3FTFBSDI

Origin of Embrittlement of Commercial Nuclear Reactor Pressure
Vessel Steels: Cu-rich Precipitates and Defects
Mechanisms of embrittlement of nuclear reactor pressure
vessel steels were studied by using positron annihilation
spectroscopy combined with 3D atom probe technique. In
surveillance test specimens of a commercial western-type
pressurized water reactor, Cu precipitation takes place atthe very begining of reactor lifetime, which is followed byformation of irradiation defects.
The irradiation-induced embrittlement of the nuclear
reactor pressure vessel (RPV) steels is currently a vital issue
for ensuring safe operation of nuclear power plants in theworld since the reactors of the first generation are going to
exceed their initially designed operation lifetimes. Thus, it isof great importance to reveal the nanostructural change,
attributed to the origins of the embrittlement, during long-term
operation of the commercial nuclear reactors [1]. Here wedemonstrate the formation of Cu-rich nano-precipitates(CRNPs) at the very begining of the reactor lifetime by using
3D atom probe (3DAP), which is followed by retarded
formation of irradiation defects (matrix defects) by using
positron annihilation, in surveillance test specimens of apressurized water reactor (PWR), Doel-2 in Belgium.
Figure 1 shows 3DAP atom maps of the solute distributions
in Doel-2 RPV surveillance test specimens of low dose ((a)
0.83×1019 n/cm2) and high dose ((b) 5.1×1019 n/cm2),
corresponding to the in-service irradiation periods for about 3
and 20 years, respectively. Local-electrode type 3DAP
apparatus was employed, which enabled us to analyze muchlarger volume than conventional type 3DAP and to reduce
the fracture of the irradiated needle specimens. In the both
specimens, CRNPs were clearly observed, which shows that
3 years of irradiation is enough to cause Cu precipitation in
the reactor, and the coarsening of CRNPs takes place withincreasing the dose. In the CRNPs, Mn and P atoms areenriched even at the low dose, while Ni atoms are enriched
only at the high dose.
The residual Cu content in the matrix are also obtained
from the 3DAP mappings directly. We found that the residual
Cu content for the high dose is similar (0.1% in weight) to thatfor the low dose, showing that the fraction of Cu atoms
consisting of CRNPs has saturated at the begining of theirradiation.
Figure 2 shows high (H)/ low (L) momentum component
fractions of coincidence Doppler broadening (CDB) spectraof positron annihilation for the Doel-2 surveillance test
specimens, normalized to those for pure Fe (unirradiated).
The (L,H) points for the well-annealed pure Cu, and neutronirradiated pure Fe are also plotted as references. Higher L
and H correspond to positron trapping at vacancy-type
defects and at CRNPs due to positron quantum-dot-like state[2], respectively. The results shows that the vacancy-type
defects have accumulated with dose so that the fraction of
positrons trapped at the vacancy-type defects increases,
accompanying decrease in the annihilation fraction in CRNPs.
By positron lifetime method [3], the vacancy-type defects are
identified mainly as small vacancy clusters (V3~V4).

Advanced Nuclear Energy Materials

Fig. 1. 3DAP atom maps of the solute distributions in RPV
surveillance test specimens of Doel-2. (a) 0.83×1019 n/cm2,(b) 5.1×1019 n/cm2.

Fig. 2. High/Low momentum component fractions of CDB
spectra for Doel-2 surveillance test specimens.
References[1] Y. Nagai et al., Appl. Phys. Lett. 87, 261920 (2005).
[2] Y. Nagai et al., Phys. Rev. B. 63, 134110 (2001).
[3] M. Jardin et al., Nucl. Instr. Meth. In Phys. Res. A 568,
716 (2006).
Contact toYasuyoshi Nagai (Irradiation Effects in Nuclear and Their
Related Materials Division)e-mail: nagai@imr.tohoku.ac.jp

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3FTFBSDI

Electrical Conductivity of Proton Conductive Ceramics
under Reactor Irradiation
Electrical charges may be transported in ceramics by
not only electrons but also by electron-holes, ions, andprotons. Especially in nuclear fusion environments, electricalconductivity by proton migration (protonic conduction) willplay an important role, as supply of hydrogen isotopes issufficient and working temperature for ceramics will be ingeneral high.
Perovskite-type oxides, yttrium doped barium-cerium
oxide, (BaCe0.9Y0.1O3-δ), were irradiated in the Japan
Materials Testing Reactor (JMTR) in the Oarai Research
Establishment of Japan Atomic Energy Agency (JAEA) [1].
The electrical conductivity was measured in-situ under thereactor irradiation, as illustrated in Fig. 1. The fast (E>1.0
MeV) and the thermal (E<0.683 eV) neutron fluxes were in
the range of 1x1016-6x1017 and 1x1017-1.6x1018 n/m2s,
respectively. The associated gamma-ray dose rates were in
the range 0.1-2.0 kGy/s. Subcapsules, each of which wascontaining one specimen, were accommodated in an
instrumented rig whose atmosphere was partially circulating
purified helium of 2 atmospheric-pressure. The temperature
was controlled by electric heaters as well as by changing thehelium gas pressure between the subcapsule and theirradiation rig, after the reactor reached its steady-state
operation mode at 50 MW [1-3]. The guard-ring configuration
was adopted but a guard ring was connected to the ground
potential through a wall of the subcapsule accommodating
the specimen. Thin zirconium films were deposited on bothsides of a disk-like specimen as electrodes and two platinum
wires were connected to the two zirconium electrodes
working as a cathode and an anode. The guard ring electrode
was connected to a stainless-made subcapsule wall througha thick copper made specimen holder which is circumventing
a leakage electrical current from coming into a currentmeasuring-side electrode. To study protonic conductivity,
10keV proton ions were implanted into the anode sidezirconium film and oxygen ions into the cathode side.
Figure 2 summarizes temperature dependence of theelectrical conductivity of BaCe0.9Y0.1O3-δ. The electricalconductivity could be composed of two mechanisms; one
having a lower apparent activation energy of about 0.2 eV(mechanism 1; electron conduction) in the temperature range
below 480K, and the other having a higher apparent activationenergy of about 0.7-0.8 eV (mechanism 2; proton conduction).
The hydrogen implantation into the anode did not alter theelectron conductivity substantially but increased the proton
conductivity by about the two orders of magnitude. Under
JMTR irradiation, measurements of temperature dependenceof the electrical conductivity below about 400 K could not berealized due to a large nuclear heating rate and the electricalconductivity by mechanism 1 could not be evaluated during
the JMTR steady-state operation. Without the hydrogen
implantation into the anode, the increase of the proton
conductivity by 2 kGy/s irradiation was about 7 times. In themeantime, with the hydrogen implanted anode, the observed
increase of proton conductivity by the 1.1 kGy/s irradiation
was more than two orders of magnitude. Under the JMTR
irradiation, further increase of the electrical conductivity(mechanism 3) from that by the mechanism 2 was observed
at higher temperatures, namely, above 623 K for the specimen
without hydrogen and above 523 K for the specimen withhydrogen. This increase of the electrical conductivity athigher temperatures was not observed in the bench top
experiments as can be seen in Fig. 2.

Fig. 2. Temperature dependence of electrical conductivity of
BaCe0.9Y0.1O3-δ, without irradiation and under JMTR
irradiation.
References[1] T. Shikama, B. Tsuchiya, S. Nagata and K. Toh, Advancesin Science and Technology 45, 1974 (2006).
[2] B. Tsuchiya, T. Shikama, S. Nagata, K. Toh, M. Narui and
M. Yamazaki, J. Nucl. Mater. 329-333, 1511 (2004).
[3] B. Tsuchiya, S. Nagata, K. Saito, K. Toh and T. Shikama,
Materials Science Forum 480-481, 579 (2005).

Fig. 1. Schematic of the electrical conductivity in-situ
measurement under JMTR irradiation.

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IMR KINKEN Research Highlights 2007

Contact to
Tatsuo Shikama (Nuclear Materials Science Division)e-mail: shikama@imr.tohoku.ac.jp

Advanced Nuclear Energy Materials

3FTFBSDI

Quantitative Analysis of the Dependence of Hardening
on Copper Precipitate Diameter and Density in Fe-Cu Alloys
In-situ TEM observations during tensile tests of dislocation
gliding through copper precipitates in thermally aged Fe-Cu alloys
were performed. The obstacle strength was estimated from thecritical bow-out angle, φc, of dislocations. A good agreement was
obtained between hardening estimation based on the critical bowing
angles and those obtained from conventional tensile tests.
It is well known that the radiation-induced defect clusters in metals
are responsible for significant changes in mechanical properties.
Examples of extended crystal defect clusters include precipitates,
voids, dislocation lines and loops, and so on. Understanding thequantitative effect of such radiation-induced defect clusters on
dislocation motion is important for construction of predictive models toestimate the lifetime of power plant components. We investigate thedependence of hardening on copper precipitate diameter and number
density by performing in-situ TEM observations during tensile tests inthermally aged Fe-Cu alloys [1,2]. The obstacle strength α wasdetermined as α= cos(φc/2) at the point of dislocation break away fromthe obstacle. The critical bow-out angle φc was taken as the angle
between two tangent lines drawn from the cusp in the line tension
approximation. The number density of obstacles was determinedfrom the dislocation segment length L between two adjacent pinningpoints. Comparison is made between number densities estimatedfrom in-situ observations and those obtained by three dimensional
atom probe (3DAP) and conventional TEM observations.
Fe-1.0wt%Cu alloy was prepared by arc melting followed by coldrolling. All samples were annealed at 825 oC for 4 h and subsequently
quenched into ice water. Some of the samples were thermally aged at
525 oC for 20 min, 1, 10, and 100 h. In-situ tensile deformation
experiments were performed at room temperature in a JEOL 4000FXTEM operating at 400 kV with a single tilt straining holder. The motionof dislocations was recorded by a CCD camera with a time resolutionof 1/30 s. The slip planes had to be determined prior to measuring

bow-out angles of dislocation. Slip systems being activated were
determined to be <111> {101} or {112}.
Fig. 1 shows the typical morphology of dislocations during an insitu TEM observation of Fe-1.0 wt % Cu after aging 20 min, 1 h, 10 hand 100 h at 525 oC. Dislocations are along <111> direction andanchored at pinning points. Most of dislocations interacting withobstacles were of screw type. Copper precipitates were not observedby TEM after 20 min and 1 h aging, and were clearly observed at
pinning points and in the matrix after 10 h and 100 h aging. The
diameter was about 1 nm after 20 min and 1 h aging. They are
considered to have a coherent b.c.c. structure after 20 min and 1 haging. On the other hand, the precipitates were in the stage oftransformation from b.c.c. structure to 9R structure, after 10 h aging.
There is a peak in number density of copper precipitates at 1 h aging.
Sometimes small defect clusters with a black dot contrast wereobserved to form during straining. These defect clusters are debris
defects left after cross slips of dislocations.
The obstacle strength parameter, cos(φc/2) was almost sameafter 20 min and 1 h aging. For increasing the aging times to 10 h and100 h, the average obstacle strength increased with broader
distributions. The obstacle strength depends on the precipitate
diameter and “the impact parameter” which is the distance from thecenter of obstacle to a glide plane [3]. The relationship between theincrease in shear stress and the increase in yield stress is given byfollowing equation using the Taylor factor for polycrystals (M = 3.06):
σy=σ0+Δσy=σ0+MΔτc,where σy is yield stress, σ0 is yield stress for Fe-1.0wt%Cu alloy
as quenched without precipitation hardening. Fig. 2 shows a qualitative
agreement between the yield stress measured by conventional tensile
tests and that obtained from in-situ TEM observations.

Fig. 2 Precipitate number density measured from 3DAP or TEM
observation and estimated from in situ TEM observations. The yieldstress measured from tensile tests and values calculated from theexperimentally obtained microstructural data.
References[1] K Nogiwa, N. Nita and H. Matsui, J. Nucl. Mater. in press.
[2] K. Nogiwa, T. Yamamoto, K. Fukumoto, H. Matsui, Y. Nagai, K.
Yubuta and M. Hasegawa, J. Nucl. Mater. 307-311, 946 (2002).
[3] T. Hatano and H. Matsui, Phys. Rev. B72, 094105 (2005).

Fig. 1 The typical morphology of dislocations under stressduring in-situ TEM observations in Fe-1.0wt%Cu aged 20
min (a), 1 h (b), 10 h (c) and 100 h (d) at 525 oC.

Advanced Nuclear Energy Materials

Contact to
Hideki Matsui (Nuclear materials engineering Division)e-mail: matsui@imr.tohoku.ac.jp

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3FTFBSDI

Intrinsic Martensite Formation in Neutron Irradiated V-1.6%Y
Alloys with Fine-grained Structure of Highly Pure Matrix
It is known that in vanadium gaseous interstitials (H, N,O) cause transformation from the bcc to bct structure byelongating one axis relative to the other two to accommodate
the induced strain. This paper describes the first finding of
intrinsic martensite formation that is not associated with theinterstitials for neutron irradiated V-1.6Y (in wt%) with a finegrained structure of a highly pure matrix [1]Development of refractory transition metals with improved
resistance against mechanical property degradation by highenergy particle irradiation is required for their use underirradiation environments. It has been shown that the mosteffective microstructure to improve the radiation resistanceconsists of fine grains and finely dispersed particles, the finer
the better because grain boundaries and particles can serve
as effective sinks for irradiation induced point defects. For
vanadium such a microstructural refinement requires aprocessing method that enables to make the vanadium matrix
free from gaseous interstitials, particularly nitrogen andoxygen. This is because vanadium is chemically very reactivewith the interstitials and their dissolution into the matrix leads
to significant embrittlement. We have proposed an advanced
powder metallurgical (P/M) method utilizing mechanicalalloying and hot isostatic pressing and successfully developed
ductile V-Y and V-Y-Cr alloys with fine grains of a highly purematrix free from gaseous interstitials and very fine dispersoidsof Y2O3 and YN [2-4]. The Y2O3 and YN dispersoids areformed as a result of consuming solute oxygen and nitrogen
impurities.
V-1.6Y alloys were irradiated at 290 and 600°C to 0.25
and 0.6 displacement per atom (dpa), respectively, in theJapan Materials Testing Reactor to examine the effect ofneutron irradiation. It has been demonstrated that the alloys
exhibit much improved radiation resistance [5] Anothernoticiable finding is that the irradiated V-1.6Y exhibits theintrinsic martensite formation of a vanadium bct structure thatis not caused by gaseous interstitials. The emergence of
intrinsic martensite is the first to be observed in vanadium or
its alloys not only in neutron irradiated states but also in
unirradiated states. The main results and conclusions are as
follows [1]:1. The martensite emerges heterogeneously in some of
the bcc grains of V-1.6Y and has a high density of {101)bct
microtwins as the lattice invariant shear strain. The bct
structure exhibits a tetragonality (a/c) of 1.06 as measured
with Kikuchi line analyses.
2. The habit planes between the bct martensite and thesurrounding bcc matrix are {101)bct // {110}bcc, which can be
regarded as the close-packed plane with coherency even forthe bct structure because of the slight changes in the lattice
constants of both the phases.
3. The features of the P/M V-1.6Y alloys, an interstitialscavenged matrix and extremely fine grains, are considered
to be responsible for the martensite formation caused byirradiation.
4. The transformed vanadium bct lattice provides

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preferential sites as the host phase for hydrogen atoms tooccupy with an ordered arrangement (Fig. 1). The orderedarrangement of hydrogen atoms occurs in a very thin
specimen area that is susceptible to hydrogen pickup
associated with TEM specimen preparation after irradiation,
but does not occur in a thicker specimen area.

Fig. 1 TEM micrograph of dark field image from a very thin
specimen area with a transformed region that provides
preferential sites for hydrogen atoms to occupy with anordered arrangements in V-1.6Y irradiated at 290°C to
0.25dpa.
5. The hydrogen content (H/V) in the ordered bct structureis estimated to be approximately H/V = 0.24. This content ismuch lower than H/V~ 0.41, the lowest content of the β1
hydride phase that can exist in the thermoequilibrium state,
supporting the absence of hydride formation in the bct
phase.
References[1] H. Kurishita, S. Kobayashi, K. Nakai, T. Kuwabara and M.
Hasegawa, J. Nucl. Mater. 358, 217 (2006).
[2] T. Kuwabara, H. Kurishita and M. Hasegawa, J. Nucl.
Mater. 283-287, 611 (2000).
[3] T. Kuwabara, H. Kurishita and M. Hasegawa, Mater. Sci.
& Eng. A 417, 16 (2006).
[4] H. Kurishita, T. Kuwabara and M. Hasegawa, Mater. Sci.
& Eng A. 433, 32 (2006).
[5] H. Kurishita, T. Kuwabara, M. Hasegawa, S. Kobayashi
and K. Nakai, J. Nucl. Mater. 343, 318 (2005).
Contact toHiroaki Kurishita (International Research Center for Nuclear
Materials Science)e-mail: kurishi@imr.tohoku.ac.jp

Advanced Nuclear Energy Materials

3FTFBSDI

Development Study of Fast Reactor Core withHydride Neutron Absorber
We are now developing hafnium hydride control rods6,7
for Fast Breeder Reactor (FBR) in order to drasticallyprolong its in-core life accounting its superb characteristicsof helium and swelling free features.A new hydride neutron
absorber of nuclear reactor has proposed based on the
results of irradiation experiments in Japan Materials Testing
Reactor (JMTR) of Japan Atomic Energy Agency (JAEA). A
metal-hydride has very high hydrogen atom density, whichis equivalent to that of liquid water. Fast neutrons in nuclear
reactors are efficiently moderated and are absorbed in themetal-hydride. The Hafnium hydride and Gd hydride areconsidered as neutron absorber in FBR.

is larger than that of B10 40% enriched B4C.
One of the most important R&D items to establish theidea of Hf-hydride control rod is the development of Hfhydride pellet, which can be safely irradiated for long time.
The idea has proposed based on the results of irradiation
experiments in JMTR of JAEA. Fig. 2 shows the structure ofcapsule for irradiation experiments, where the actinide
hydride pellet were loaded. After irradiations, non-destructive
examinations (X-ray photograph, gamma scanning and soon) and destructive examinations (electron microprobe
analysis, physical property measurements and so on) were
performed and the integrity of the hydride pellet during
irradiation was assured.

Fast breeder reactor is considered as the major nuclear
power source for the future. The B4C has been mainly used
for control and shut down material in FBRs. But the life of
B4C control rods is restricted by Pellet-Cladding MechanicalInteraction (PCMI) failure due to the He gas swelling of B4C
pellet which is caused by the following nuclear reaction,
B10 (n, α) Li7.

(1)

Fig. 2 Schematic of HfHx application concept to FBR control rod.

The concept how to prolong control rod life of FBR byusing HfHx is illustrated in Fig. 1. In order to prolong thecontrol rod life–time we propose to use HfHx as absorbermaterial for FBR by the reasons that He gas is not generated
in nuclear reaction of HfHx, and that HfHx can be expected toabsorb neutron for more than 40 years owing to the fact that
Hf178 and Hf179 which are generated by neutron captures of
Hf177 and Hf178 respectively, also have large neutron capture
cross sections which are expressed by the following
reaction,.

The development program of hydride neutron absorberhas been started, which is accepted as an innovative nuclear
research and development program of Ministry of Education,
Culture, Sports, Science and Technology of Japan.

Hf177 (n, γ) Hf178 (n, γ) Hf179 (n, γ) Hf180.

(2)

References[1] K. Konashi et al., “Development of advanced control rodof hydride hafnium for fast reactor”, Proc. of the ICAPP2006,
Reno, USA (2006).
[2] K. Konashi et al., “Study on application of hafnium hydride
control rods to fast reactors”, Proc. of the 15th Pacific Basin
Nuclear Conference, Sydney, Australia (2006).
[3] K. Konashi and M. Yamawaki, “The development of thorium
hydride fuel”, Characterization and quality control of nuclear
fuels edited by Ganguly C. and Jayaraj R.N., 92, Allied
Publishers Pvt. Ltd. (2004).
Contact toKenji Konashi (International Research Center for Nuclear
Materials Science)e-mail: konashi@imr.tohoku.ac.jp

Fig. 1 Schematic of HfHx application concept to FBR control rod.
The control rod assembly, which contains Hf-hydride, isconsidered to enhance neutron absorption in fast reactors.
In general, a metal-hydride has very high hydrogen atom
density, which is equivalent to that of liquid water. Fastneutrons in nuclear reactors are efficiently moderated andare absorbed in the metal-hydride. As the ratio of H to Hf
increases, the Hf-control rod can more efficiently absorb
neutrons in fast reactor core. For example, the worth of HfH1.0

Advanced Nuclear Energy Materials

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3FTFBSDI

Defect Structure of Nonstoichiometric Plutonium Oxide
Nuclear fuel systems typically require decades tooptimize, since the assessment of behavior of nuclear fuel ismade based on empirical data and irradiation testing. As ageneral trend, it is tried to assess them on basic scientificunderstanding. In this work, an oxygen defect of plutoniumoxide has been studied by the first principle calculations,which predict the electron structure of materials without
experimental data.
Plutonium, which is generated in an operating nuclear
reactor due to transmutation of uranium, is one of importantmaterials as nuclear fuel. Plutonium is mainly used in the
nuclear reactor in the chemical form of oxide. Plutoniumoxide has a wide range of non-stoichiometry, determined bytemperature and oxygen partial pressure. Non-stoichiometricplutonium oxide is represented by PuO2-x, where x is nonstoichiometry. The oxygen partial pressure is a key parameter,
which controls chemical reactions in the fuel pin.
It is very important to develop a model, which is able topredict well the correlation between the non-stoichiometry of
PuO2-x and the partial pressure of oxygen. In the design of
nuclear fuel, the model based on defect chemistry has beenused for calculation of oxygen potential. The defect chemistrymodel is essentially empirical one, since the parameters in
the defect chemistry model are adjusted by the experimental
data. In this study, the first principle calculations were used tocheck the validity of the defect chemistry model.

Fig.1 Supercel for plutonium oxide with oxygen defect. Red
and white spheres indicate plutonium and oxygen atoms,
respectively.
The total energy and the electronic structure of PuO2-x
have been calculated by PAW method within the GGA. We
also used the DFT+U methodology to check localization off-electron. Figure 1 shows a supercell of 96 atoms containing
32 plutonium atoms, 63 oxygen atoms and one oxygen
vacancy. The calculations were performed with the HITACHI/
SR-8000 in Tohoku University. The calculated lattice
parameter and the enthalpy of formation of plutonium dioxide
are in good agreement with the experimental values. The

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vacancy formation energy, which was defined as the energy
required for extraction of an oxygen atom from the perfectlattice, was calculated to be 4.20eV. This value is close tothat calculated from the experimental data.

Fig. 2 Differential charge densities near oxygen vacancy
Difference in the charge density between the defectmodel and the perfect model has been calculated to makethe effect of vacancy clear (Fig. 2). The six nearest neighbor
oxygen atoms moved from original site toward the oxygen
vacancy site. The two atoms of those were shown in Fig. 2. It
is also shown that the charge densities around the twonearest neighbor plutonium atoms are affected by the oxygen
vacancy. The calculation results show that only the nearest
neighbor plutonium atoms are affected by the oxygen
vacancy. This agree with the following picture used in theclassical defect chemistry, that is, the oxygen has twoelectrons provided by plutonium, which are left behind whenan oxygen atom is leaving lattice position. These twoelectrons are localized on two plutonium atoms, turning Pu+4
into Pu+3.
References[1] K. Konashi H. Matsui, Y. Kawazoe, M. Kato and S.
Minamoto, J. Phys. Soc. Jpn. 75 suppl., 143 (2006).
[2] K. Konashi, M. Kato, S. Minamoto, ‘First Principle PlaneWave Pseudopotential Calculation of Point Defects in PuO2X’, ANS 2005 Winter Meeting, Nov. 15-19, 2005, Washington,
D.C., USA
Contact toKenji Konashi (International Research Center for Nuclear
Materials Science)e-mail:konashi@imr.tohoku.ac.jp

Advanced Nuclear Energy Materials