.1 0.2 0.45 9 0.25 2 0.05 0.07

Heat Treatment

F82H: Normalization; 1040 C for 38 min, Tempering; 750C for 1 h

JLF-1: Normalization; 1050 C for 1 h, Tempering; 780C for 1 h

Materials Transactions, Vol. 46, No. 3 (2005) pp. 394 to 404Special Issue on Fusion Blanket Structural Materials R&D in Japan#2005 The Japan Institute of Metals OVERVIEW


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in many organizations. JLF-1 has also been tested in the

Japan-US collaborative research on fusion materials for

understanding radiation effect mechanism of RAF/M steels.

Basic property data of Japanese RAF/M steels before and

after irradiation have been thus accumulated through the

efforts of many organizations in the world, and they are now

being compiled into a database.20)The influence of ferromagnetism on plasma operation has

often been worried about. Recent operation tests of the JFT-

2M facility with F82H plates in it showed no deteriorating

effects on plasma but reduction of the toroidal field ripple.21)

Since high-temperature strength of the RAF/M steels is

limited, many efforts have been made to increase the strength

for highly thermal-efficient blanket system. One of the

effective methods is to design the RAF/M steel-blanket in

combination with oxide dispersion strengthening (ODS)

ferritic/martensitic steels, which have a remarkably higher

strength than RAF/M steels. Recent irradiation experiments

clearly showed that 9Cr and 12Cr-ODS steels, which havebeen developed for the fuel cladding of fast breeder reactor

(FBRs),2225) were highly resistant to irradiation embrittle-

ment at temperatures between 573 and 773 K up to 5 and

15 dpa, respectively.26) The ODS steels never showed

irradiation-induced reduction of tensile elongation, which

was generally observed in the other metallic materials.

The ODS steels developed for FBR fuel cladding material

contain 9% or 12% chromium. It is well known that the

corrosion resistance in high-temperature water is reduced

significantly by decreasing chromium concentration below

13%. Thus, for ODS ferritic/martensitic steels, the most

critical issue for the application to water-cooling system is to

improve their corrosion resistance.27,28)

As for the hydrogen embrittlement of ferritic/martensitic

steels, it was reported that in the reduced-activation 9Cr

martensitic steels, several mass ppm of hydrogen caused a

considerable reduction in the tensile ductility accompanied

by a change in the fracture mode from ductile fracture to

intergranular or cleavage brittle fracture.29,30) However, little

attention has been paid to the hydrogen effects on ODS steels,

so far. As ODS steels have been manufactured by hot

extrusion processing, they will have an anisotropy in the

mechanical properties due to elongated grains parallel to the

extruded direction.31)

Transmutation helium sometimes causes severe embrittle-ment in materials. Helium-induced embrittlement is consid-

ered to be due to helium bubbles, namely helium-vacancy

clusters, formed at grain boundaries. In the temperature range

where a single vacancy is almost immobile, transmutation

helium is trapped by a single vacancy and helium bubbles are

not formed. At higher temperatures, however, vacancies and

helium atoms become mobile and aggregate into helium

bubbles.13,16)

In this paper, the performance of RAF/M steels under

fusion relevant environment is overviewed, and it is

demonstrated that the steels are most promising for fusion

blanket structural material towards early realization of

DEMO reactors. Current status of ODS steels R&D to

increase the thermal efficiency of power reactors is also

shown.

2. Performance Under Fusion Relevant Environment

Ferritic/martensitic (9Cr or 12Cr) steels have been

successfully used for fast-breeder-reactor (FBR) core com-

ponents to a damage level of 100 displacement per atom

(dpa) that corresponds to a wall loading of 10 MW(annual)/

m2

. This is very encouraging for the application of RAF/Msteels to the fusion core components. The response of the

Japanese RAF/M steels to irradiation has been studied using

fission reactors and ion accelerators within the framework of

domestic programs and international collaborations, which

include collaborations in the IEA RAM/S steel working

group and Japan-US bilateral collaboration programs.

2.1 Irradiation hardening

2.1.1 Neutron dose and irradiation temperature de-

pendence

The dependence of irradiation-induced hardening on

neutron dose and irradiation temperature of Japanese RAF/M steels is shown in Fig. 1, which indicates irradiation

hardening appears to saturate at a dose that depends on

irradiation temperature. The tensile tests were carried out at

room temperature. Below 700 K, irradiation induces harden-

ing that turns to softening above 700 K. At 648 K, irradiation

hardening is saturated at 10 dpa where the y 110MPa.

At above 733 K, irradiation softening occurs and appears to

saturate at 30 dpa where the y 170MPa. At below

648 K, hardening becomes large even at low neutron doses.

Especially, the hardening observed in F82H is much more

significant than a modified JLF-1 steel, which is considered

to be due to the differences in the Cr concentration (7.6% for

F82H and 9% for mod. JLF-1) and heat treatment con-

y

0

-

(523 543K,

HFIR, F82H)

dpa

IrradiationHardening,

/MPa

Tested at RTCross -head speed

= 0.5 mm/min

0 20 40 60

0

200

400

600

800

-200 733

673 703K, FFTF

648K, FFTF

573K, HFIR

Mod. JLF-1F82H

JLF-1Ni

543K, ATR

873K, FFTF

Fig. 1 The dependence of irradiation hardening of mod. JLF-1 steel

(circles) and F82H (squares) on the neutron dose (dpa).7) The irradiation

temperatures and facilities are shownin the figure. A significant irradiation

hardening was observed for F82H but not for mod. JLF-1 steel.

Current S tatus of Reduced-Activation F erritic/Martensitic Steels R&D for Fusion Energy 395


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ditions.9) It was reported that softening was due to irradiation-

induced recovery of the martensitic structure,6) and hardening

was mainly due to interstitial dislocation loops and partly due

to small precipitates of (Cr, Fe)23C6, M6C and (Ta, W)C

induced by irradiation.3234) Figure 2 shows the TEM micro-

graph of modified JLF-1 steel irradiated with neutrons up to

13 dpa at 646 K, showing dislocation loops observed from

{100} direction, indicating presence of the loops with the

Burgers vector ofah100i and 1=2ah111i. Fine (Cr, Fe)23C6carbides were also observed along lath boundaries and in the

matrix. The investigation of saturation neutron dose andstress level at a low temperature below 523 K would be

necessary to understand the detail of hardening mechanism.

Whichever irradiation induces hardening or softening, the

RAF/M steels suffer irradiation embrittlement, namely, an

increase in ductile-brittle transition temperature (DBTT) and

reduction of upper shelf energy (USE), which will be

discussed in the next session.

As for the contribution of microvoids to irradiation

hardening, positron annihilation spectrometry (PAS) studies

showed that they were ineffectual in the hardening of RAF/

M steels. Figure 3 shows the annealing behavior of irradi-

ation hardening and intensity of long lifetime component(300400 ps) of PAS in the neutron irradiated (triangles) and

helium implanted (circles) RAF/M steel, indicating that no

correlation was observed between the recovery stage of

hardening and the number density of microvoids. It is

considered that microvoids are not the main factor control-

ling irradiation hardening of RAF/M steels.35)

2.1.2 Correlation between y-DBTT

As shown in Fig. 4, there is a linear relationship between

irradiation hardening, y, and shift in DBTT, DBTT, in

neutron-irradiated ferritic/martensitic steels at various irra-

diation conditions.6) The fracture mode was not changed by

irradiation. The irradiation hardening was measured by

tensile tests at room temperature except for (c) and (e). The

shift in DBTT was normalized into that for 1/3 size Charpy

V-notch specimens following the previous work.36) This

trend suggests that irradiation embrittlement is controlled by

so called hardening-mechanism, and the similar fracture

mechanism with that of unirradiated steel plays a role in the

fracture process of the irradiated ones. Irradiation-induced

softening is also accompanied by DBTT shift. It is expected

that the fracture stress was reduced by a high-temperature

(>700K) irradiation that caused a growth of the carbides.

As for neutron dose dependence, the DBTT shift is

saturated at 10 dpa when the irradiation temperature is

Fig. 2 TEM micrograph of mod. JLF-1 steel irradiated with neutrons up to

13dpa at 646 K. The dislocation loops were observed from {100}

direction, indicating the presence of the loops with the Burgers vector of

both ah100i and 1=2ah111i.

Hv

0

20

40

60

Intensity,I2(%)

0

20

40

60

80

100

120

IncreaseinHV,

HV

As-irr. 400 500 600 700 800 900

Annealing Temperature, T/ K

PAS, Intensity, I 2

Hv

He-implanted (580appm)0.223dpa, < 423K,

JMTR, 0.15dpa,493K,

Fig. 3 Recovery behavior of the hardening (closed) and PAS intensity of

the long lifetime component (open) during isochronal annealing for30 min. in the neutron irradiated (triangles) and helium-implanted (circles)

RAF/M steel. There is no correlation between the recovery behavior of the

hardening and the number density of microvoids measured by PAS.

(1) JMTR/363K, 0.006dpa: RT

(2) MOTA/663K, 22dpa: RT




(a) JMTR/493K, 0.15dpa: RT

(b) EBRII/663K, 26dpa: RT

(c) MOTA/638K, 7dpa: 638K

(d) HFIR/323K, 5dpa: RT

(e) HFIR/673K, 40dpa: 673K

(f) ATR/543K, 2.2dpa: RT

0

50

100

150

200

250

300

-200 -100 0 100 200 300 400 500

9Cr-2W

ShiftinDBTT(K)

Irradiation Hardening,

(1)

(2)

(3)

(b)

(4)

(5)

(a)

(c)

(d)

(e)

(b)

580appm He-implanted120,

(f)

(f)

1w%Ni added9Cr-2W

: 9Cr-2W

: 9Cr-1Mo

: 9Cr-2W (He)

: 9Cr-2W-1Ni

9Cr-1Mo-(Ni)

Irr. Temp., Test Temp.

y / MPa

Fig. 4 Relation between the irradiation hardening and shift in DBTT of

neutron irradiated RAF/M steels and 9Cr-1Mo steel.6) The irradiation

hardening was measured by tensile tests at room temperature except for (c)

and (e). The shift in DBTT was normalized into that for 1/3 size Charpy

V-notch specimens following the previous work.36)

396 A. Kimura


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between 656 and 693 K, as shown in Fig. 5, which is similar

to the trend of irradiation hardening. It was found that

irradiation embrittlement was reduced very much by sub-

stituting molybdenum and nickel with tungsten and vana-

dium, respectively. Microstructural observations revealed

that the significant irradiation embrittlement was due to (Ni,

Mo) silicides induced by neutron irradiation.37)

It is of notice that the relation between y-DBTT of aRAF/M steel implanted with 580 appm of helium is almost

same with that of the steels without helium.38) This suggests

that the DBTT of helium-implanted steel is reasonably

interpreted in terms of not helium embrittlement but

displacement damage-induced embrittlement where helium

implantation simply plays a role of atomic displacement.

Helium issues will be discussed in section 2.2.

2.1.3 Varying-temperature irradiation effects

Since irradiation hardening remarkably depends on irradi-

ation temperature, varying-temperature irradiation effects

were investigated for a RAF/M steel. Two sets of temper-

ature combination was selected, namely, 498/613 K and633/793 K, because there were two important critical

temperatures for RAF/M steels: the vacancies become

mobile by dissociating from vacancy-carbon complexes at

around 520 K, and irradiation hardening turns to softening

above 700 K. Eight-cycles of 0.5 dpa of varying-temperature

irradiation was performed in the HFIR. The tensile properties

were compared with those irradiated at a constant temper-

atures of 613 K and 793 K, and the results showed no

significant effect of varying temperature irradiation on the

tensile behavior, although small enhancement of the harden-

ing was observed for the 498/613 K irradiation. This

indicates that the irradiation hardening of the RAF/M steel

is not affected by the pre-irradiation at a lower temperature

and the final irradiation temperature determines the irradi-

ation hardening of the RAF/M steel.

Effect of one cycle varying-temperature irradiation was

investigated using the JMTR where the irradiation capsule

was removed on the way of the irradiation.35) Elevation of

irradiation temperature from 493 to 693 K caused almost full

recovery of the hardening, while PAS measurement indicated

that microvoids still exists with an increase in the number

density. This again clearly indicates that microvoids are not

the main factor controlling irradiation hardening as men-

tioned in section 2.1.1.2.1.4 Fatigue properties

Fatigue properties of RAF/M steels before irradiation are

almost similar to those of non-reduced activation commercial

9Cr-1MoVNb steels.20) One of the matters of interest is the

irradiation-induced mechanism change of fatigue damage.

Post-irradiation fatigue tests were performed for F82H at

room temperature.39) Irradiation was carried out in the JMTR

to 3.8 dpa at 523 K. Neutron irradiation did not cause any

changes in fatigue life, except for the results with smallest

plastic strain range of 0.04%. The number of cycles to failure

decreased to one seventh that of the unirradiated specimen at

the strain range. Based on the observations of fracturesurface, it was suggested that the reduction might be

attributed to channel deformation under cyclic loading.

2.2 Effects of helium on irradiation embrittlement

Another main concern is transmutation helium effect on

the y and DBTT. Although helium might cause further

degradation, no difference was observed in the relationship of

y-DBTT between fast neutron irradiation and He ion

implantation, as shown in Fig. 4,6,38) indicating that the

observed embrittlement is interpreted in terms of hardening-

mechanism in both cases. Previous works on helium effects

are introduced as follows.

2.2.1 Boron-10 technique

Utilizing nuclear transmutation reaction, 10B(n; )7Li,

helium effects were often investigated for fusion structural

materials in previous works. In order to detect helium effects

0

50

100

150

200

250

300

ShiftinDBTT(K)

9Cr-1Mo

9Cr-2W

0 10 20 30 40 50dpa

Neutron Irradiatedat 656-693K (hardening)

733K (softening)

Fig. 5 The dependence of the shift in DBTT of a 9Cr-1Mo steel and a 9Cr-

2W reduced activation ferritic/martensitic (RAF/M) steel on the neutron

dose (dpa).7) The irradiation temperatures are in the range between 656and 693 K in which the irradiation caused hardening more or less. The

irradiation embrittlement is significantly reduced by the substitution of Mo

and Ni with W and V, respectively.

200

F82H

JLF-1

High Purity Ferritic Alloy (HPF)

HPF + n-B(16 appmHe)

HPF + 10-B

(80 appmHe)

150 250

DBTT (K)

JMTR

0.12dpa, 563K1.5mmCVN

Fig. 6 The shift in DBTT of each material caused by neutron irradiation.45)

Open and closed symbols are of the DBTT before and after the irradiation.

No enhancement of the embrittlement was recognized for HPF alloy by the

80 appm of helium generation in the 10-B doped steel. JLF-1 showed

better impact properties than F82H.



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on F82H, the DBTT shift induced by neutron irradiation was

compared for the F82H steels that contain either natural

boron that consists of about 20% of boron-10 and 80% of

boron-11, or boron-10 so as to vary the amount of trans-

mutation helium.40) Since the DBTT was larger in the steel

containing boron-10, they concluded that transmutation

helium enhanced irradiation embrittlement of the steel. On

the contrary, the reverse results were obtained for the

DBTT of high purity ferritic (HPF) alloys, as shown in

Fig. 6, in which open and closed symbols are of the DBTT of

each material before and after irradiation, respectively. Thediscrepancy between the above two can be explained in terms

of difference in the nitrogen concentration between the

materials they used. The nitrogen concentration in F82H and

HPF alloy is 0.015 and


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2.3 Swelling

Figure 8 shows the swelling behavior of RAF/M steels46)

as well as an austenitic stainless steel. This indicates that the

swelling of the RAF/M steels are much smaller than that of

the austenitic stainless steel. It was reported32) that, the swel-

ling becomes maximally at the irradiation temperature of

around 700 K in the RAF/M steels, reflecting that the peak

position of the average void size is located around 700 K.

Below 700 K the defect clusters, such as voids and loops

are stable. With decreasing irradiation temperature from

700 K, void size and density appear to decrease. However,

it is expected that the fine scale microvoids as well as dislo-cation loops are probably missed by TEM observation. At

temperatures above 700 K, the defect clusters no longer ther-

mally stable, which means the generated point defects anni-

hilate through the absorption by dislocation or V-I pair anni-

hilation. The absorption of point defects to the dislocations

causes the recovery of dislocation or martensite structure, re-

sulting in irradiation softening. Figure 9 summarizes irradia-

tion hardening and swelling behavior of the RAF/M steel.

2.4 Compatibility

It is well known that the susceptibility to environmentally

assisted cracking in high-temperature water is relatively low

for 912% Cr martensitic steels. On the other hand, it has

been indicated that irradiation hardening often increases the

susceptibility to IASCC of austenitic stainless steel.47) Slow

strain rate test (SSRT) in high-temperature water environ-ments was carried out for irradiated F82H specimens to 2 dpa

in JMTR at 523 K. Almost no indication of environmentally

assisted cracking was detected.

2.5 Tensile properties of joints

Post irradiation tensile tests have been performed on TIG

weld joint and HIP bonded specimens of F82H. TIG joint

specimens irradiated to 5 dpa at 573 K and 773 K exhibited

almost similar tensile properties with base metal specimens,

although hardening by irradiation for joint specimen at 573 K

was slightly smaller.48) It has been reported that HIP process

at 1313 K for 5 h with 150 MPa hydrostatic pressuresuccessfully produced joints with strength comparable to

base metal.49) According to the tensile test results of HIP joint

tensile specimens irradiated in JMTR to 1.5 dpa at 523 K,

both strength and elongation of HIPed specimens were

comparable to those of base metal specimens.50)

3. Steel Design Towards High Thermal Efficiency

Radiation-induced softening that limits the highest oper-

ation temperature is critical for RAF/M steels toward high

thermal efficiency. Many efforts to improve the high-

temperature strength have been made via: 1) alloying

control,5153) 2) ODS steel development.24,28) The formationof helium bubbles at elevated temperatures is another

concern, although the previous ion-irradiation experiments

suggest that RAF/M steels have a much better resistance to

helium bubble-induced embrittlement than the austenitic

steels.

3.1 Alloy design5153)

Based on the experimental results obtained in the Japan-

873

Experimental Results

Void Size

Swelling

Void Density

Softening

Hardening

M23

C6

(L)

MC (M)

Ni addition

y(+) () Swelling,etc Irradiation

Temperature(K)

No VoidsNo Loops

Voids & Loops

Microvoids

InvisiblePrecipitatesand/orLoops

Elemental Processes

Carbon Migrationto C-V pairs

Dissociation ofC-V complexes

Fe3C precipitation

M23

C6

(L)

M23

C6(S)

Void Swelling

Laves-P

773

673

573

473

373

Fig. 9 Summary of neutron irradiation hardening and swelling behavior of RAF/M steel.4)

0

1

2

3

4

5

0 50 100 150 200 250

Swelling(%)

dpa

12Cr FerriticsFFTF 695K

by D.S. GellesA. Kimura

SUS316-Ti773K

9Cr FerriticsFFTF 695K

Fig. 8 The dose dependence of the swelling of various steels. Swelling

resistance is much larger in ferritic steels than in austenitic steel.46) The

peak irradiation temperature for swelling is around 695 and 773 K for the

ferritic and austenitic steel, respectively.



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US collaboration, modifications of steel alloy compositions

were conceived for JLF-1: 1) an intended addition of boron,2) reduction of nitrogen concentration, 3) an increase in

tantalum concentration and 4) titanium addition. Transmu-

tation helium from natural boron, which was generally added

in the steel for the purpose for increasing hardenability, was

verified to induce no embrittlement in the RAF/M steel.

Boron can be added to increase hardenability of the steels for

production of large size of ingot. Nitrogen was reduced for

the sake of reduction of radioactivity. The role of tantalum in

precipitation hardening is characteristic. It is shown that the

size of (Ta,Ti) carbide observed by replica technique in a

RAF/M steel after FFTF/MOTA irradiation at 693 K up to

33 dpa is ranging from 10 to 200 nm. The fine carbides are

considered to be formed by the irradiation, since only largecarbides (100200 nm) were observed before the irradiation.

It is expected that the fine tantalum carbides are benefit to

improve high-temperature strength with keeping ductility at

low temperatures. Addition of titanium reduces void swel-

ling. Although the DBTT of the RAF/M steel that contained

no boron was increased by the titanium addition, a simulta-

neous addition of titanium and boron to the steel showed no

remarkable degradation of ductility by neutron irradiation.10)

3.2 ODS steels R&D

3.2.1 Improvement of corrosion resistance

Corrosion rate in a super critical pressurized water(SCPW) (783 K, 25 MPa) was measured for several kinds

of the ODS steels by means of weight loss and/or gain

measurement. Nine ODS steels were investigated: a 12Cr-

ODS steel, three 9Cr-ODS steels and five high-chromium

ODS steels of which the titanium and yttrium concentrations

were partly changed.27) The chemical compositions were

shown in Table 2. The details of the fabrication process of

the ODS steels were given in the previous paper.2225)

Figure 10 shows the dependence of mass gain on the

corrosion test period for each steel. Mass gain of the 12Cr-

ODS steel of the first generation of ODS steels is almost the

same as that of a ferritic/martensitic steel, and about one

order of magnitude larger than that of SUS316 that contains

18% of chromium. In comparison with the first generation of

12Cr-ODS steels, the second generation of high-Cr ODS

steels showed high corrosion resistance.27,28) Furthermore, an

increase of chromium and an addition of aluminum at the

same time resulted in further suppression of corrosion. It

shows that addition of chromium over 14mass% andaluminum (4.5 mass%) are very effective in suppressing

corrosion in SCPW. This is considered to be due to the

formation of Al2O3 or delaying the formation of Fe2O3.54,55)

Since increasing Cr concentration often causes aging

embrittlement due to decomposition to Cr-rich secondary

phases, impact fracture tests were also carried out for the

thermally aged high-Cr ODS steels. Figure 11 shows that the

DBTT shift occurred in 22Cr-ODS steels after aging at 773 K

for 100 h.27) Almost no change was observed for the impact

properties of 14 and 16Cr-ODS steels after the aging, while

the 22Cr-ODS steel was significantly embrittled. The DBTT

of 19Cr-4Al-ODS steel is lower than that of 19Cr-ODS steel

by about 30 K, indicating that the addition of aluminum is

also effective in improving the impact property.56,57) The

improvement of Charpy properties by the addition of Al is

considered to be due to diminishing anisotropy in the tensile

Table 2 The chemical compositions of ODS steels. (a) First generation of ODS steels: 9Cr-ODS steels for in core structural material of

FBR. (b) Second generation of ODS steels: high-Cr ODS steels for water-cooling systems.

mass% C Si Mn P S Cr W Al Ti N Y Y2O3

(a) 1st Generation

12Cr-ODS 0.02 0.05 0.046 11.97 2.02 0.30 0.010 0.19 0.24

9Cr-ODS(1) 0.13 0.05 0.044 9.00 1.95 0.20 0.013 0.29 0.37

9Cr-ODS(2) 0.13 0.006 0.01 9.11 1.95 0.18 0.017

9Cr-ODS(3) 0.13 0.005 0.01 8.84 1.91 0.29 0.010 0.35 0.44

(b) 2nd Generation

19Cr (K1) 0.05 0.041 0.06


8/11

properties, which showed almost similar ductility in both the

axial and radial direction of extruded bar. The mechanism

diminishing anisotropy by the addition of Al is not fully

understood yet. It is expected that the addition of Al causes

the increasing and decreasing the size and density of yittria

through the formation of alumina.

3.2.2 Hydrogen embrittlement

In order to evaluate susceptibility to hydrogen embrittle-

ment, tensile tests were carried out for ODS steels after and/or during hydrogen charging in an electrolyte of 1 N H2SO4with an addition of a recombination poison.58) Miniaturized

tensile specimens (gauge: l 5mm, w 1:2mm, t 0:25

mm) were sampled from the extruded rod so that the axis

direction is parallel to longitudinal (L) or transverse (T)-

direction with respect to the extruded direction. Hydrogen

charging caused a reduction of tensile elongation, which

became large with increasing charging current density or

hydrogen concentration.59)

The hydrogen concentrations in 14Cr-ODS and 9Cr-2W

steel, which were measured by TDS analysis, were shown in

Fig. 12 as a function of cathodic current density as well as thecalculation results following to McNabb & Foster trapping

model.6164) According to the previous research, the critical

hydrogen concentration required to induce cracking of

martensitic steels is reported as a several mass ppm,30) and

our analysis results of 9Cr ferritic/martensitic steels revealed

that the critical concentration, which corresponded to the

current density of 4A/m2, is also in the range of 1$2

mass ppm. On the contrary, hydrogen concentration in 14Cr-

ODS steel is much higher than that of the RAF steels at the

same charging condition, and critical hydrogen concentration

for the 14Cr-ODS was in the range of10$12mass ppm that

was almost 10 times higher concentration than the 9Cr

ferritic/martensitic steels. When hydrogen atoms are intro-

duced into the specimen by cathodic charging, they are

considered to be dissolved in solution and/or trapped at

various kinds of trapping sites such as vacancies, alloy

elements, grain boundaries, second-phase particles and

dislocations etc. However, these trapping sites have different

binding strength.22) Among them usually grain boundaries

and second-phase particles have relatively strong binding

energies in the range from 60 to 90 kJ/mol. Although the

binding energy of hydrogen to yttria is not known, it could be

large because of high affinity between yttrium and hydrogen,130 kJ/mol. Hence, it is considered that the high concen-

tration of hydrogen in ODS steel is attributed to the high

density of hydrogen trapping sites such as grain boundaries

and oxide/matrix interfaces, because the ODS steel has a

very fine size of grain, and a high number density of yittria

particles.60)

3.2.3 Helium trapping behavior

Typical examples of the thermal desorption spectrum of a

RAF/M steel (broken line) and the ODS martensitic steels-

FC (dotted line) and NT (solid line) are shown in Fig. 13 after

bombardment of 8 keV helium up to 2 1019 and 2 1020

He

/m2

.65)

Generally, there were six main peaks in theTHDS of each the steel. It is of notice that the peak IV was

only observed in the ODS steels and not in the RAF steels.

This trend became more significant with increasing the

amount of helium implantation. The peak position appeared

to be independent of amount of helium. The Peak IV was not

observed even in the 12Cr-ODS steel. Since the 12Cr-ODS

steel also contains a high density of oxide particles, the peak

IV is not directly attributed to the oxide particles. In RAF/M

steels, the martensite phase becomes unstable above 873 K,

while the martensite phase in the 9Cr-ODS steels is much

more stable than that in RAF/M steels under neutron

irradiation.66) Therefore, it is considered that the peak IV is

due to helium desorption from the -phase in which the

defects still remain and play a role in trapping helium.

Finally, in the ODS steels, the fraction of helium desorption

by bubble migration mechanism was smaller than in the

HydrogenConcentration(m

ass.ppm)

Current Density,

I/ (Am2)0.5

Hydrogen Concentration, CL/ m-3

Fig. 12 Hydrogen concentration measured by TDS after cathodic charging

to 19Cr-ODS steel (closed squares) and RAF/M steel (closed circles). 59)

Following the hydrogen trapping model,6164) the amount of trapped

hydrogen was calculated and shown as dotted lines in the figure.

0

AbsorbedEn

ergy,E/J

Temperature, T/ K

100 150 200 250 300 350 400

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

14Cr-4Al

As-received

16Cr-4Al

22Cr-4Al

773K 100h0.9

1.0

Fig. 11 Effects of thermal aging at 773K for 100 h on the impact

properties of ODS steels with different concentration of chromium.27)

Almost no effect was observed for 14Cr and 16Cr-ODS steels, while 22Cr-ODS steel showed a significant aging embrittlement. Small specimens of

1.5 mm square rod with a sharp V-notch were used for the impact test.



9/11

RAF/M steel. This suggests that the bubble formation and

growth were suppressed in the 9Cr-ODS steels.

4. Materials Development Road Map to DEMO and

Beyond

The Japanese RAF/M steels R&D road map toward

DEMO (Fig. 14) is consistent with the scheduled operation

of ITER test blanket and DEMO. Important steps includehigh-dose irradiation experiments in fission reactors, such as

the High Flux Isotope Reactor (HFIR) at the Oak Ridge

National Laboratory (ORNL), 14 MeV neutron irradiation

tests in the International Fusion Materials Irradiation Facility

(IFMIF) and application to ITER-TBM to provide an

adequate database of RAF/M steels for the design of DEMO.

4.1 Toward the ITER test blanket module

The verification of fusion blanket designing technology by

the operation of ITER-test blanket module (TBM) are

necessary to construct the DEMO blanket.6769) For the

designing and licensing of ITER-TBM, design criteria

compatible with the irradiation-induced property changes

of the materials should be prepared.

Designing and licensing of ITER-TBM using RAF/M

steels are expected to be highly promising but still required to

perform irradiation experiments with 14 MeV neutrons usingIFMIF to confirm several key issues. IFMIF will provide the

best simulation of fusion environment for testing materials

and components. Component tests in IFMIF will be followed

by the operation of the ITER-TBM, where the blanket system

performances will be verified as an integral system including

coolant, tritium breeder and neutron multiplier in real fusion

irradiation conditions.

4.2 Expected future progress in RAF/M steels toward

and beyond DEMO

Judging upon the basis of knowledge so far obtained,

RAF/M steels can be employed up to a fluence levelcorresponding to the peak DEMO wall loading (1015

MWa/m2) at temperatures below 773 K.8) DEMO will have

to demonstrate, however, that fusion is an attractive option in

many aspects, such as energy conversion efficiency and

environmental safety. Note that higher operating temperature

and less exchange frequency of the blanket should lead to

higher efficiency and reduction of nuclear waste disposal.

ODS steels27) are quite promising for high-temperature

operation of the blanket system with a high potential in

resistance to irradiation and corrosion. Application of ODS

steels to blanket structural components provides a larger

design tolerance and a higher energy conversion efficiency. A

combined utilization of ODS steels with RAF/M steels willbe effective to realize fusion power with a reasonable thermal

efficiency.8) For this purpose, development of joining

technology of these materials is necessary.

5. Summary

1) The US/Japan collaboration has been effective in accu-

mulating irradiation database and in understanding the

mechanism of irradiation effects of RAF/M steels. The

irradiation data obtained so far indicates rather high

20402020

(Irradiation)

(tests in components)

(973K/10k hr, 120MPa)

(Ferritic Steel R&D)

2005

ITER TBM

Fission reactor (Irradiation database, modeling irradiation effects)

DEMO / PROTO Design Construction Operation

IFMIF (Engineering verification, 14MeV n-effects, Helium effect)

ITER Construction Operation

RAFS Database (fast realization: Thermal Efficiency=34%)

ODS steel R&D (high performance: Thermal Efficiency.= 47%)

Combined utilization of

RAF/M steel and ODSS toward

commercial reactor

Attained an excellent creep property

Fig. 14 The RAF/M steel R&D road map toward DEMO.8)

3000

2

4

6

8

10

8keV

0

2

4

6

8

(x10 )-3 (x10 )-3

Temperature, T/ K Temperature, T/ K

2 x 1019

He+/m

22 x 10

20He

+/m2

DesorbedFractionofHeliu

m

,(DesorbedHelium)/(Total

Retention)

700 1100

9CrODS(NT)

JLF1

IIIIII

IV

V

VI

1500 300 700 1100 1500

JLF1

9CrODS(FC) 9CrODS(NT)

9CrODS(FC)

Fig. 13 THDS of 9Cr-ODS steels and JLF-1.66)

The peak IV was onlyobserved for the 9Cr-ODS steels. In the ODS steels, the fraction of helium

desorption by bubble migration mechanism was smaller than JLF-1.

402 A. Kimura


10/11

feasibility of the steels for application to fusion reactors,

because of their high resistance to degradation of material

performance by both the displacement damage and trans-

mutation helium.

2) The martensitic structure of RAF/M steels consists of a

large number of lattice defects before the irradiation, which

strongly reinforce the resistance to displacement damagethrough absorption and annihilation of the point defects

generated by the irradiation. Transmutation helium can be

trapped at those defects in the martensitic structure so that the

formation of helium clusters at grain boundaries is sup-

pressed. The martensitic structure of the RAF/M steels is

considered to be appropriate for fusion structural material.

3) The critical issue of the RAF/M steels is, therefore, to

increase the phase stability of the martensitic structure during

high-temperature irradiation. Since high-temperature per-

formance is a measure of phase stability, many efforts to

improve high-temperature strength have been made.

4) The 9 and 12Cr-ODS steels showed almost similarcorrosion behavior with the ferritic/martensitic steel in a

SCPW. However, increasing chromium (>14mass%) and

addition of aluminum (4.5 mass%) are very effective to

suppress the corrosion in SCPW. Addition of aluminum

improves the Charpy impact property of the high-Cr ODS

steels.

5) The critical hydrogen concentration required to induce

brittle fracture in ODS steels is in the range of 1012

mass ppm that is almost one order of magnitude higher value

than that of 9Cr ferritic/martensitic steels. The high capacity

of hydrogen trapping in ODS steels is considered to be due to

very fine size of grain and yittria in the steel.

6) In 9Cr-ODS steels, the fraction of helium desorption bybubble migration mechanism was smaller than that in RAF/

M steel. This suggests that the bubble formation and growth

were suppressed in the ODS steels.

7) The verification of fusion blanket designing technology by

the operation of ITER-TBM is necessary to construct the

DEMO blanket. Designing and licensing of ITER-TBM

using RAF/M steels are expected to be highly promising but

still required to perform irradiation experiments with 14 MeV

neutrons using the IFMIF to confirm several key issues.

8) Application of ODS steels to blanket structural compo-

nents provides a larger design tolerance and a higher energy

conversion efficiency. A combined utilization of ODS steelswith RAF/M steels will be effective to realize fusion power

with a reasonable thermal efficiency.

Acknowledgements

The author wishes to sincerely thank Dr. Taro Morimura,

ULVAC Inc. and Dr. Ryuta Kasada, Kyoto University,

for their vital efforts to perform neutron irradiation experi-

ments. The author would like to express his appreciation to

Prof. Katsunori Abe, Tohoku University, Prof. Akira

Kohyama, Kyoto University, and Prof. Shiro Ishino, Tokai

University, for giving opportunities to participate in the

Japan-US collaborations on Fusion Materials, such as FFTF/

MOTA and JUPITER programs. The author is very grateful

to Mr. Minoru Narui and Prof. Hideki Matsui, Tohoku

University, for their substantial support for carrying out post-

irradiation experiments at the hot laboratories. Helium

implantation study was done under the collaborative research

with Prof. Akira Hasegawa, Tohoku University. The author

thanks to Dr. Shiro Jitsukawa and his staff, the Japan Atomic

Energy Research Institute, for their collaboration on the

Japanese RAF/M steels. The author wishes to thank Dr.

Shigeharu Ukai, the Japan Nuclear Cycle DevelopmentInstitute, Dr. Masayuki Fujiwara, KOBELCO, Inc., Dr. Joo-

Suk Lee, Korean Advanced Institute of Science and

Technology, and Mr. Hang-Sik Cho, Graduate School of

Energy Science, Kyoto University, for their collaboration of

ODS steels R&D.

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Current Status of Reduced-Activation Ferritic-Martensitic Steels R&D for Fusion Energy

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