Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels Comprehensive nuclear materials 4 08 oxide dispersion strengthened steels
Trang 1S Ukai
Hokkaido University, Sapporo, Japan
ß 2012 Elsevier Ltd All rights reserved.
Abbreviations
CTT Continuous cooling transformation
CEN-SCK Centre d’Etude de l’e´nergie Nucleaire –
Studiecentrum voor Kernenergie
CVN Charpy V-notch EFTEM Energy-filtered transmission electron
microscopy EPMA Electron probe microanalysis
241
Trang 2FFT Fast Fourier transformation
HIP Hot isostatic pressing
HRTEM High-resolution transmission electron
microscopy
INCO International Nickel Company
JAEA Japan Atomic Energy Agency
LBE Lead–bismuth eutectic
LFR Lead fast reactor
LMP Larson–Miller parameter
MA Mechanical alloying
MOX Mixed oxide
ODS Oxide dispersion strengthened
PMW Pulse magnetic Welding
PRW Pressurized resistance welding
SCPW Super critical pressurized water
SEM Secondary electron microscopy
SFR Sodium fast reactor
TIG Tungsten inert gas welding
UTS Ultimate tensile strength
Recent progress in oxide dispersion strengthened
(ODS) steels produced by mechanical alloying (MA)
techniques allows them to be used as fuel cladding in
sodium-cooled fast reactors (SFR) The thermally
stable oxide particles dispersed in the ferritic matrix
improve the radiation resistance and creep resistance
at high temperature As a result, ODS steels have a
strong potential for high burnup (long-life) and
high-temperature applications typical for SFR fuels The
attractiveness of ODS steels is due not only to the
nanosize oxide particles composed of Y–Ti–O atoms
but also to their controlled micron-size grain
mor-phology We review existing knowledge on the
crys-talline structure and lattice coherency of these
nanosize particles with their surrounding matrix,
since these factors dominate the dispersion and
strength-determining mechanism through dislocation
interaction The development of manufacturing
pro-cesses is a principal issue for hardened ODS steels to
realize long, thin-walled ODS steel cladding on
pro-duction scales There was the long-standing problem in
low hoop strength due to the extremely elongated fine
grains parallel to the rolling direction To soften
hard-ened cold-rolled products and modify their grain
morphology, martensitic 9Cr-ODS steels and ferritic
12Cr-ODS steels have been developed Current
prog-ress in the development of these ODS steel claddings,
including their relevant mechanical properties, for
example, tensile and creep rupture strengths in thehoop directions, and irradiation performance, isreviewed The development of Al-added high Cr-ODS steel cladding is also addressed, with a focus onsuperior resistance to oxidation and corrosion in a lead–bismuth eutectic (LBE), and supercritical pressurizedwater (SCPW) in the international Generation IVadvanced nuclear power system Nanocluster ODS
structure materials, are not addressed in this chapter
Control
essential to improving the high temperature strength
of ODS steels, is attained by the dissociation of oxide
into the ferritic steel matrix during the MA process.Subsequent annealing induces oxide particles to pre-cipitate finely at elevated temperature of around
the precipitation of Y–Ti complex oxide particles
emis-sion ion micro-probe (FIM) analysis confirmed thatthis type of complex oxide is constituted of several
was investigated by means of a small angle neutron
)plots for the milled U14YWT(Fe–14Cr–0.4Ti–3W–
that the hot isostatic pressing (HIP) of U14YWT at
density of nanoclusters, as designated by Odette
Figure 1(b)shows the effects of HIP (filled symbols)and powder annealing (open symbols) at tempera-
in magnitude and decrease in slope of the dS/dO
nanoclus-ters decreases and their number density increaseswith decreasing temperature at HIP and powder
scattering and lowest sloping, which indicates that thesmallest-sized nanoclusters precipitate with the high-est number density at lower temperatures In terms of
an X-ray diffraction experiment using Super Photon
Trang 3ring-8 eV (Spring-8) constructed in Japan, Kim et al.
recently reported that nanoclusters could be in a
noncrystalline state and can be transformed to
With regard to ODS steels without Ti, high
resolu-tion (HR) TEM investigaresolu-tions were performed by
The crystallographic lattice of the metal matrix
Figure 2 shows an HRTEM image taken from an
lattice This image was taken from the grain, oriented
Fourier transformation (FFT) of the image shows the
matrix lattice as a hexagonal pattern with diffraction
is rectangular, with diffraction spots of the {2 2 2}
type and a corresponding atomic planes distance
of dYO(2 2 2)¼ 0.306 nm The angle of 70.5 between
Figure 2(a)confirms that the Y2O3particle is oriented
and the ferritic matrix can be estimated by Klimiankou
as follows:
3dM ð1 1 0Þ 2dYO ð2 2 2Þ
This result suggests that a coherency could be satisfied
Concerning the Y–Ti–O complex oxide particles
an energy-filtered (EF) TEM micrograph from aY–Ti–O particle in which two atomic planes are simul-
phases with the [1 1 0] zone axis In fact, the measured
10
1
850 ⬚C HIP (No Y 2 O3)
850 ⬚C HIP (U14YWT) As-MA (U14YWT) As-MA (No Y2O3) 0.1
0.01
2 0
HIPed materials Powder anneals
(2 2− ) (1 1−0) 3.06 Å
(0 1 1−)
(1 0 1−) (−1 1 0)
(−2 2 2)
(−2 2 2−)
(−1 0 1) 70.5 ⬚
Figure 2 HRTEM micrograph of the Y 2 O 3 particle with surrounding matrix (a) and FFT image of micrograph (b) The diffraction spots from Y 2 O 3 particle of {2 2 2} type form the rectangle, whereas diffraction spots from the matrix of {1 1 0} type form the hexagon at the [1 1 0] zone axis and [1 1 1] of matrix Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A J Nucl Mater 2004, 329–333, 347–351.
Trang 4data are equal to the following data calculated from the
and an angle between the (0 0 4) and (2 2 2) atomic
composition
These findings suggest that nano-oxide particles
precipitate from the ferritic matrix, maintaining
crys-talline coherency or partial-coherency with a ferritic
matrix In general, the nucleation and growth of
pre-cipitates proceeds, as both interfacial and strain
ener-gies become minimal In the case of ODS steels,
interfacial coherency could be maintained between
thermodynamically stable nanoparticle precipitates
and the ferritic matrix in order to decrease the free
energy in the system from the extremely high energy
state induced by MA Elucidation of the details of the
nanoscale precipitation is important not only as basic
materials science research but also as the
develop-ment of high-strength engineering materials
Microstructure
9Cr-ODS steels are being developed by the JAEA
( Japan Atomic Energy Agency) for application to
SFR fuel cladding Their standard chemical
The chromium concentration was determined to be
9 wt% in terms of ductility, fracture toughness, and
corrosion resistance based on a series of irradiation
data of ferrite steels The addition of titanium
pro-duces the nanoscale dispersion of oxide particles,
which leads to a markedly improved high-temperature
strength If titanium is added to excess, however, it
creates too much strength, which negatively impacts
cold rolling and cold workability To achieve a ance between strength and workability, a value of0.2 wt% was selected Tungsten of 2 wt% is alsoadded in order to improve high-temperature strength
bal-by means of solid solution hardening
easily controlled by a reversible a–g transformationwith a remarkably high driving force of a few hundred
steels By inducing reversible a–g transformations,9Cr-ODS steel cladding for fast-reactor fuel elements
is currently being manufactured at the JAEA
The microstructure of 9Cr-ODS steel cladding isbasically tempered martensite However, it has beenrecognized that 9Cr-ODS steel cladding manufactured
in an engineering process possesses a dual-phase ture that comprises both tempered martensite andferrite phases An example of their microstructure is
and the elongated phase is indicated by arrows Their
The formation of a ferrite phase in 9Cr-ODS steel issomewhat unusual, because only the full martensitephase can be expected in 9Cr-ferritic steel withoutyttria under normalizing and air-cooling conditions.Moreover, the high-temperature strength of manu-factured 9Cr-ODS steel is significantly improved by
Figure 3 EFTEM images of Y 2 Ti 2 O 7 particles.
Reproduced from Kliniankou, M.; Lindau, R.; Moslang, A.
J Nucl Mater 2004, 329–333, 347–351.
Residual ferrite
Tempered martensite
20 μm
Figure 4 Microstructure of 9Cr-ODS steel showing residual ferrite and tempered martensite.
Trang 5Therefore, the control of ferrite phase formation is a
key to the realization of high-temperature strength in
9Cr-ODS steel cladding
Strength Characterization
characterization
The computed phase diagram of the Fe–0.13C–2W–
respect to carbon content For a carbon content
of 0.13 wt%, a single austenite g-phase containing
TiC carbide exists at a normalizing temperature of
at this temperature corresponds to a carbon content
of 0.08 wt%, beyond which d-ferrite is not stable The
full martensite structure, whereas the specimens with
compris-ing both martensite and ferrite phases Digital image
analyses show that the area fraction of the ferrite phase
High-temperature X-ray diffraction measurement at
only to the austenite g-phase, whereas specimens
corresponding not only to an austenite g-phase but to a
ferrite phase as well The austenite g-phase transforms
to the martensite phase, but the ferrite phase remainsunchanged by quenching Considering that the ferritephase is formed only in the specimens containing 0.35
have an identical chemical composition except for
a–g reverse transformation
Figure 722
shows the results of dilatometric surement when 9Cr–0.13C–2W–0.2Ti is heated
a–g-phase, which corresponds reasonably well with thecomputed phase diagram The addition of 0.35 wt%
degree of reduction in linear thermal expansion duringthe reverse transformation of the a–g-phase; this obser-vation indicates that the entire a-phase could not betransformed to a g-phase This untransformed ferritephase was designated as a residual ferrite
particlesAlinger’s results indicate that the mechanically
radius and highest density in Y–Ti complex oxide
Solid: with residual ferrite Open: without residual ferrite
Mechanically milled without Y2O3
700 ⬚C
Figure 5 Uni-axial creep rupture strength of 9Cr-ODS
steels at 700C after the normalizing-and-tempering
(1050C 1 h, Ar-gas cooling (AC) = > 780 C 1 h, AC)
with and without residual ferrite Reproduced from
Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T.
Mater Trans 2005, 46, 487.
0
1500 1400 1300 1200 1100 1000 900 800 700 600 500 400 300
Trang 6Y2O3particles are decomposed during MA, subsequent
annealing results in the formation and precipitation of
Y–Ti complex oxide particles at elevated temperatures
which is higher than the precipitation temperature of
Y–Ti complex oxide particles, it is possible that the
retention of the residual a-ferrite can be attributed
to the presence of Y–Ti complex oxide particles in
9Cr-ODS steels These particles could block the
motion of the a–g interface, thereby partly
suppres-sing the reverse transformation from a- to g-phase
This section presents a quantitative evaluation of
this process
The chemical driving force (DG) for the reverse
transformation from a- to g-phase in the Fe–0.13C–
terms of Gibbs energy versus carbon content curves at
each temperature These curves were derived using
the Thermo-Calc code and the TCFE6 database The
The peak value of the driving force for the reverse
The pinning force (F ) against the motion of
the a–g interface can be expressed as the following
equation, which was derived from the modified Zener
p
a- and g-phases, and its value was selected to be
the oxide particles (m) in the a-phase; its value wasdetermined as 1.5 nm by using TEM observation The
oxide particles (), and was derived on the basis of theexperimental evidence that oxide particles consist of
afore-mentioned equation, the value of pinning force F was
The velocity of the a–g interface motion (v) isproportional to the difference between F and DG, asshown in the following equation:
M is the mobility of the interface DG and F arecompetitive, and DG > F indicates a positive velocityfor the interface motion, that is, the reverse trans-formation from a- to g-phase On the other hand,
DG < F indicates that the a–g interface can be
Figure 7 Results of linear thermal expansion
measurement between 700 and 1100C at temperature
rising of 0.33C s1for 0 mass % and 0.35 mass % Y 2 O 3 in
9Cr–0.13C–2W–0.2Ti specimens Reproduced from
Yamamoto, M.; Ukai, S.; Hayashi, S.; Kaito, T.; Ohtsuka, S.
Mater Sci Eng A 2010, 527, 4418–4423.
0 2 4 6 8 10 12
Trang 7pinned by oxide particles so that the a-phase is,
thus, retained The results of the calculation shown
inFigure 822
con-tents of 0.35 and 0.7 wt%, the pinning force is larger
than the driving force for 0.13 wt% C These results
are reasonably consistent with our observation of
the retainment of residual ferrite during a–g reverse
transformation
On the basis of the aforementioned discussion, the
formation process of the residual ferrite in Fe–0.13C–
Figure 9 At the AC1 point, the carbide begins to
decompose, and a–g inverse transformation takes
place in the area of higher carbon content around
the decomposed carbide, where the driving force of
the reverse transformation exceeds the pinning force
because the carbon content may be >0.2 wt% (see
Figure 8) The g-phase could be enlarged by these
carbon content achieves equilibrium at 0.13 wt%,
driving force (0.13C), and the velocity of the a–g
interface motion is markedly reduced due to dragging
by the oxide particles Thus, the a-ferrite could be
Nanoindentation measurements were conducted in
order to evaluate the mechanical properties of the
residual ferrite itself The trace of a Berkovich tip can
be placed within the interiors of the residual ferrite
regions, while conventional micro-Vickers diamond
shows the hardness change in the individual phases
measured by this nanoindentation technique as a
in hardness is significantly restricted in the residualferrite as compared to that of the martensite phase
in terms of increasing the tempering conditions.The overall hardness measured by the micro-Vickerstester is also shown by the broken line which coversboth the residual ferrite and martensite, therefore,representing the average hardness of both phases
(Larson–Miller parameter), hardness can be converted
to yield stress at room temperature for the individualphases: 1360 MPa for the residual ferrite and 930 MPafor the tempered martensite The yield strength of theresidual ferrite is 1.5 times higher than that of mar-
A full ferrite ODS steel and full martensite ODSsteel were manufactured, and the oxide particle dis-tribution in both ODS steels was measured by TEM
that a few nanometer-sized oxide particles are finelydistributed in the full ferrite ODS steel, whereas theirsize is coarsened in the bi-modal distribution in themartensite ODS steel Considering that the residualferrite phase belongs to full ferrite ODS steel, resid-ual ferrite contains fine (nanosized) oxide particleswhich are responsible for higher strength in residualferrite containing ODS steels In regard to the bi-modal distribution of oxide particles in martensite
2.0 3.0 4.0 5.0 6.0
Average covering residual ferrite and tempered martensite
Tempered martensite Residual ferrite
Figure 10 Hardness change at room temperature as
a function of tempering conditions for the residual ferrite and tempered martensite NT: normalizing and tempering; FC: furnace cooling Ukai, S.; Ohtsuka, S.; Kaito, T.; Sakasegawa, H.; Chikata, N.; Hayashi, S.; Ohnuki,
S Mater Sci Eng A 2009, 510–511, 115–120.
Figure 9 Formation process of residual ferrite in
9Cr-ODS steel (Fe–0.13C–2W–0.2Ti–0.35Y 2 O 3 ).
Reproduced from Yamamoto, M.; Ukai, S.; Hayashi, S.;
Kaito, T.; Ohtsuka, S Mater Sci Eng A 2010, 527,
4418–4423.
Trang 8ODS steels, the a–g-phase transformation could
induce the coarsening of oxide particles by disturbing
the interface coherency between these particles and
the g-phase matrix
diagram
The preparation of a CCT (continuous cooling
trans-formation) diagram is essential to the microstructure
diagram that was experimentally constructed for
matrix phase in order to fully transform to martensite
is extremely higher in 9Cr-ODS steel (solid circularsymbol) than in mechanically milled EM10 (opendiamond symbol) that does not contain added
the process of continuous cooling transformation.The minimum cooling rate is known to increasewith a decrease in the size of prior austenite (g) grains.This smaller size of prior g grains provides morenucleation sites (grain boundaries) for a g–a-phasetransformation, so that a higher cooling rate isrequired to enable steel with small prior g grains tofully transform to a The presence of residual ferriterestricts the growth of g grains; the prior grain size ofresidual ferrite-containing steel is roughly 5 mm, thusincreasing the minimum cooling rate to produce a fullmartensite matrix
In steel that does not contain residual ferrite andthe mechanically milled EM10, the size of the prior
g grains is roughly 10 mm and 35 mm, respectively
by the relationship between the size of prior g grains
normal-izing heat treatment used in commercial furnaces, the
100 200 300 400 500 600 700 800 900 1000
300 500 700 900 1100 1300
30 (K/h) Austenite ( γ) => Ferrite (α)
Figure 12 CCT diagram of 9Cr-ODS steel Reproduced from Ohtsuka, S.; Ukai, S.; Fujiwara, M.; Kaito, T.; Narita, T J Nucl Mater 2006, 351, 241.
Figure 11 TEM photograph of the oxide particles:
(a) finely distributed oxide particles in full ferrite ODS steel
and (b) bi-modal distribution of oxide particles with larger
size in the full martensite ODS steel Yamamoto, M.; Ukai, S.;
Hayashi, S.; Kaito, T.; Ohtsuka, S J Nucl Mater 2011,
417, 237–240.
Trang 9the matrix phase of 9Cr-ODS steel cladding consists
of residual ferrite, martensite, and a small amount of
transformed ferrite from the g-phase
9Cr-ODS steels are promising materials to enable
fast reactor fuel cladding to realize a high burnup
radiation resistance and high temperature strength
Figure 13shows a series of manufacturing processes
of fuel cladding that is 8.5 mm in diameter by 0.5 mm
in thickness by 2 m in length The element powders
and yttria powder are mechanically alloyed for 48 h in
an argon gas atmosphere using an attrition type ball
mill with a capacity of 10 kg batch The mechanically
alloyed powders are sealed in hollow-shaped cans
The hollow shape of the bars is consolidated by
the dimensions of 32 mm in outer diameter, 5.5 mm in
wall-thickness, and 4 m in length After machining
to the precise dimensions, claddings are produced at
their final dimension (8.5 mm in outer diameter,
0.5 mm in thickness, and 2 m in length) by four-pass
rolling with about a 50% reduction ratio on each pass
by using a pilger mill
Without heat treatment, it is too difficult to
man-ufacture cladding for ODS steels by the cold-rolling
process Using the CCT diagram of 9Cr–0.13C–2W–
intermedi-ate heat treatment in order to induce the ferrite phase
at room temperature without martensite tion This phase has a lower degree of hardness Hard-ened cladding due to the accumulation of colddeformation can be sufficiently softened by this inter-mediate heat treatment, and cold rolling can then be
represents the typical hardness change of 9Cr-ODSsteel in the process of cladding manufacturing byrepeated cold rolling and intermediate heat treatment.The elongated grain structure induced by the fourthcold rolling can ultimately be made into equi-axedgrain structure by the final heat treatment, which
Mechanical alloying (MA)
Figure 13 Cladding tube manufacturing process developed for 9Cr-ODS steel.
350 400 450
Figure 14 Hardness change in the process of cold rolling and intermediate and final heat treatments for cladding tube manufacturing of 9Cr-ODS steels.
Trang 10consists of normalizing at 1050C for 1 h, followed by
The lifetime of a fast reactor fuel pin is most strongly
determined by the internal creep rupture strength
of the cladding induced by the internal pressure of
9Cr-ODS steel cladding, internal creep rupture data
Additionally, the best fit lines for hoop stress versus
rupture time at each temperature are shown by solid
lines These results confirmed that creep rupture
strengths in the hoop and longitudinal directions of
cladding are almost the same, due to their equi-axed
grains The corresponding creep rupture curves for
for comparison PNC316 is a typical austenitic
clad-ding developed by JAEA in the fast reactor program
Notably, superior performance in rupture time is
shown in 9Cr-ODS steel cladding The slope of
PNC316 is steeper, and there is a cross-over before
stress condition of the fast reactor fuel pin gradually
increases due to the accumulation of fission gases and
reaches around 120 MPa at its final service milestone
that 9Cr-ODS steel cladding is of advantage
The ultimate tensile strength (UTS) of 9Cr-ODSferritic cladding in the hoop direction as measured
in a temperature range from room temperature to
corresponding data for the ferritic–martensitic
fast reactor fuel cladding The strength of 9Cr-ODSsteel is superior to that of conventional PNC-FMS.The uniform elongation that takes place from room
fast reactor is commonly operated, the measureduniform elongation exhibits adequate ductility Thisadvantage of superior elongation in the produced clad-dings can probably be ascribed to the pinning of dis-locations by oxide particles, which retard recovery andsustain work-hardening
When JAEA started to develop ODS steels in
phase and does not include the martensite Based onthe results of R&D conducted for several years, threekinds of claddings, 63DSA, 1DK, and 1DS, weremanufactured in 1990 Their chemical compositions
Time to rupture (h)
Stress range for SFR fuel cladding
9Cr-ODS (923 K) 9Cr-ODS (973 K) 9Cr-ODS (1023 K)
500 400 300
200
100 90 80 70 60 50
PNC316 (1023 K)
Figure 15 Creep rupture curves of 9Cr-ODS steel claddings in hoop direction by using internally pressurized specimens
at temperatures of 650, 700, and 750C, compared with those of HT9 and PNC316 Reproduced from Allen, T.; Burlet, H.; Nanstad, R K.; Samaras, M.; Ukai, S Mater Res Soc Bull 2009, 34(1), 20–27.
Trang 11are 13Cr–0.02C–3W–0.7Ti–0.46Y2O3(63DSA), 13Cr–
except for the rolling process and intermediate heat
treatment, because cold-rolling processing can be hardly
applied to these ODS steels In the case of the 1DK
thin-walled cladding in the dimension of 7.5 mm outer
diameter, 0.4 mm thickness, and 1 m length In the case
of the 63DSA and IDS claddings, only six warm rolling
and 63DSA and 1DS claddings at 1100 for 3.6 ks.The uni- and bi-axial creep rupture strengths of
in Figure 17, where the uni-axial corresponds tothe hot working direction and bi-axial belongs
Mol-ODS (DT2203Y05, see figure 28)
1000
600 500 400
Trang 12there is strong strength anisotropy, and the bi-axial
creep rupture strength is considerably lower than
that of the uni-axial direction Microstructure
obser-vations of these claddings exhibited the elongated
grains like a bamboo structure in parallel to the
working direction The strength degradation in the
bi-axial/internal hoop direction, which is essential
for the fuel elements, should be mainly attributed to
the grain boundary sliding and crack propagation due
to stress concentration
Based on the aforementioned finding in ODS steels,
the recrystallization processing was extensively
stud-ied to change the substantially elongated grain
recrys-tallized structure in the ODS ferritic steels The two
types of 12Cr-ODS steels in the chemical
From these heat-treated bars, the internal creep
Figure 1835
exhibits a comparison of the creep
rupture strength of recrystallized (A3) and
von Mises’ equivalent stress was estimated for theinternal hoop stress The unrecrystallized specimenshows significant strength anisotropy in uni-axial andinternal creep rupture strength, whereas the recrystal-lized specimens reveal decrease of anisotropy, whereuni-axial creep rupture strength decreases and inter-nal strength approaches the uni-axial strength Theseresults demonstrate that the recrystallization processadequately improves the creep rupture strength inthe internal hoop direction Furthermore, softening
by recrystallization makes it possible to manufacturecladding by cold-rolling processing
ManufacturingBased on the aforementioned finding, claddingmanufacturing tests were conducted using cold-rolling pilger mill in 12Cr-ODS ferritic steels with
recrystallized structure, and their internal creep
speci-mens are denoted as F1 to F4 The four levels of
mass %, and the titanium content ranged from 0.13
to 0.31 mass % Cold rolling by a pilger mill was
A3-2(1123 K extrusion) A3-1(1423 K extrusion) A15
Solid: uni-axial Open: internal
100
200 300 400 500 600
Trang 13repeated twice with a reduction ratio of about 50%
per rolling The intermediate heat-treatment to
soften the cold-rolled cladding was performed at
Figure 19 shows the optical microstructures of
the manufactured claddings in the longitudinal and
to be recrystallized However, the extent of
recrys-tallization depends on the yttria and titanium
con-tents In the transverse cross-section, the grain size
becomes slightly finer with increasing yttria and
titanium contents from F1 to F4 In the F4
speci-men, the elongated grains can be still seen and the
aspect ratio in the longitudinal (L) and transverse
(T) directions is large, whereas the aspect ratio
of specimen F1 appears to be nearly unity These
findings show that F4 specimen with higher yttria
and titanium contents did not attain the perfectly
recrystallized grain structure by the annealing of
The internal creep rupture properties of the
Figure 20.36Increasing the yttria and titanium tents improves the internal creep rupture strength(F4 > F3 > F2 > F1) The uni-axial creep rupturestrength for F4 is also plotted; there is the strengthanisotropy between the uni-axial and internal hoopdirections This strength anisotropy can be associatedwith the slightly elongated grain structure shown in
con-Figure 19.The stress–strain rate relationship was investi-gated for ODS ferritic claddings to evaluate thecreep deformation mode The results of the analyses
general, the creep strain rate in the steady-state
where n is the stress exponent and A is the
Table 1 Chemical composition of F1–F4 specimens with different titanium and yttria contents (mass %) in 12Cr-ODS steels
Specimen no Chemical composition (mass %)
Trang 14the uni-axial creep mode, a significantly high stress
sensitivity of n ¼ 43.7 appears This stress exponent
that initiates the strain is clearly located around
250 MPa; this stress corresponds to the so-called
threshold stress for deformation On the other hand,
the stress exponent, n, is 10.4 for the internal creep
mode of F4, and a higher strain rate is found evenbelow a stress of 200 MPa A transverse section of thisspecimen shows finely equi-axed grains of 5–10 mm(Figure 19) Apart from pinning the gliding dislocationsdue to oxide particle-dislocation interaction, the defor-mation mechanism associated with grain morphologymay be the dominant factor that induced acceleratedstrain in the hoop stress mode of the tubular specimen
In order to characterize the high temperaturestrength of manufactured 12Cr-ODS steel cladding,its strength mechanism was evaluated from the view-
and dislocations This interaction could be lated by the void-hardening mechanism proposed
re-placed by voids The oxide particle-hardening stress
into account the interaction between the branches of
for screw dislocation;
B ¼ 0:6for edge dislocation;
B ¼ 0:7where G is the Shear modulus, v is Poisson’s ratio, M
is the Taylor factor, b is the magnitude of Burgers
which the dislocation detaches from the particles
particles on a slip plane and is given as a function of
where the averages are calculated by considering thesize distribution of the particles The factor 1.25 isthe conversion coefficient from regular square distri-
particles by means of TEM D is the harmonic mean
the oxide particle-hardening stress was estimated by
Figure 21 Stress–strain rate relationship for internal creep
of specimens F1–F4 and PNC-FMS, and for uni-axial creep
of specimen F4 at 700C Reproduced from Ukai, S.;
Okuda, T.; Fujiwara, M.; Kobayashi, T.; Mizuta, S.;
Nakashima, H J Nucl Sci Technol 2002, 39(8), 872–879.
Uni-axial
Internal, bi-axial direction
F1 F3 F4 (uni-axial, this work)
Figure 20 The creep rupture strength in hoop direction
for pressurized F1 to F4 specimens at 700C Reproduced
from Ukai, S.; Okuda, T.; Fujiwara, M.; Kobayashi, T.;
Mizuta, S.; Nakashima, H J Nucl Sci Technol 2002,
39(8), 872–879.
Trang 15substituting l, M ¼ 3.0,41 n ¼ 0.334, b ¼ 2.48
Figure 22 shows the results of analyses in
The oxide particle-hardening stress levels estimated
are represented by vertical bars, with the upper and
lower bars derived from an estimate of edge and screw
uni-axial mode of the F4 specimen is shown by anopen circle These results imply that the higher oxideparticle-hardening stress for specimen F4 is due to its
The lower band represents the stress corresponding
direc-tional mode For the F1 specimen, as a stress level
the oxide particle-hardening stress, the strong ropy tends to disappear However, for the F3 andF4 specimens with a shortened distance between par-
hoop direction are degraded from the oxide hardening stress The strong anisotropy still remains inthe F4 specimen The accelerated deformation in theinternal hoop direction could be the result of grainboundary sliding, since finely equi-axed grains with asmall size of 5–10 mm are formed, and the grain bound-aries occupy a large fractional area in the transverse
Based on these results, it seems to be difficult to controlinternal creep rupture strength by recrystallizationprocessing in 12Cr-ODS steel cladding
Generation IV advanced nuclear power systems areproposed; the temperature and dose regimes for their
supercritical water-cooled reactor (SCWR) and thelead fast reactor (LFR) require a higher neutron dose
Face-to-face distance between particles, l (nm)
Stress at strain rate of
10−9 S−1 in the internal
hoop direction
F4
F3 F2
F1
F4 (uni-axially longitudinal) Oxide particle-hardening stress (sp ) from particle distribution by TEM
Figure 22 Comparison of oxide particle-hardening stress
estimated from dispersion parameters of F1, F2, F3, and
F4 specimens, uni-axially longitudinal creep strength of
F4 specimen, and internal creep strength in hoop direction
at a strain rate of 109s1for F1, F3, and F4 specimens,
as functions of face-to-face distance between particles.
Each stress was obtained at 973 K Note that internal creep
strength is located below the oxide particle-hardening
stress due to the grain boundary sliding in the hoop stress
mode Reproduced from Ukai, S.; Okuda, T.; Fujiwara, M.;
Kobayashi, T.; Mizuta, S.; Nakashima, H J Nucl Sci.
Technol 2002, 39(8), 872–879.
0 200 400 600 800 1000 1200 1400
MSR GFR