Energy-related Materials

Impact of Al-doped ZnO layer on growth rate andmicrostructures in polycrystalline Si thin film obtained by
Al-induced layer exchange process

Buried Al-doped ZnO (AZO) layer was found to give a drastic impact on Al-induced layer exchange

process of amorphous Si (a-Si) and Al, and resultant formation of polycrystalline Si thin film. The growthrate of polycrystalline Si was largely enhanced by the presence of AZO through the increased
interface roughness between a-Si and Al, which promoted diffusion of Si in Al. Furthermore, a fractionof <100>-oriented crystal grains was found to increase with the presence of AZO.

Realizationofwafer-equivalentpolycrystalline Si (poly-Si) thin film on largearea low-cost substrates will give a great
impact since it could dramatically improvesemiconductor devices such as solar cells andflat panel displays. The film should be of highquality containing large crystal grains with
controlled grain orientations and grainboundary characters. Especially, if one could
realize <111> or <100>-oriented wafer-equivalentpoly-Si thin films, they can be used for subsequentepitaxy of a-axis oriented BaSi2 thin films, which
could be applied to novel high-efficiency,
earth-abundant thin film solar cells [1].
A possible growth method to realize such ahigh-quality poly-Si thin film is Al-induced layer
exchange process (ALILE) of Al and amorphousSi (a-Si) [2]. ALILE starts with stacked structuresof a-Si/oxide/Al/substrate followed by annealing
below the eutectic temperature (577 oC) in the
Al-Si system. After annealing, Al and a-Si layers
are exchanged and the initial a-Si is transformed
into the continuous poly-Si film on the substrate[3-5]. Since ALILE proceeds under nearequilibrium conditions with a small amount ofsupersaturation, the nucleation rate is relatively
slow and large crystal grains in the order of 10100 nm can be obtained. The grain size andthe preferential orientation are strongly
dependent on the interfacial structures [6]since dissociation of Si atoms in Al proceeds
through the a-Si/Al interface. However,controlling mechanisms of microstructures aswell as the growth rate have not been fully
understood in spite of the importance.
In order to explore these mechanisms, we
attempted to modify the interface by using
buried aluminum-doped zinc oxide (AZO) toact as the transparent conductive layer andcarried out in-situ observation of the growth
process [7,8]. Glass substrates were cleaned
with acetone, methanol, and deionized water.
For a part of samples, 650 nm-thick AZO wasdeposited by radio-frequency magnetron
sputtering. The 100 nm-thick Al layers were thendeposited using a thermal evaporator.
Subsequently, samples were exposed to air for
48 hrs to form the native Al oxide layer followed
by deposition of 100 nm-thick a-Si layer. Finally,19

IMR KINKEN Research Highlights 2011

samples were annealed at 500-550 oC underargon gas ambient. The growth process wasmonitored in-situ by using a system with a digital
optical microscope and a furnace..
Fig. 1 shows snapshots during growth ofpoly-Si, which indicate surface coverage of Si
as a function of the annealing time. The top
and bottom images of Fig. 1 express the poly-Si
films formed without/with presence of AZO
layer. At the same annealing temperature, thecrystallization process time of samples with AZO
was found to be much shorter than that without
AZO. With increasing annealing temperature,
the crystallization process time was also rapidly
shortened.
One of the possible reasons for the faster
growth rate in the sample with AZO is
acceleration of the diffusion rate of Si into Al
layer. By atomic force microscopy, we havedisclosed that the root-mean-square roughnessof the glass/Al and glass/AZO/Al layers before
annealing was 2.0 and 3.1 nm, respectively.
The initial interface roughness affected theFig. 1 Snapshots during growth of poly-Si display
increase of the Si coverage as a function of
annealing time for samples without and with AZO,
and different annealing temperatures of 500 oC
and 550 oC.

IMR KINKEN Research Highlights 2011

Fig. 2 Cross-sectional scanning electron microscope
images of samples (a) without and (b) with AZO
after annealing at 500 oC

roughness after annealing as shown in Fig. 2.
The Al/Si interface of sample without AZO layer
is relatively flat and clean, whereas that ofsample with AZO layer has a remarkably large
roughness. Two-dimensional finite element
calculation supported that the increased
interface roughness results in the increase in
the Si concentration in Al, which enhances thegrowth rate of poly-Si [8].
Fig. 3 shows the crystalline orientation
mapping images for the surface normal
direction to poly-Si films. Remarkable variationof crystallographic orientation for each sample
was found depending on growth parameters.
The presence of the AZO layer and the increaseof the annealing temperature were found tolead to increase of <100>-oriented crystal
grains. On the other hand, lower annealing
temperature was found to result in increase of<111>-oriented crystal grains. Here, we introduce
a possible mechanism of the preferential
orientation of poly-Si films, which can be
determined from the growth rate. Generally,
the Si{111} plane with the lowest surface energy
tends to have higher stability; however, theSi{100} plane has a higher surface energy with

Fig. 3 EBSD crystalline orientation mapping images
for poly-Si film

unstable state. In the case of low growth rate,the {111} coverage with lower surface energy islarger than the coverage of other planes. On
the other hand, when the growth rate is fast,the {100} coverage with a rough surface isincreased. As illustrated by this study, control of
the interface can modify ALILE andmicrostructures in poly-Si thin film [9], which
could be a promising route toward realizationof wafer-equivalent poly-Si thin film.
References[1] T. Suemasu and N. Usami, Japanese Patent 2009-115337[2] G. Majni, and G. Ottaviani, Appl. Phys. Lett. 31, 125 (1977).
[3] O. Nast, T. Puzzer, L. M. Koschier, A. B. Sproul, and S. R.
Wenham, Appl. Phys. Lett. 73, 3214 (1998).
[4] A. G. Aberle, N. P. Harder, and S. Oelting, J. Cryst. Growth226, 209 (2001).
[5] W. Fuhs, S. Gall, B. Rau, M. Schmidt, and J. Schneider,
Solar Energy 77, 961 (2004).
[6] M. Kurosawa, N. Kawabata, T. Sadoh, and M. Miyao,
Appl. Phys. Lett. 95, 132103 (2009).
[7] M. Jung, A. Okada, T. Saito, T. Suemasu, and N. Usami,Appl. Phys. Express 3, 095803 (2010 ).
[8] M. Jung, A. Okada, T. Saito, T. Suemasu, C.Y. Chung, Y.
Kawazoe, and N. Usami, Jpn. J. Appl. Phys. 50, 04DP02 (2011).
[9] H. Suzuki, N. Usami, A. Nomura, T. Shishido, K. Nakajima,
and T. Suemasu, J. Cryst. Growth 312, 3257 (2010)

Keywords: crystal growth, thin films, solar cells
N. Usami (Crystal Physics Division)
E-mail:usa@imr.tohoku.ac.jp
URL:http://www.xtalphys.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201120

Energy-related Materials

Observation of the Spin Seebeck Effect in Insulators

Spin Seebeck effect (SSE), generation of spin voltage as a result of a temperature gradient, is shown

to appear in magnetic insulators, which has been measured only in metals. The observation of the
SSE in insulators has been demonstrated in two different configurations. Our findings extend therange of potential materials for thermoelectric applications and provide a crucial piece of
information for understanding the physics of the SSE.

21

IMR KINKEN Research Highlights 2011(a)Electricity

Magnetism

Seebeck effectspin Seebeck effectConductor

V

Vs
Vs

∇T

∇T
ferromagnetic metal

metal or semiconductor

spin Seebeck effectVs

Insulator

For thermoelectric power generation, we
often rely on the Seebeck effect, which is thegeneration of electric voltage as a result of atemperature gradient. In contrast, recent
studies on spintronics revealed that a spin
analogy for the Seebeck effect, named thespin Seebeck effect (SSE) [1-9], appears inmagnets. The SSE stands for the generation of‘spin voltage’ as a result of a temperature
gradient. Spin voltage refers to the potential forelectron’s spins, which drives spin currents. The
SSE has been reported in several metallic
magnets [1], although consensus about its
microscopic mechanism has not yet beenreached [2].
The conventional Seebeck effect requires
itinerant charge carriers, and therefore exists
only in electric conductors. It appears natural
to assume that the same should hold for the
SSE. Here, we show that a magnetic insulator
exhibits the SSE despite the absence ofconduction electrons (Fig. 1(a)) [3]. This resultclearly indicates that thermally induced spin
voltage is associated with magnetization
dynamics. This new mechanism for the thermospin conversion allows us to narrow down
scenarios for the physical origin of the SSE.
The observation of the SSE in a magnetic
insulator is realized by the inverse spin Hall
effect (ISHE) in Pt films [3]. The ISHE enables toconvert the spin signal generated by heat flows
in the insulator into electric voltage. Fig. 1(b)shows a schematic of the sample used in thisstudy. The sample consists of a single-crystal
insulator LaY2Fe5O12 film with several Pt wires.
An external magnetic field H and a temperature
gradient were applied along the x direction(Fig. 1(b)). Since the localized magnetic
moments in LaY2Fe5O12 and conduction
electrons in Pt are coupled by the interface
spin-exchange interaction [10], the SSE in the
LaY2Fe5O12 layer should generate a spin current
in the attached Pt layer (Fig. 1(b)). In the Pt
layer, the ISHE converts this spin current into anelectric field via the spin-orbit interaction.
Therefore, we can detect the SSE by measuring
electric voltage V between the ends of the Pt
wire.
In Fig. 2(a), we show the H dependence of
V in the Pt wires, placed at the lower- andhigher- temperature ends of the LaY2Fe5O12
layer, for various values of the temperatureVs∇T
magnetic insulator

(b)
Pt

V
EISHE

V
Pt
H

σ
Js

T + ∆T
T (300 K)

LaY2Fe5O12

z
y

x

M

∇T
y

z
x

LaY2Fe5O12∇T

Fig. 1: (a) Seebeck and spin Seebeck effects (SSE).
(b) Measurement mechanism for the SSE in an
insulator.

difference ΔT. The magnitude of V is observed
to be proportional to ΔT in both the setups. The
sign of V at finite values of ΔT is clearly reversed
between these setups and is also reversed byreversing H; these distinctive behaviors of V are
the key features of the SSE [1]. Since the
LaY2Fe5O12 is an insulator, complicating
thermoelectric artifacts in itinerant magnets
do not exist at all. These results confirm theexistence of the spin current injected into thePt layer which is generated from thetemperature gradient in the insulator LaY2Fe5O12[3].
Next, we measured the spatial distributionof the thermally induced spin voltage in thesame sample structure. Fig. 2(b) shows V as afunction of the Pt-wire position for various
values of ΔT. When a finite ΔT is applied, V varies
linearly along the temperature gradient, similar
to the behavior of the SSE in metallic films [1,3].
We demonstrated the observation of the
SSE in a magnetic insulator also in longitudinal

IMR KINKEN Research Highlights 2011

T(300 K)

V (µV)

Higher T

25 K

25 K

20 K

20 K

15 K

15 K

10 K

10 K

5K

5K

∆T = 0 K

00
H (Oe)

100

(b)
2
1
V (µV)

-100

0

2.8

0
H (Oe)

9.2 K

θ
∇T

10
∆T (K)zxy13.8 K

H

4.6 K
0

20

∆T = 0 K-1

0
H (kOe)1Fig. 3: Temperature-difference ΔT dependence (a)and H dependence (b) of the longitudinal SSE
signal V in the Y3 Fe5O12/Pt sample.

-2
-1

0

100

∆T = 20 K
10 K
0K
-2

V

18.4 K

xPt

-1

-3

Pt

Y3Fe5O12

0

-4

θ=0

0

-20

。

θ = 9023.0 K。

-10∆T = 0 K

10.0

θ = 90

2 µV

H

∇T

-2.8

10

V (µV)

2 µV(b)

H = 1 kOe

T + ∆T

LaY2Fe5O12

Lower T-100(a) 20Pt
H

V (µV)

∇T

(a)

0
1
xPt (mm)2

3

4

Fig. 2: Magnetic field H dependence (a) Pt-wireposition xPt dependence (b) of the SSE signal V in
the LaY2Fe5O12/Pt sample.

configuration, in which a spin current is injected
parallel to a temperature gradient from aferromagnet into an attached paramagnetic
metal [5, 7]. In Fig. 3, we show the experimental
results on the longitudinal SSE in an Y3Fe5O12/Pt
system. Since the longitudinal SSE requires only
simple and versatile systems, it will extend therange of device application and experimental
investigation of the SSE.
The SSE in magnetic insulators can beapplied directly to the design of thermo-spin
generators and, in combination with the ISHE,
thermoelectric generators, allowing new
approaches towards the improvement ofthermoelectric generation efficiency. In
general, the efficiency is improved bysuppressing the energy loss due to heat
conduction and Joule dissipation, which arerealized respectively by reducing the thermal
conductivity κ for the sample part where heatcurrents flow and by reducing the electric
resistivity ρ for the part where charge currents
flow. However, in electric conductors, theWiedemann-Franz law (κρ~constant) limits thisimprovement. In contrast, in the present SSE
setups, κ and ρ are independent since theheat and charge currents flow in different partsof the sample [3, 5, 7]. The thermoelectric
generation efficiency of our SSE devices is still
small, but can be enhanced by reducing thephonon contribution to the thermal conductivityof the insulator layer and by improving theinterface spin-exchange efficiency.
References[1] K. Uchida, S. Takahashi, K. Harii, J. Ieda, W. Koshibae, K.
Ando, S. Maekawa, and E. Saitoh, Nature 455 (2008) 778.
[2] J. Xiao, G. E. W. Bauer, K. Uchida, E. Saitoh, and S.
Maekawa, Phys. Rev. B 81 (2010) 214418.
[3] K. Uchida, J. Xiao, H. Adachi, J. Ohe, S. Takahashi, J.
Ieda, T. Ota, Y. Kajiwara, H. Umezawa, H. Kawai, G. E. W.
Bauer, S. Maekawa, and E. Saitoh, Nature Mater. 9 (2010)894.
[4] C. M. Jaworski, J. Yang, S. Mack, D. D. Awschalom, J. P.
Heremans, and R. C. Myers, Nature Mater. 9 (2010) 898.
[5] K. Uchida, H. Adachi, T. Ota, H. Nakayama, S. Maekawa,
and E. Saitoh, Appl. Phys. Lett. 97 (2010) 172505.
[6] H. Adachi, K. Uchida, E. Saitoh, J. Ohe, S. Takahashi, and
S. Maekawa, Appl. Phys. Lett. 97 (2010) 252506.
[7] K. Uchida, T. Nonaka, T. Ota, and E. Saitoh, Appl. Phys.
Lett. 97 (2010) 262504.
[8] C. M. Jaworski, J. Yang, S. Mack, D. D. Awschalom, R. C.
Myers, and J. P. Heremans, Phys. Rev. Lett. 106 (2011)
186601.
[9] K. Uchida, H. Adachi, T. An, T. Ota, M. Toda, B. Hillebrands,
S. Maekawa, and E. Saitoh, Nature Mater. (in press).
[10] Y. Kajiwara, K. Harii, S. Takahashi, J. Ohe, K. Uchida, M.
Mizuguchi, H. Umezawa, H. Kawai, K. Ando, K. Takanashi,
S. Maekawa, and E. Saitoh, Nature 464 (2010) 262.

Keywords: spintronic, spin current, thermoelectricity
K. Uchida and E. Saitoh (Materials Property Division)
E-mail: saitoheiji@imr.tohoku.ac.jp
URL: http://saitoh.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201122

Energy-related Materials

Advance and Recent Activity of Neutron Group of IMRI

MR is a unique institute even in the world which has a neutron facility based on the strong backgroundof advanced material science. IMR neutron scattering group is investigating issues in materials
science, such as ferroelectric materials, high-TC superconductivity, as well as developing highperformance devices in neutron science. In particular, the group is aiming at constructing an
originative neutron spectrometer with oncoming-generation technologies in J-PARC, which will bea flagship instrument of material science in IMR.
1) Advanced Spectrometer Project(K. Ohoyama)
Metal Physics with Quantum Beam
Spectroscopy group (Yamada Lab.) andCenter of Neutron Science for Advanced
Materials are advancing a project to constructa world top class spectrometer at the brightest
neutron beam facility of Japan Proton
Accelerator Research Complex (J-PARC),
based on collaborations with KEK (project
leader: K. Ohoyama) (Fig.1).
The spectrometer is characterised bypolarised neutrons, which are very sensitive toobserve pure magnetic correlations. Since there
exist just a few polarised neutron spectrometersin the world, this spectrometer will be a key
instrument to generate breakthroughs in novel
material science in Japan. In 2010, the project
team succeeded in preliminary polarised
neutron experiments, using a 3He gas typeneutronspin
polarisersdevelopedbycollaborations with KEK, JAEA and IMR. Now,
the spin polariser/analyser for this project is in
the progress of production. Moreover, we haveestimated performances of the spectrometer,
such as flux, resolution, and polarisation, basedon test measurements on IMR neutron
instruments, and Monte Carlo simulation. By
optimization of designs, satisfactory performance
have been obtained; the polarised neutron fluxof the spectrometer is expected to be 1.1E+5(n/s/cm2/meV) at 100meV, and the neutron
polarisation to be ~90%.
To make this spectrometer originative inthe world, we have been developing anepoch-making
technology:
the
cross
correlation method, which will give drastic
improvement of efficiency and S/N ratio of
measurements of spectroscopy.
In 2011, this project succeeded in gaining
authorisation by J-PARC Center. Aiming at
beginning the construction from 2012, detailed
engineering designs have already begun.
2) Mechanism of Relaxor Ferroelectrics
(M. Matsuura)
Ferroelectric materials are widely used inour daily life, such as condensers in cellular
phone, computer, etc., almost all electric
devices. Relaxor ferroelectrics have gainedhuge attention from industrial and scientificpoints of view, because they show extremely
high dielectric susceptibility and piezoelectric
responses over a wide temperature range[1]which enable thinner and smaller condensers
with the same quality.
In relaxor ferroelectrics, it is known that
randomly oriented polar nanoregions (PNRs)
are formed from high temperatures[2], asshown schematically in Fig.2a, and play animportant role in the giant dielectric response
and its thermal evolution. From inelastic neutron
scattering studies on a single crystal of relaxor
ferroelectric
0.7Pb(Mg1/3Nb2/3)O3 -0.3PbTiO3,
we have clarified that flipping motion of electric
dipoles (relaxation) in PNRs are coupled withlattice vibrations (phonon) in surrounding
regions[3], as schematically shown in Fig. 2b.
Through the mode-coupling between phonon
andrelaxation
modes,
some
regions

Detector Bank
Stair to Upper Floor

Spin Analyser

Entrance

Beam Stopper

Spin Polariser

Radiation Shield
Beam Transport Tube
Neutron BeamFermi Chopper

Sample Position
Radiation ShieldFig. 1 A rendering view of the spectrometer of the collaboration project of Tohoku Univ. and KEK.

23

IMR KINKEN Research Highlights 2011

IMR KINKEN Research Highlights 2011

Fig.2: (a) Schematic diagram of polar nanoregion(PNR). Arrows indicate direction of electric dipoles
in PNRs. (b) Schematic diagram of transverse
acoustic (TA) and soft transverse optical (TO)phonondispersion
curvesinrelaxor
0.7Pb(Mg1/3 Nb2/3)O3 -0.3PbTiO3 (0.7PMN-0.3PT). The
gradated thick lines indicate the q-dependence ofthe relaxation frequency at T=600K (red) and 400K
(blue). [3]surrounding PNRs become polar by atomic shift
along an electric dipole of each PNR from
ideal cubic positions, which leads toenlargement of PNRs. The new growthmechanism of PNRs not only explains therelaxor properties over a wide temperature
range but also unifies different approaches
based on order-disorder-type and displacivetype lattice dynamics.
3) Development of Novel Neutron Devices
(H. Hiraka)
A development of neutron monochromatoris one of the instrumental key issues in neutronscience. Stimulated by recent development inplastically deforming techniques for Si-crystal
wafers by Nakajima et al. [4], we explored its
potential of a deformed-wafer crystal for a novel
neutron monohromator or analyzer crystal.

Fig. 3. Rocking curves of (111) Bragg reflection using
deformed Si wafers with R=600 mm. The reflected
intensity is normalized by the direct-beam intensity.
For reference, a resolution-limit Bragg reflection ofa flat wafer is also plotted by open squares. Thelines are drawn as a guide to eyes.

Fig. 3 shows that the angle-integrated
intensity is amplified by a factor of ~4 for the
0.5-mm wafer during the deformation
process[5]. Moreover, the scattered intensity
successively increases as the thickness of thedeformed-wafer stack (ts) increases. Therefore,
further stacking can provide a way of increasingreflected neutrons more. Besides, this plastic
deformation has a pronounced advantage informinganintensecurvature,unlikemechanically (or elastically) bent crystals. We
are planning to equip these unique wafers forneutron analyzer elements because a relatively
wide-angle divergence of neutron beam is
acceptable for a detector. We have taken out
a patent on this technique[6].
References[1] S.-E. Park and T. R. Shrout: J. Appl. Phys. 82 (1997) 1804.
[2] G. Burns and F. H. Dacol: Solid State Commun. 48 (1983)853.
[3] M. Matsuura et al., submitted to JPSJ
[4] K. Nakajima et al., Nat. Mater. 4 (2005) 47.
[5] H. Hiraka et al., Nucl. Instr. And Meth. A 635(2011)137.
[6] H. Hiraka et al. Japan patent 2009-066185.

Keywords: J-PARC spectrometer project, relaxor ferroelectrics, novel neutron devices
K.Yamada, K.Ohoyama, H. Hiraka, M. Matsuura, T. Sato, and J. Sungdae (Metal Physics with Quantum Beam
Spectroscopy)
URL: http://www.yamada-lab.imr.tohoku.ac.jp/jp/index.html

IMR KINKEN Research Highlights 201124

Energy-related Materials

Microstructure Changes in Nuclear Reactor Pressure
Vessel Steels by Neutron Irradiation and their Effects on
Mechanical Properties

Understanding of the embrittlement mechanism of nuclear reactor pressure vessel (RPV) steels by
neutron irradiation is one of the most important subjects for safety operation of nuclear power plants.
Here, we investigate the correlation of microstructure evolution and hardening in RPV steels irradiatedin a material testing reactor. Three-dimensional atom probe and positron annihilation are used toclarify the contributions of solute nano-clusters and matrix damages to the irradiation hardening.
Irradiation-induced changes of the atomic distributions of solute and impurity elements aroundcarbides were also revealed.

The irradiation-induced embrittlement of
the nuclear reactor pressure vessel (RPV) steels
is currently a vital issue for ensuring the safe
operation of nuclear power plants. The
embrittlement is considered to be mainly
caused by the irradiation-induced changes in
the microstructure such as the formation of thesolute nano-clusters (SCs) and matrix damages(MDs) due to irradiation-induced point defectclusters and dislocation loops. However, in spiteof the extensive studies, the details of the
microstructural changes and their effects on
the mechanical properties are not wellclarified.
One of the reasons is due to the difficulty inthe observation of the microstructural changes
even by state-of-the-art transmission electron
microscopy. In this study, we overcome thedifficulty by employing three-dimensional atom
probe (3DAP) and positron annihilation (PA).
3DAP, which can map out the alloying elementsin three-dimensional space with near-atomic
resolution, enables us to reveal the detailmorphology of the SCs. However, 3DAP gives us
no information about MDs. On this issue, PA
technique is a powerful tool for vacancyrelated MDs. Here, we investigate the correlationof microstructure evolution and hardening in
the RPV steels by neutron irradiation using 3DAPand PA to clarify the contributions of SCs and
MDs to the irradiation hardening.
The samples are A533B type steels withhigh Cu concentration (0.16wt.%) irradiated in
the JMTR (Japan Materials Testing Reactor) in
the wide range of doses (0.32 - 9.9 × 1019 n cm -2)
under almost the same flux (1.6-1.9 × 1013 n cm -2s-1) at 290 °C. Fig. 1 shows the dose dependenceof irradiation hardening (variation of Vickers
microhardness). Rapid hardening at low doses
below 1 × 1019 n cm -2 and gradual hardening
with further increasing irradiation dose are
exhibited.
Typical elemental maps of solute atoms inthe samples subjected to different doses
obtained by 3DAP are shown in Fig. 2 [1]. These
maps reveal approximately spherical Cu rich(Mn, Ni, Si cosegregated) SCs with about 1-2
25

IMR KINKEN Research Highlights 2011

nm radius. We can estimate the contribution to
the hardening due to the SCs based on RussellBrown model [2] using the obtained 3DAP data
(sizes and volume fraction of the SCs). The
results are shown in Fig. 1, suggesting that theSCs formation is responsible for the origin of thegradual hardening above 1 × 1019 n cm -2, while
the initial rapid hardening below 1 × 1019 n cm -2
could be due to the other origin: i.e. MDs. This
claim is strongly supported by PA experiments
because the positron lifetime rapidly increases
with the dose below 1 × 1019 n cm -2, due to thepositron trapping at MDs [1].

Fig. 1: Dose dependence of irradiation hardening
for A533B steels. Estimated SC contribution for thehardening using 3DAP data based on RussellBrown model [2] is also shown.

We have also studied irradiation-induced
microstructural changes of the atomic
distributions of solute and impurity elements
around carbides in a RPV steel, a surveillance
test specimen of Doel-4 reactor in Belgium,using 3DAP [3]. Fig. 3 shows the elemental mapsof solute atoms around the carbide in theunirradiated specimen. Nano-scale Fe-Mn-CrMo carbides (~M6C type) are observed. Mn, Cr
and Mo were enriched around the innercarbide-matrix interface. The depletion of Si inIMR KINKEN Research Highlights 2011

the carbide and the segregation of P at theinterface are also observed.

Fig. 2: Elemental maps of Cu, Ni, Mn and Si atoms in
A533B type steels neutron-irradiated with various
doses.

By neutron irradiation for 12 years, the Mnconcentration within the carbide is increased
and the enhancement of Mn, Cr and Mo
concentration around the carbide- matrix
interface as well as P segregation at theinterface are markedly intensified [3].

These changes could be one of thepossibilities for degradation of the interface,which would contribute to the embrittlementof RPV steels.

Fig. 3: Elemental maps of C, Si, Mn, P, Ni, Cr and Mo
atoms around the carbide in the unirradiatedDoel-4 specimen.

References[1] T. Takeuchi, A. Kuramoto, J. Kameda, T. Toyama, Y. Nagai,
M. Hasegawa, T. Ohkubo, T. Yoshiie, Y. Nishiyama and K.
Onizawa,J. Nucl. Mater. 402, 93 (2010)[2] K.C. Russell, L.M. Brown, Acta. Metall. 20, 969 (1972)[3] T. Toyama, N. Tsuchiya, Y. Nagai, A. Almazouzi, M.
Hatakeyama, M. Hasegawa, T. Ohkubo, E. van Walle, R.
Gerard, J. Nucl. Mater. 405, 177 (2010)

Keywords: neutron irradiation, embrittlement, atom probe tomography
Y. Nagai (Irradiation Effects in Nuclear and Their Related Materials)
E-mail: nagai@imr.tohoku.ac.jp
URL: http://wani.imr.tohoku.ac.jp
Category: Energy-related Materials

IMR KINKEN Research Highlights 201126

Energy-related Materials

Ion-induced Luminescence and its Application to Analysis
of Dynamic Damage ProcessI

rradiation induced luminescence was used to investigate dynamic process of damage accumulationin dielectric materials. The evolution curves of the luminescence intensity during the H and He ion
bombardment were analyzed to examine the defect production and annihilation process causedby the nuclear collisions and electronic excitation.
Drastic changes in the structural andelectrical properties of dielectric materials
under irradiation are of great interest for theirapplication to fission and fusion devices [1].
Electronic excitation processes play
important roles for creating defects in ceramics,
known as color centers and excitons in ionic
crystals and oxides. Those radiation induced
phenomena such as optical absorption andradio-luminescencecanbepotentiallypowerful tools capable of examining
modification and its evolution under irradiation
condition. For example, in silica glasses, wehave been successfully investigated neutral
oxygen deficiency centers (ODCs), which are
not able to be detected by a probe sensitive toparamagnetic defects such as E’centers, theperoxy radicaland non-bridging oxygen hole
center (NBOHC), by using the ion-induced
luminescence methods [2-4]. An advantage ofthe ion-beam induced luminescence is thatconventional ion-beam analysis techniques
are simultaneously applicable to obtaincompositional depth distribution of lattice
atoms and displaced atoms up to several μm.
The combined techniques were used toinvestigate the interaction between hydrogenand irradiation induced damages in tritium
breeding ceramics [5, 6], and? tungsten
oxides[8-10]. Recently, we are trying to apply
the ion-induced luminescence to polymer

Fig.1: A Schematic illustration of an in-situ
measurement of ion–induced luminescence
combined with ion-beam analysis techniques.

27

IMR KINKEN Research Highlights 2011

films[10-12], to search a possibility of constructing
a tow-dimensional monitoring system for theionizing radiation.
In the experiment, the evolution of the ioninduced luminescence spectrum is obtained
under the irradiation with MeV energy light ionssuch as H and He ions. Simultaneously, backscattered and forward-recoiled particles fromsamples are detected to evaluate variation of
atomic composition in the near surface layer,
as schematically shown in Fig. 1. In silica glasses,
we observed three broad peaks of ion-induced
luminescence, originated from NBOHC,ODCII(B2α) and ODCII(B2β) centered at around
1.9, 2.7 and 3.1 eV, respectively. The B2βband
immediately diminished at a very small ion
fluence, and the B2α band started to increase
with successive irradiation, depending on its
OH content. The higher intensity of ODCII(B2α)
in the lower-OH silica may be explained by aspeculation that the glass network was morestrained in low-OH silica, where Si-O-Si bonds
are easily broken into oxygen deficiency
centers to stabilize the glass network. The 1
MeV H and He ion induced luminescence froma synthesized silica glass containing 800 wt.ppm
OH plotted as a function of incident ion fluence.
The ODCII(B2α) emission intensity increased
from zero with an increase of the ion fluence,and tended to saturate at higher fluence.
Assuming a first order reaction for the increaseof the luminescence intensity under the local
energy deposition along the ion trajectory, theformation and annihilation rates of theluminescence centers can be estimated fromthe evolution curves as shown in Fig. 2. The
production rates by H ion irradiation, however,are lower than the annihilation rates, indicatingthat the annihilation by H ion irradiation is more
prominent in comparison with He ion
irradiation. This result was reasonably explainedby the relation between the nuclear stoppingand electronic stopping for the different ionsand energies.
Ion bombardment of the polymer
materials, in particular, causes importantchemical transformations, depending on theenergy deposition mechanisms. An aromatic
polymer similar to polyethylene terephthalate(PET) is polyethylene naphthalate (PEN),
consisting of two benzene rings in a monomeric
structure. Both films exhibited characteristic

800

1 MeV Helium
I (t )

1000

R
0

dN ( x, t )dtNormalized IBIL intensity (arb. units)

K ( E ) N ( x, t )dx

pN 0

N ( x, t )1 MeV Proton

1000

440 nm from PEN

100

1 MeV proton

600
400

10

2.8 MeV Helium

1
0

100

200

300

12

400

2

500

Ion fluence (10 /cm )

600

200
0

500

200

150
100

0

1

10

100
16

50

1000
-2

Incident ion fluence (10 ions m )Fig. 2: Evolution of the luminescence intensity under
1 MeV H and He ionirradiation plotted as a functionof incident ion fluence. Inserted solid curves are
fitted by a model with parameters of production
and annihilation rate constants.

luminescence bands mainly in the blue region
and the luminescence intensity monotonicallydecreased with ion fluence. According to thesimultaneous
Fourier-transform
infraredabsorption measurements, the CH2 bending
vibration of the trans-configuration of theethylene glycol residue and typical vibration
bands of the para-substituted benzene rings
were assigned to the bands at 1471 and 1504
cm−1, respectively. The absorption due to thesebands gradually decreased with increasing
incident ion fluence, suggesting the loss of
crystallinity and the scission of the main chainat the para position. The ratio of the damaged
cross-sections of H to He is comparable for the
PET and PEN films, suggesting that the damage
processes of the luminescence centers in bothfilms are essentially the same. On the other
hand, the ratios of the stopping power of H to
He are considerably smaller in both films. The
luminescence centers along the H+ ion
trajectories effectively damaged, and an
excess of the energy was deposited aroundthe He+ ion trajectory for altering theluminescent structure. The back-scattering and
Forward-scattering analysis revealed that H, C,
and especially O were lost during the ion
irradiation. An Al coating of about 100 nm was
successfully applied to reduce the degradationof the PET and PEN films. The suppression of theloss of elements from the films indicated thatthe thin Al layer works as a barrier for volatile
molecules caused by local heating andcharging effects by incident ions.

300

400

500

600

700

0

Ti
m
e(
s)

Normalized 2.7eV luminescnece intensity(arbitrary units)

1500

Luminescence intensity

IMR KINKEN Research Highlights 2011

Wavelengths (nm)Fig. 3: Ion induced luminescence spectra from a
PEN film under 1MeV H ion bombardment plotted
as a function if irradiation time. The inserted graphshows the Evolution of the luminescence intensity
at 440 nm from the PEN film irradiated by1 MeV H+
and 2. 8 MeV He++ ions.

References[1] T. Shikama, K. Toh, S. Nagata, B. Tsuchiya, M. Yamauchi,
T. Nishitani, T. Suzuki, K. Okamoto, N. Kubo, Nuclear Fusion
46, 46(2006).
[2] S. Nagata S. Yamamoto, K. Toh, B. Tsuchiya, N. Ohtsu, T.
Shikama, H. Naramoto, J. Nucl. Mater. 329,1507(2004) .
[3]M. Watanabe, T. Yoshida, T. Tanabe, S. Muto, A, Inouye, S.
Nagata, Nucl. Instrum.and Meth.B 250, 174 (2006)[4] S. Nagata, H. Katsui, B. Tsuchiya, A. Inouye, S. Yamamoto,
K. Toh, T. Shikma, J. Nucl. Mater. 367, 1009 (2007).
[5] H. Katsui, S. Nagata, K. Toh, B. Tsuchiya, T. Shikama, J.
Nucl. Mater. 386, 1074 (2009).
[6] H. Katsui , S. Nagata, B. Tsuchiya, T. Shikama., Nucl.
Instrum. and Meth. B 268, 2735 (2010).
[7] S. Nagata, A. Inouye, S. Yamamoto, B. Tsuchiya, K.
Takano, K. Toh, T. Shikamaet., J. Alloys and Comp. 446,
558 (2007).
[8] A. Inouye, S. Yamamoto, S. Nagata, M. Yoshikawa, T.
Shikama, Nucl. Instr. and Meth. B 267, 1480 (2009).
[9] S. Nagata, H. Fujita, A. Inouye, S. Yamamoto, B. Tsuchiya,
T. Shikama., Nucl. Instrum. and Meth. B 268, 3099(2010)[10]S. Nagata, K. Takahiro, B. Tsuchiya, H. Katsui, T. Shikama,
Nucl. Instrum. and Meth .B 267, 1553 (2009).
[11] S. Nagata, H. Katsui, K. Takahiro, B. Tsuchiya, T. Shikama,
Nucl. Instrum. and Meth. B 268, 3099(2010)[12]M.Saito, F.Nishiyama, K.Kobayashi, S.Nagata, K.
Takahiro, Nucl. Instrum. and Meth. B 268, 2918(2010)

Keywords: luminescence, ion-solid interactions, nuclear material
T.Shikama S.Nagata and M.Zhao (Nuclear Materials Science)
E-mail:nagata@imr.tohoku.ac.jp
URL:http:// www-lab.imr.tohoku.ac.jp/~wshikama/

IMR KINKEN Research Highlights 201128

Energy-related Materials

Effects of Displacive Radiations and Hydrogenation in Zrand Fe-based Alloys Investigated by Ion-accelerator,
HVEM and Nano-hardnessTo develop materials for advanced light water reactors and fusion reactors, we investigated Zr- andFe-based alloys by means of ion irradiation, HVEM and nano-hardness. Hardening of Zr-Nb alloys
was evident at less than 1 dpa (displacement per atom) followed by saturation whose value depends
on composition. In Fe-alloy we have found that one-dimensional motion of defect clusters whichhave been noticed in high purity iron and FCC metals. Effect of helium upon irradiation hardening
was observed even at relatively lower doses in F82H.

(1) Effects of ion irradiation in Zr-Nb alloys
Zr-1.8Nb (J1), Zr-1.6Nb-0.1Cr (J2) and Zr2.5Nb (J3) (in mass %) were subjected to ion
irradiation,
followedbynano-hardness
measurement. Irradiations with 4 MeV Ni3+ ionsto the dose up to 11 dpa (displacement peratom) at 573K were performed at HIT facility,
University of Tokyo. Fig. 1 shows dose
dependence of irradiation hardening [1]. The
steep increase at the early stage of irradiation
followed by saturation roughly above 1 dpa

was evident. The level of saturation depended
on materials, and the effect of solution, Cr,
presumably resulted in reducing irradiation
hardening. In this study, we investigate effectof hydrogenation in the alloys which will be
described elsewhere [2].
(2) Motion of defect clusters in Fe-alloy
Recent theoretical and experimental
studies have confirmed that small interstitialclusters in pure metals cause a gliding motion
along the direction of the Burgers vector withlow activation energy, called one-dimensional(1D) migration [3]. This phenomenon has beensuggested to affect defect structuraldevelopment under irradiation with highenergy particles, such as void swelling. Although
1D migration processes in practical alloys have

SUS316L

10-1
1D migration frequency (s-1)

The mechanical degradations of nuclear
materials
strongly
dependon
theirenvironments. For instance, Zr-alloys forcladding in LWRs are affected by distribution of
hydrides as well as irradiation defects, and Fealloy for fusion blanket is done both by radiations
and helium accumulation. Therefore, it isindispensable to understand the microstructural
evolutions
andresultingmechanicaldegradations and to clarify their correlations.
We investigate hardening induced by ion
irradiation and/or He/H implantation in Zr- andFe-based alloys by nano-hardness tester,electron microscopy, as well as fundamental
study on stability of defect clusters underelectron irradiation in HVEM.

Fe-Cr-Ni
+C
+Mn

Fe

10-2

10-3
1024
1025
electron beam intensity (e m-2 s-1)
Fig. 1 Irradiation hardening in Zr-Nb alloys underirradiation with 4MeV Ni3+ ions at 573K.

29

IMR KINKEN Research Highlights 2011

Fig. 2. Frequency of 1D migration of interstitial clusters
in SUS316L and its model alloys as a function ofelectron beam intensity.

IMR KINKEN Research Highlights 2011technical and fundamental significance fordeveloping materials for nuclear power
applications, they are not well understood. In
the present study, 1D migration of interstitialclusters in austenitic stainless steel SUS316L andits model alloys, namely, Fe-18Cr-13Ni, Fe-18Cr13Ni-0.012C, and Fe-18Cr-13Ni-1.7Mn, wasexamined using in-situ observation by highvoltage electron microscopy. As shown in Fig.
2, such 1D migration was confirmed to occur
along the <110> direction at irregular intervals
in all these alloys under 1250 kV electron
irradiation at room temperature. The frequency
of 1D migration was found proportional toelectron beam intensity, and was about 1/10
that in high-purity iron under the same
irradiation intensity.
The distance of 1D
migration in the four alloys was less than 10 nm,
which was much shorter than that in high-purity
iron. No clear difference in the frequency or
distance of 1D migration was observed among
the four alloys, suggesting that minor solute/impurity elements have no apparent effect on
1D migration in SUS316L.
(3) Irradiation and He-induced hardening in
RAFM steel
Reduced activation ferritic/martensitic(RAFM) steel has been recognized as a primarynear-term candidate material for nuclear
fusion reactors [4]. In the fusion reactor
environment, irradiation of 14 MeV neutrons
might produce a large amount of displacement
damage and transmutant helium (He) atoms instructural materials. For example, 100 dpa of
displacement damage and 1000 appm ofhelium concentration will be raised to RAFM
steel after 10 MW/m2 of neutron wall loading[5].
In the previous research [5,6], ion- irradiatedF82H steel to high He/dpa at 633 K caused theradiation hardening accompanied by theformation of black dot defects, which were
probably dislocation loops, radiation induced
precipitates and nano-voids. Under neutron
irradiation, the radiation hardening caused byirradiation- induced dislocation loops andprecipitates changes occurred at irradiation
temperatures up to 693 K.
The present work investigated theirradiation hardening behaviors and themechanisms in single and dual ion irradiatedF82H steel by combining multi-ion irradiation,
ultra-micro-indentation
techniqueandmicrostructural evolution. F82H steel (IEA-heat)Fig. 3 Irradiation hardening behavior by single anddual ion irradiation at 573K for F82H steel.

was irradiated in various conditions with single/
dual ion beam at the HIT facility. For the single
ion irradiation experimental, 4 MeV Ni3+, 2.8
MeV Fe2+ and 1 MeV O+ ions were induced at
573 K, and 2.8 MeV Fe2+ and 1 MeV He+ ionswere induced for dual ion irradiation test at thesame conditions. And nano-hardness test was
performed
followedbymicrostructuralobservation in order to evaluate and identify
the irradiation hardening behavior as shown in
Fig. 3.
References[1] Y. Satoh et al, Philos. Mag. 89 (2009) 1489.
[2] K. Shiba et al., J. Nucl. Mater. 329–333 (2004) 243.
[3] M. Ando et al., J. Nucl. Mater. 329–333 (2004) 1137.
[4] H. Ogiwara et al., J. Nucl. Mater. 367-370 (2007) 428.

Keywords: ion irradiation, electron irradiation, mechanical properties, electron microscopy
Hiroaki ABE (Nuclear Materials Engineering)E-mail: abe.hiroaki@imr.tohoku.ac.jp
URL: http://www-lab.imr.tohoku.ac.jp/~abeken/

IMR KINKEN Research Highlights 201130

Energy-related Materials

New Approach for Designing Mg Based BMG Matrix
Composite with in-situ α-Ti Dispersoids

A

Mg61Cu28Gd11 BMG matrix composite with in-situ α-Ti dispersoids was successfully synthesized bymixing a prealloy comprising Ti2Cu and Cu-Gd compounds and pure Mg melt followed by coppermold casting. The Cu-Gd phases along with just the Cu from the Ti2Cu phase were selectively
dissolved into the Mg melt to form Mg61Cu28Gd11 bulk metallic glass former, and the remaining Ti
formed porous structures that were dispersed in the melt. The size of the in-situ α-Ti dispersoids is
directly related to that of the Ti2Cu in the prealloy, and is thus controllable by the cooling rate usedin the prealloy preparation. The maximum fracture stress in four-point bending mode was increased
from 230 MPa to 387 MPa upon refining the α-Ti dispersoids. The relationship for improving tensile
properties, L ≈ Rp, where L is the size of the porous α-Ti and Rp is the size of the process zone, wasachieved.
Mg-based bulk metallic glass (BMG) is apromising structural material because of its
high specific strength (~2.3×105 Nm/kg) despite
the low fracture toughness (a few MPa m1/2) as
low as that of brittle ceramics [1]. To amend thislow toughness, many efforts have been madeto develop Mg-based BMG matrix composites(BMGMC) [2-5]. The mechanical property
assessment of the composites reported in thoseworks, however, was mostly performed undercompressive mode instead of the tensile one.
A composite with in-situ precipitates
bonded strongly with the matrix has become aparticular interest because they have clean
oxide-free surfaces. In addition, theseprecipitates tend to have higher aspect-ratio
shapes, i.e. dendritic, ellipsoidal or flaky, which
are effective for bridging any separation of
BMG at the shear planes under tension.
Accordingly, we developed Mg-based BMG
composite containing the in-situ precipitationof α-Ti to improve the tensile property. The alloy
design was based on the formula which
effectively enhances tensile plasticity of the
BMGMC [6], i.e. L ≈ S ≈ Rp, where L is the dendrite
size, S is the inter-dendrite spacing and Rp is thelength scales of the process zone. Because of
the small Rp may lead to the very low toughnessof Mg-based BMG, the microstructuralrefinement down to the optimal size of 0.1 to1.0
μm is necessary. The similar in-situ Mg-based
BMGMC was also developed previously by Li etal. [5], but the microstructural length scale oftheir composite was much larger (~40 μm)
compared to the optimal one. In the present
work the microstructure of the BMGMC was
adjusted to L ≈ S ≈ Rp by controlling the in-situ
formation process of α-Ti.
For inducing the in-situ formation of α-Ti in
the Mg-Cu-Gd, the alloy was prepared by twostep procedures designed based on theenthalpy of mixing ( ΔH{mixAB}) which reflects thechemical interaction among atoms. In the firststep, by pre-alloying Cu, Gd and Ti by arcmelting method, one can expect the formationof Ti-Cu and Gd-Cu phases because of the31

IMR KINKEN Research Highlights 2011

Fig. 1: SEM image and corresponding elemental
maps of the mother alloy using slow-cooledprealloy for Ti (green), Mg (red) and Cu (yellow).

negative enthalpy of mixing for Ti-Cu and GdCu and the positive one for Ti-Gd [7]. The resultof the cross sectional analysis of the pre-alloyedingot shows that the acicular and circularprecipitates of Ti2Cu were surrounded by
Cu7Gd2 and Cu2Gd (at.%), which is in goodaccordance with our prediction [8]. The second
step is preparing the mother alloy by mixing
the pre-alloy with pure Mg. In this step, the insitu formation of pure Ti from the Ti-Cu phase is
expected due to the negative enthalpy of
mixing of Mg-Cu and Mg-Gd and the positive
one for Mg-Ti [7]. Fig. 1 shows the SEM image
and the corresponding elemental mapping of
the cross section of the mother alloy [8]. Cu
element (yellow) was not observed in the Ti richarea (green), but in the same areas (red) as theMg rich matrix. This clearly confirms that Cu
separated from the Ti2Cu precipitates observedin the pre-alloy to form the Mg-Cu-Gd matrix
with α-Ti dispersoids. An Mg-Cu-Gd BMGMC
with in-situ α-Ti precipitates was then successfully
formed by casting the mother alloy into a Cumold. The XRD pattern of the as-cast sample
confirms the weak peaks of the α-Ti phase inIMR KINKEN Research Highlights 2011

addition to the typical broad peak of the glassy
structure.
It was also observed that the α-Ti dispersedin the glass matrix of the Mg-Cu-Gd BMGMC
has porous plate-like structure and the size wellcorresponds to the plate-like Ti2Cu in the prealloy. Accordingly, microstructural control ofthe pre-alloy may lead to the formation of finer
Ti2Cu, which in turn results in the in-situ formationof finer α-Ti dispersoids in the BMGMC. We thenemployed arc-melt type mold casting method
during pre-alloying which enables the prealloy ingot to solidify at a much faster coolingrate. (Hereafter, the pre-alloy prepared by arcmelt type mold casting is called as rapidlycooled (RC) pre-alloy, while that produced byslow solidification on a the copper hearth iscalled as slow-cooled (SC) pre-alloy).
Fig. 2 shows SEM image of an α-Ti dispersoidof the BMGMC prepared from the RC pre-alloy
after the immersion in aqueous HNO3. The
dipersoid has a porous plate-like structure witha pore-size of ~500 nm, width of ~4 μm andthickness of ~500 nm [8]. It was much finer thanthat of the BMGMC prepared from the SC prealloy. The effective design for improving thetoughness of the BMGMC was partiallyachieved, i.e. by controlling L ≈ Rp.

Fig. 3: Stress-Strain curves of the BMGMCs preparedusing rapidly-cooled prealloy (RC prealloy) andslow-cooled prealloy (SC pre-alloy) under the
four-point bending test. The curve of the monolithic
Mg61Cu28Gd11. BMG is inserted for comparison.

The size-refinement of the porous α-Ti in the
BMGMC results in the increase of the maximum
fracture stress under the four-point bending
mode from 230 MPa (the SC pre-alloy BMGMC)
to 387 MPa (the RC pre-alloy BMGMC) as shownin Fig. 3.
Moreover, the localized plastic
deformation via continuous slips of multiple
shear bands was also observed in the tensileside surface [8].
References

Fig. 2: a-Ti dispersoids obtained from the BMGMC
made using the RC prealloy filtered after immersion
in aqueous HNO3.

[1] M.F. Ashby, A.L. Greer., Scripta. Mater. 54, 321 (2006).
[2] J.S.-C. Jang, S.-R. Jian, T.-H. Li, J.C. Huang, C.Y.-A. Tsaoand C.-T. Liu., J. Alloys Compd. 485, 290 (2009).
[3] M. Kinaka, H. Kato, M. Hasegawa, A. Inoue., Mater. Sci.
Eng. A. 494, 299 (2008).
[4] H. Ma, J. Xu and E. Ma., Appl. Phys. Lett. 83, 2793 (2003).
[5] F. Li, S. Guan, B. Shen, A. Makino and A. Inoue., Mater.
Trans., JIM 48, 3193 (2007).
[6] D.C. Hofmann, J-Y. Suh, A. Wiest, G. Duan, M-L. Lind,
M.D. Demetriou and W.L. Johnson., Nature 451, 1085(2008).
[7] A. Takeuchi, A. Inoue., Mater. Trans., JIM 46, 2817 (2005).
[8] H. Oka, T. Wada, A. Inoue and H. Kato., J. Mater. Res .
under review.

Keywords: composite, mechanical properties, metallic glass
Hideki Oka (Non-Equilibrium Materials)
E-mail: hideki-o@imr.tohoku.ac.jp
Category: Energy-related Materials

IMR KINKEN Research Highlights 201132

Energy-related Materials

Growth Mechanism of Si-faceted Dendrites

T

he growth mechanism of faceted dendrites has not been clarified since the discovery of the
faceted dendrite of germanium in the 1950s. We successfully observed the growth processes of Sifaceted dendrites by in situ observation technique. The growth model of the faceted dendrites isestablished based on our experimental results.
Dendrites of faceted materials, so-called‘faceted dendrites’, which are distinguishedfrom the dendrites of nonfaceted materials in
metals and alloys, were discovered in the 1950s[1]. Recently, Si-faceted dendrites haveattracted much attention from both scientificand technological viewpoints due to theirunique crystal structures. The surface of the
faceted dendrite is bounded by {111} habit
planes, and at least two parallel twins exist at
the center of the faceted dendrite. It is also
known that the preferential growth direction of
faceted dendrites is <112> or <110>. Suchfeatures can be applied in technologies for
growing thin Si ribbon crystals and polycrystalline
Si ingots for solar cells.
The growth model of faceted dendrites
preferentially grown in the <112> direction,
hereafter referred to as <112> dendrites, wasfirst proposed in 1960 by Hamilton and
Seidensticker [2]. The model had been widely
accepted although the actual growth behavior
had remained mystery because of the lack of
direct evidence. We recently succeeded inobserving the growth behavior of <112>
dendrites by in situ observations, and proposeda growth model based on experimental resultsof the growth behaviors of <112> dendrites [3].
On the other hand, no growth model of <110>dendrite has been proposed.
We succeeded in directly observing the
growth processes of <110> dendrites by using
an in situ observation system, which consistedof a furnace and a microscope [4].
Fig. 1(a) shows a <110> dendrite and a<112> dendrite growing from a part of a faceted
crystal-melt interface. The shape of the tip of
the growing dendrite is markedly different
between the <110> and <112> dendrites.
Although the tip of the <112> dendrite becomeswider during growth, that of the <110> dendrite
remains narrow during growth.
We presented a growth model of <110>dendrites based on our experimental evidence,as shown in Fig. 1(b). The first illustration in Fig.
1(b) shows the equilibrium form of the Si crystal
with two parallel {111} twin boundaries. The
crystal is bounded by {111} habit planes. It isfound that reentrant corners with an externalangle of 141° (type Ⅰ) exist at both twins. Rapid
growth occurs at these type Ⅰ corners, andleads to the formation of triangular corners.
Although the rapid growth is inhibited after the33

IMR KINKEN Research Highlights 2011

Fig. 1: (a) Growth shape of Si <110> and <112>
dendrites. (b) Growth model of the <110> dendrite[4].

formation of triangular corners, the crystal
continuously grows on the {111} flat surface.
When the triangular crystals propagated acrossthe other twin, two type Ⅰ corners appear,
and rapid growth occurs there again. The later
growth processes are the same as described
before. Note that the tip of the dendrite remains
narrow during dendrite growth. Thus, we can
explain the shape of the <110> dendrite during
growth, shown in Fig. 1(a), by this mechanism.
We have developed an equation for thetheoretical growth velocity of a Si <110> faceted
dendrite on the basis of the growth model
shown in Fig. 1(b) [5]. The theoretical growthvelocity is described asV <110 > =h+V 2 ,
3 (h/V1 + d/V2)(1)

IMR KINKEN Research Highlights 2011

where h is the height of the triangular corner, dis the twin spacing, V1 is the growth velocity ata type Ⅰ corner, and V2 is the growth velocity
on the {111} plane. It is found that the growthvelocity of the dendrite is inversely proportional
to twin spacing.
Fig. 2 (a) shows the growth of three dendrites(d1, d2, d3) from the same crystal-melt interface
at the same time. The preferential growth
direction of these three dendrites is <110>. It isfound that the growth velocities are obviously
different among the three dendrites. Fig. 2(b)shows the parallel twins observed at the centerof d1, d2, and d3 by EBSP measurement. The
average twin spacings in d1, d2, and d3 are 4.5
μm, 10 μm, and 16.5 μm, respectively. Dendrite
growth velocity as a function of twin spacing isplotted in Fig. 2 (c). As twin spacing decreases,
dendrite growth velocity nonlinearly increases.
We use Eq. (1) to fit the experimental results, as
shown by the red line in Fig. 2(c). Equation (1)fits quite well with the experimental results. The
fitting results of h, V1, and V2 are 45±5 μm,
4600±500 μm/s, and 120±5 μm/s, respectively. It
is found that the growth velocity at the type Ⅰ
corner is much faster than that on {111} plane. It
is also found that the growth velocity of adendrite eventually approaches a certain
value as twin spacing increases to infinity.
We also investigated the dependence ofthe growth velocity of faceted dendrites onundercooling [6]. The growth velocity increased

linearly with increasing undercooling. We foundthat the relationship between the growthvelocity and undercooling is most sensitive totwin spacing.
In summary, we investigated the growthmechanism of Si-faceted dendrites. The growthmodel of the Si-faceted dendrite wasestablished based on the experimental results
by in situ observations. Our in situ observation
technique is a powerful tool for revealing the
crystal growth mechanism.
References[1] E. Billig, Proc. R. Soc. London Ser. A 229 (1955) 346.
[2] D. R. Hamilton and R. G. Seidensticker, J. Appl. Phys. 31(1960) 1165.
[3] K. Fujiwara et al., Phys. Rev. Lett. 101 (2008) 055503.
[4] K. Fujiwara et al., Phys. Rev. B 81 (2010) 224106.
[5] X. Yang et al., Appl. Phys. Lett. 97 (2010) 172104.
[6] X. Yang et al., Appl. Phys. Lett. 98 (2011) 012113.

Fig. 2: (a) Growth behavior of faceted dendrites. Three dendrites grew from the same crystal-melt interface [5].
(b) Parallel twins observed at the center of the three dendrites [5]. (c)Dendrite growth velocity as a functionof twin spacing [5].

Keywords: crystal growth
K. Fujiwara (Crystal Chemistry Division)
E-mail: kozo@imr.tohoku.ac.jp
URL:http://www.uda-lab.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201134

Energy-related Materials

High-speed Preparation of YBCO Superconducting Film byLaser Chemical Vapor DepositionIn high-temperature superconducting wire, superconducting layer (YBa2Cu3O7-δ; YBCO) has
generally been prepared by pulsed laser deposition or chemical vapor deposition (CVD). However,their low deposition rate and low yield constant are issues in practical production process. Chemicalcomposition and orientation of the superconducting layer should also be precisely controlled inorder to archive a high superconducting performance. We have developed a new laser CVD andsuccessfully prepared a-axis oriented CeO2 film for buffer layer and c-axis oriented YBCO film forsuperconducting layer on multilayer-coated Hastelloy tape substrate at high deposition rate.

IMR KINKEN Research Highlights 2011

]
[0

01

500 nm]35

10Preparation of CeO2 film
a-axis oriented CeO2 films epitaxially grownon (100) SrTiO3 single crystal substrate at thewhole laser power from 52 to 182 W (deposition

a

[0

b
(111)
(111)

(111)

High-temperature superconducting (HTS)materials having a critical temperature (TC)above 77 K (boiling point of liquid nitrogen)opens up a new horizon for applying
superconductivity to electrical devices. One of
the most promising HTS materials is YBa2Cu3O7-δ
(YBCO). The TC of YBCO is approximately 90 K
and its high critical current density (typicallyover 1 MA/cm2) is the best to use for powerdevices.
In HTS practical application including
power cables, generators, motors and sensing
devices, a flexible superconducting wire with along length up to several kilometers is required.
In order to meet this requirement, Ni-Cr alloy(Hastelly® C-276) tape is used as a substrate.
On the tape substrate, oxide buffer layers, suchas Gd2 Zr 2O7, MgO, LaMnO3 and CeO2 are
generally deposited by physical vapor
deposition. These buffer layers can prevent theoxidization of the metal substrate at hightemperature, reduce lattice mismatches andthe thermal coefficient mismatches between
YBCO superconducting layer and tapesubstrate, and provide an in-plane epitaxial
growth of c-axis oriented YBCO layer.
Various deposition techniques, such as solgel,
metalorganic
deposition
using
trifluoroacetates (TFA-MOD), pulsed laser
deposition(PLD),
sputtering,thermalevaporation, molecular beam evaporation(MBE) and metalorganic chemical vapor
deposition (MOCVD) has been employed for
preparation of CeO2 buffer layer and YBCO
superconducting layers. Main issues for above
techniques are relatively low deposition rates(below 1 μm/h). In practical application, highspeed preparation of YBCO and CeO2 films isrequired.
We have developed new coating Laser
chemical vapor deposition (laser CVD) [1,2],
which could be an effective approach toprepare highly oriented YBCO film and its buffer
layer CeO2 film at high deposition rate.

(111)

Ce
Fig. 1: (a) Surface microstructure, (b) schematic ofin-plane orientation and (c) atomic arrangementof CeO2 film prepared on (100) SrTiO3 substrate bylaser CVD.

temperatures of 952–1119 K). As for the (100)CeO2 films on (100) SrTiO3 substrate, their XRD
pole figure patterns showed fourfold reflections
at α = 45°, attributing to the {220} planes. The
pole figure pattern from SrTiO3 (110) showed afourfold reflection at α = 45°, attributing to the{110} planes. The azimuth angle β of the fourpole peaks from the CeO2 film showed a 45°
rotation on the basal plane of SrTiO3 substrate,
implying the in-plane epitaxy of (100) CeO2 //(100) SrTiO3 and CeO2 [010] // SrTiO3 [011].
Microstructure of (100) CeO2 films preparedon (100) SrTiO3 substrate consisted ofrectangular-shaped grains (Fig. 1(a)) andshowed a columnar growth in the cross section.

IMR KINKEN Research Highlights 2011

The pyramidal caps of the rectangular-shaped
grains were attributed to CeO2 {111} planes (Fig.
1(b)). The {111} planes had lowest surface energy(9.6 eV/nm2) among the low index planes of(110), (211), (100), (210) and (310) which had
surface energy of 15.3, 16.7, 20.3, 20.4 and 22.4
eV/nm2, respectively. CeO2 {111} planes were
therefore energetically favorable to be formedat the late deposition stage of CeO2 film.
Preparation of YBCO film
Fig. 2(a) shows the YBCO films prepared on
multilayered Hastelloy C-276 tape with (100)CeO2 top layer. c-axis oriented YBCO film was
epitaxially grown on (100) CeO2 top layer (Fig.
2(b)). The electrical resistivity of c-axis-oriented
YBCO films decreased linearly with lowering
the temperature in normal state and showedsharp zero resistance transitions at TC = 90 K (Fig.
2(c)).
The XRD pole figure from the (102) reflectionof the c-axis-oriented YBCO film showed afourfold pattern at around α = 33°, which was
attributed
to{102}
planeswith
thecomplementary angle of 57° to the (001) plane(Fig. 2(d)). The azimuth angles β of the polepeaks from (102) reflection of the c-axisoriented YBCO film rotated 45° with respect tothose of CeO2 (220) reflection (Fig. 2(d)),
indicating the in-plane alignment growth of
YBCO (001) // CeO2 (100) and YBCO [100] //CeO2 [011] (Fig. 2(e)). a-axis- and c-axis-oriented
YBCO films were also obtained by changing
laser power. The preferred orientation changed
from a-axis to c-axis with increasing laser
power.
The deposition rate of c-axis oriented YBCO
films was 83 μm/h, and this was about 3-900
times higher than that of YBCO by conventional
MOCVD [3].

Fig. 2: (a) Cross-sectional microstructure of c-axis
oriented YBCO film prepared on (100) CeO2 top
layer of multilayer-coated Hastelloy tape, (b)selected area electron diffraction patterns, (c)electrical resistivity against temperature, (d) pole
figure patterns and (e) atomic arrangement of
YBCO film on CeO2 (100). TEM observation courtesyof JFCC.

References[1] T. Goto et al., Thin Solid Films 515 (2006) 46.
[2] A. Ito et al., J. Alloys Compd. 489 (2010) 469.
[3] P. Zhao et al., Supercond. Sci. Technol. 23 (2010) 125010.

Keywords: chemical vapor deposition, ceramic, superconducting
T. Goto, R. Tu and A. Ito (Multi-Functional Materials Science Division)
E-mail: goto@imr.tohoku.ac.jp
URL:http://www.goto.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201136

Energy-related Materials

Production of Ni-Cr-P-B Alloy-Coated Bipolar Plates for
PEMFC by HVOF Spray-Coating and Their Surface Analysis
by XPSI

n this study, the newly designed bipolar plates for proton exchange membrane fuel cells (PEMFC)
were produced by the High Velocity Oxy-Fuel(HVOF) spray-coating the Ni65Cr15P16B4, Ni60 Cr 20 P16B4
and Ni55Cr 25P16B4 alloys on Al plates having a flow field. The I-V performance of a single fuel cell withthese bipolar plates produced in this work was studied and we found that the single fuel cells withthe alloy-coated bipolar plates showed high I-V performance as well as the cell with the carbon
bipolar plates. After that, the 24h durability tests were conducted. As the result, the cell voltage of asingle fuel cell with the alloy-coated bipolar plates did not show significant voltage drop during thetests. Furthermore, XPS analysis of the surface layer of the Ni-Cr-P-B alloy-coated bipolar plates was
conducted after the 24h durability tests. It was found that the surface passive film of Ni-Cr-P-B alloy
contains the Cr 2O3 of very high concentration, resulting in high I-V performance and good durability.
Bipolar plates are one of the most importantcomponents of PEM fuel cells and are
multifunctional as they conduct electricity fromcell to cell, they separate the fuel gas from theoxidant gas, and their flow field supplies thegases to the electrodes. Bipolar plates areconventionally made of carbon graphite.
However, high production cost of graphite
bipolar plate inhibits wide spread use of thefuel cells. Moreover, it cannot be used to makea thinner bipolar plate because of its significant
brittleness.
Compared to the carbon graphite, themetallic materials have many advantages
such as low-cost, good electrical and thermal
conductivities,
excellentmechanicalproperties and good workability. However,corrosion may occur when the metallic bipolar
plates are used under acidic atmosphere in
the fuel cells, which leads to the dissolution of
metallic ions.
So, we focused on metallic glass which is
known for its extremely high corrosion resistanceand tentatively produced the Ni-Cr-P-B metallic
glass-coated bipolar plates by using the High
Velocity Oxy-Fuel (HVOF) spray coating
technique in this study. The electricity
generating properties of the single fuel cellwith the metallic glass-coated bipolar plates
were also examined [1,2]
Production of metallic glass-coated bipolar
platesThe bipolar plates were produced by the
HVOF
spray-coating
theNi65Cr15P16B4,
Ni60 Cr 20 P16B4 and Ni55Cr 25P16B4 alloys on Al
plates having a flow field.
Fig. 1 shows the outer view of theNi60 Cr 20 P16B4 metallic glass-coated Al bipolar
plate produced by HVOF spray-coating andsubsequent hot-pressing. The glass film wasdeposited on both sides of the plates. As aresult of XRD observations, it was found that theand
HVOF
spray-coated
Ni65Cr15P16B4
Ni60 Cr 20 P16B4 alloy films deposited on Al plates
showed mainly a broad hallo peak comingfrom the glassy matrix and a small peaks fromthe crystalline phases. The spray-coated
Ni55Cr 25P16B4 alloy showed sharp distinct peaks
coming from crystalline phases.
Power generation tests with metallic glasscoated bipolar plates
Fig. 2 shows the I-V performance of a single
fuel cell with these bipolar plates produced inthis work. It was found that the single fuel cells

Fig. 1 : Schematic illustration of the flow field.

37

IMR KINKEN Research Highlights 2011

IMR KINKEN Research Highlights 2011

the surface layer of the Ni-Cr-P-B alloy-coatedbipolar plates after the 24h durability tests. As aresult, the Cr 2O3 passive film was found in all
the Ni-Cr-P-B alloy surface films and also the
P2O5 in the Ni65Cr15P16B4 and Ni55Cr 25P16B4 alloy
surface films and Ni(OH)2 in the Ni65Cr15P16B4
alloy surface film. This observation indicatesthat the surface passive film of Ni60 Cr 20 P16B4
alloy contains the Cr 2O3 of the highest
concentration among the three alloy surface
films, resulting in the highest I-V performanceand good durability.
Fig. 2 : I-V curves measured with a single fuel cellwith carbon, Al, Ni65Cr15P16B4 35), Ni60 Cr 20 P16B4 andNi55Cr 25P16B4 alloy-coated bipolar plates measured
after 50 cycles repetition.

Fig. 3 : 24h durability tests conducted at a constantcurrent density of 200 mA•cm -2 with a single fuelcell with carbon, Al, Ni65Cr15P16B4 35), Ni60 Cr 20 P16B4
and Ni55Cr 25P16B4 alloy-coated bipolar plates.

with the alloy-coated bipolar plates showed
high I-V performance as well as the cell withthe carbon bipolar plates. Among them, theNi60 Cr 20 P16B4 alloy showed the highest I-V
performance, showing the largest current
density at 0.3 V.
Fig. 3 shows the 24h durability tests
conducted at the constant current density of
200 mA•cm -2. The cell voltage of a single fuelcell with the alloy-coated bipolar plates did
not show significant voltage drop during thetests. It can be concluded that the most suitablealloy in this study is Ni60 Cr 20 P16B4 alloy.
XPS Analysis of metallic glass-coated bipolar
platesFig. 4 shows the results of the XPS analysis ofFig. 4 : XPS Ni 2p, Cr 2p and P 2p narrow region scansof the Ni65Cr15P16B4, Ni60 Cr 20 P16B4 and Ni55Cr 25P16B4
alloy.

References[1] S.C.Kim, S.Yamaura, Y.Shimizu, K.Nakashima, T.Igarashi,
A.Makino and A.Inoue, Mater. Trans. 51, 1609 (2010).
[2] S.C.Kim, S.Yamaura, T.Igarashi, Y.Shimizu, K.Nakashima,
A.Makino and A.Inoue, J. Japan Inst. Metals 75, 122(2011).

Keywords: fuel Cell, metallic glass, coating
S. Yamaura (Metallic Glass NEDO Laboratory)
E-mail: yamaura@imr.tohoku.ac.jp
URL: http://nedo-koza.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201138

Energy-related Materials

Point Defect Processes before Steady-state Void Growth inAustenitic Stainless Steels

Defect structures during the incubation period in austenitic stainless steels, both commercial alloys
such as SUS316SS and model alloys such as Fe-Cr-Ni, were studied using positron annihilation lifetime
measurements. By fission neutron irradiation to 0.2 dpa at 363 K, single vacancies were formed in thecommercial alloys and large voids were formed in the model alloys. After irradiation at 573 K, stacking
fault tetrahedra and/or precipitates were detected in the commercial alloys, while large voids wereformed in the model alloys. Defect evolution during the incubation period was discussed on thebasis of neutron irradiation results.
Nuclear materials are vulnerable to void
swelling during prolonged neutron irradiation.
Before swelling, there exists an incubation doseas shown in Fig. 1. After the dose, a steady statevoid swelling, 1%/dpa, starts in both model
alloys and commercial alloys of austenitic
stainless steels. The dose determines theduration of service under neutron irradiation in
nuclear reactors. The start of void swellingstrongly depends on alloying elements. In Ni,
voids were observed by transmission electron
microscopy (TEM) after fusion neutron irradiationof 10 -3 dpa at 573 K [1]. Ni is the simplest model
metal of austenitic stainless steels from theviewpoint of crystal structure and lattice
constant. In a model alloy of austenitic stainlesssteels, Fe-Cr-Ni ternary alloy, voids wereobserved at a dose of 0.1 dpa at 573 K [2], andin a commercial alloy of Ti added modified
SUS316SS, no voids were observed at a dose of
15 dpa at 701 K [3].

Fig. 1: Schematic illustration of incubation dose forvoid swelling in austenitic stainless steels.

As most experimental studies have been
performed after void formation using TEM,
point defects and their clusters smaller than
the resolution limit were impossible to detect.
Positron annihilation spectroscopy makes it
possible to detect small defects, which are
below the resolution limit of TEM. In this study,
point defect processes occurring during the
incubation period of austenitic stainless steels
39

IMR KINKEN Research Highlights 2011were studied using positron annihilation lifetime
measurements.
Specimens used were 99.99 % pure Ni, fourmodel alloys, specimen ID A: Fe-16.1Cr -17.0Ni,
B: Fe-15.4Cr-15.9Ni-2.68Mo-1.89Mn, C: Fe-15.3Cr15.8N-2.66Mo-1.88Mn-0.53Si and D: Fe-15.3Cr15.8N-2.66Mo-1.88Mn-0.53Si-0.24Ti, and fourcommercial alloys, J: Ti added modified
SUS316SS(Fe-15.27Cr-15.8Ni-2.66Mo-1.88Mn0.53Si-0.24Ti-0.055C-0.024P), E: SUS316L SS, F:
SUS316SS and G: SUS304SS. Positron annihilation
lifetimes of the specimens were measured atroom temperature using a conventional fastfast spectrometer with a time resolution of 190
ps (full width at half maximum) and each
spectrum was accumulated to a total of 1×106
counts. The positron lifetime spectra wereanalyzed using the PALSfit program [4] and thelifetimes were decomposed into two (τ1, τ2) or
three components (τ1, τ2, τ3) if possible.
Otherwise only an average lifetime (τav) was
obtained.
Positron annihilation lifetimes of Fe-14Cr13Ni were almost the same as those of Ni [5-7]and we used the lifetime of Ni for theidentification and the estimation of size of
defect clusters such as voids and stacking fault
tetrahedral in austenitic stainless steels. The
lifetimes of them are as follows. Single vacancy:
176 ps, di-vacancies: 195 ps, voids of 79
vacancies: 427 ps, stacking fault tetrahedra of
21 vacancies: 155 ps, stacking fault tetrahedraof 28 vacancies: 130 ps.
Before neutron irradiation, the positronannihilation lifetimes were between 106 and109 ps. After neutron irradiation using fission
neutrons from BR-2 of SCK•CEN at 363 K and
573 K to a dose of 0.2 dpa, the lifetimes changed
as shown in Tables 1 and 2. Dashes in thesetables indicate that two- or three- component
analysis of the positron annihilation lifetime
spectra was impossible.
The long lifetime of positron annihilation
depends on a size of vacant site where positrons
were annihilated. The longer and shorter
lifetime than that of single vacancies (176 ps[6]) suggests the existence of larger vacancy
clusters (voids) and smaller spaces than single
vacancies, respectively. Following 363 K

IMR KINKEN Research Highlights 2011

Table 1: Positron annihilation lifetimes of model alloys
and commercial alloys of austenitic stainless steels
irradiated at 363 K to 0.2 dpa.

ID
Ni
A
B
C
D
J
F
G

τav(ps)
271±1
183±1
176±1
177±1
176±1
154±1
181±1
185±1

τ1(ps)
161±2
169±1
135±8
128±8
114±9
50±3
-

τ2(ps)
403±4
486±28
195±4
193±3
187±2
174±1
-

I2(%)
46±1
4±1
67±9
73±7
82±4
84±1
-

irradiation as listed in Table 1, the lifetimes of Ni,
A, B, C, D and J were decomposed into two
components. Large voids were formed in Ni
and the ternary model alloy (A), and only single
vacancies were detected in model alloys B, C,
and D, and in the commercial alloys J, F andG.
At the 573 K irradiation, the mean lifetimes
were lower than those at 363 K as listed in Table
2. The lifetime τ3 of Ni and of the four model
alloys indicated the formation of voids. The
lifetimes in the commercial alloys were lower
than those of single vacancies, and the lifetime
corresponds to that of stacking fault tetrahedra[6] and grain boundaries. Therefore at low doseof 0.2 dpa, voids were not formed in thecommercial alloys but other defects wereformed.
As metal carbides which are commonly
formed by irradiation of austenitic stainlesssteels have incoherent grain boundaries withthe matrix, vacant spaces smaller than single
vacancies were formed. By TEM observation,
the formation of metal carbides in austenitic
stainless steels was observed only at higher
temperatures and higher irradiation doses thanwere used in the present experiment [8].
However, there is a strong possibility of theformation of very small precipitates at low
temperatures and low doses, and these canbe detected by positron annihilation lifetime
measurements.
363 K neutron irradiation formed large
voids in Ni and A, while only single vacancies
and di-vacancies were formed in other alloys.
Although voids were formed in Ni and A, theintensities l2, which present the ratio of positrons
annihilated at voids, indicated that the numberof voids in A was very small compared to Ni.

Table 2: Positron annihilation lifetimes of model alloys
and commercial alloys of austenitic stainless steels
irradiated at 573 K to 0.2 dpa.

ID
Ni
A
B
C
D
J
E
F
G

τav(ps) τ1(ps) τ2(ps)195±1 54±7 154±3
233±1 39±5 164±2
273±1 46±4 162±3
148±1 47±2 157±1
156±1 53±2 161±1
148±1 53±16 149±1
149±1 84±23 153±2
149±1
146±1
-

I2(%) τ3(ps) I3(%)
70±1 433±7 17±1
68±1 511±5 21±1
55±1 523±4 29±1
79±1 489±33 2±1
72±1 549±14 4±1
96±1
93±5
-

573 K irradiation formed voids in Ni and the fourmodel alloys. The intensity of long lifetimes τ3
was, however, decreased with the addition of
Si (C and D) and no voids were formed with theaddition of C and/or P (E, F,G and J).
In all model alloys and commercial alloys,
single vacancies can move at 573 K, andvacancy clusters, voids, are only formed inmodel alloys at 573 K even with the low dose of
0.2 dpa. Therefore, the formation of voids isprevented in the commercial alloys, andstacking fault tetrahedra and/or precipitates
are formed instead. For the growth of voids inthe commercial alloys at higher dose (above
incubation dose), other factors such as
purification of the matrix by irradiation (byirradiation induced segregation, for example)
may be required, because voids are easily
formed in the model alloys.
References[1] S.Kojima, T.Yoshiie, M.Kiritani, J. Nucl. Mater. 155-157(1988) 1249.
[2] M. Horiki, M. Kiritani, J. Nucl. Mater., 212-215 (1994) 246.
[3] T. Yoshiie, X.Z. Cao, K. Sato, K. Miyawaki, Q. Xu, J. Nucl.
Mater., Accepted.
[4] P. Kirkegaard, J. V. Olsen, M. Eldrup, N. J. Pedersen, Risø
DTU, February 2009, ISBN 978-87-550-3691-8, 44 p, http://
palsfit.dk/ .
[5] B .L. Shivachev, T.Troev, T.Yoshiie, J. Nucl. Mater. 306(2002) 105.
[6] E. Kuramoto, T. Tsutsumi, K. Ueno, M. Ohmura, Y.
Kamimura, Comp. Mater. Sci., 14 (1999) 28.
[7] H. Ohkubo, Z. Tang, Y. Nagai, M. Hasagawa, Y. Tawara,
M. Kiritani, Mater. Sci. Eng. A, 350 (2003) 92.
[8] T. Yoshiie, X. Cao, Q. Xu, K. Sato, T. D. Troev, Phys. Status
Solidi C 6 (2009) 2333.

Keywords: radiation effects, neutron irradiation, nuclear materials
T.Yoshiie (Research Reactor Institute, Kyoto University)
E-mail:yoshiie@rri.kyoto-u.ac.jp
URL:http:// http://www.rri.kyoto-u.ac.jp/mre/

IMR KINKEN Research Highlights 201140

Energy-related Materials

Development of New Nanocrystalline Soft Magnetic
Materials Exhibiting ultra Low Electric Power Loss and High
Magnetic Flux Density.

F

e-rich Fe85-86 Si1-2B8 P4Cu1 nanocrystalline ribbons (~ 6 mm in width) made from industrial raw
materials (containing some impurities) were developed. These ribbons exhibit high B of 1.82 - 1.85 T(at 800 A/m), almost comparable to commercial Fe- 3 mass% oriented Si alloy. Excellent magnetic
softness (low Hc of 2.6-5.8 A/m, high permeability of 2.4-2.7x104 at 1 kHz and small saturation
magnetostriction of 2.3-2.4x10 -6) along with high electrical resistivity (0.67-0.74 μΩm) of these alloys
resulted in superior frequency characteristics of magnetic core losses (W) and a much lower W at 50
Hz up to the maximum induction of 1.75 T, in comparison to the silicon steels now in practical use forpower applications.
For power applications such as transformers
and motors, non-oriented silicon steel andoriented silicon steel are still now mainly usedbecause of their high B, despite of their muchinferior magnetic softness. Magnetic core loss(W) in silicon steels have been improved for along time, which recently seems to reach thelimitation [1, 2]. Continued growth in electrical
power generation and the recent energy
problems strongly require decreasing wastefuldissipation of energy. Making efforts to reduce
CO2 emissions and conserve resources hascome into an issue of global importance, andeven among the electronic components
industry consideration of the environment is asubject that cannot be ignored. Electrical
power loss resulting from W in motors andtransformers is estimated to be about 3.4% of
the
total
domestic
electricalpower
consumption in Japan[3]. Besides, in view ofincreasing demands of high efficiency andhigh performance electric motors for hybrid/electrical-vehicles (HEV/EV), reduction oftremendous amount of W in many kinds ofmotors and transformers requires developmentof new soft magnetic materials with highsaturation magnetic flux density (Bs) combinedwith low W.
We have developed a new Fe85-86
Si1-2B8 P4Cu1 nano-crystalline soft magnetic
material composed of ~95 mass% Fe without
rare-earth metals [4]. The new low-cost soft
magnetic material exhibits high magnetic flux
density (1.8 ~ 1.9 T) along with extremely low
core-losses (1/2 ~ 1/3 of the ordinary materials).
Fig. 1(a) shows the photograph of as-quenched
Fe83.3 Si4B8 P4Cu 0.7 alloy ribbon (width - 50 mm
and thickness ~ 23 μm). The wide ribbons canbe easily melt-spun in air due to its lower melting
temperature and its substantial resistance tooxidation. Nearly 6-mm-wide and ~ 20-μmthick ribbons can be easily obtained for Fe-rich
Fe85 Si2B8 P4Cu1 alloy. The as-quenched ribbons
show a complete ductility which should be agreat advantage as an industrial material.
Transmission electron microscopy showed a41

IMR KINKEN Research Highlights 2011

large number of extremely small α-Fe clusters(~3 nm) embedded in an amorphous materix,
i.e. a hetroamorphous for as quenched ribbons[Fig. 1(b)]. An unusual uniform nanocrystalline
structure composed of α-Fe grains with 10 - 20
nm in size [Fig. 1(c)] accompanied with anintergranular amorphous phase with ~1 nm in
width was obtained on annealing at 460 oC.

Fig. 1: (a) Outer appearance of the as-quenched
Fe85 Si2B8 P4Cu1 alloy ribbon with width of 6 mm andthickness of about 20 μm. Bright field TEM images
with selected area electron diffraction patterns
for the (b) as-quenched alloy ribbons , and (c)ribbon annealed at 460 oC.

Fig. 2 shows the changes in physical andmagnetic properties (B800, Hc, ρ and R) as afunction of annealing temperature for the
Fe85 Si2B8 P4Cu1 alloy. ρ and R exhibit a drasticchange with the precipitation of α-Fe in the

IMR KINKEN Research Highlights 2011

amorphous matrix. B800 increases from 1.13 T to
1.82 T with the structural change. Here it isnoteworthy that high B800 and low Hc are
exhibited on the crystallized Fe85 Si2B8 P4Cu1
alloy, simultaneously.

Fig.2: Annealing temperature dependence ofmagnetic flux density (B800), coercivity (Hc), density(ρ) and resistivity (R) for the melt-spun
Fe85 Si2B8 P4Cu1 alloy. Structural change withannealing is also shown for comparison.

Fig. 3 shows the changes in magnetic coreloss (W) at 50 Hz for the Fe85-86 Si1-2B8 P4Cu1
nanocrystalline alloys as a function of maximum
magnetic induction (Bm). The data of Fe78 Si9B13
amorphous, highest grade oriented Fe-3 mass%
Si and non-oriented Fe-3 mass% Si and Fe-6.5
mass% Si alloys are also shown for comparison.
W is a very important characteristic for thematerials in power applications such astransformers and motors. The amorphous alloy
with very low Hc of 2.6 A/m exhibits a very low
W in the lower Bm range than around 1.2 T, andshows a rapid increase of W with further
increasing Bm due to the magnetic saturation.
The nanocrystalline alloys show significantly
smaller W than all the silicon steels over thewhole Bm range. Here, the very low W of thenanocrystalline alloy in the Bm range of 1.5 1.75 T is 1/2-1/3 smaller than that of the oriented
Fe-3 mass% Si alloy, and overwhelmingly smaller
than those of the non-oriented Fe-3 mass %Si
and Fe-6.5 mass% Si alloys, which is important
for the power applications. Additionally,
Fe85 Si2B8 P4Cu1 nanocrystalline alloy exhibit
superior frequency characteristics of W
comparable to the typical amorphous alloyand high Tc comparable to Fe-3 mass% Si alloys,
higher R than Fe-3 mass% Si alloys and larger εfthan the ordinary nanocrystalline and theamorphous alloys. It should be noted that theseoutstanding feature never seen in other soft
magnetic materials.

Fig.3: Magnetic core losses (W) of nanocrystallized
FeSiBPCu alloys as a function of maximum
magnetic induction (Bm). The data of Fe78 Si9B13
amorphous and Fe-Si crystalline alloys are also
shown for comparison shown.

References[1] Japan Iron and Steel Federation, Bulletin of Nishiyama
Memorial Lecture, 155-156, (1995) 111-149.
[2] Nippon Steel Company Co. LTD, Nippon Steel Monthly
141, (2004).
[3] JFE 21 Century Foundation, [Online]. Available: http://
www.jfe-21st-cf.or.jp/chapter_1/1d_2.html.
[4] A. Makino et al., J. Appl. Phys. 109 (2011) 07A302.

Keywords: nanocrystalline, soft magnetic, extremely low magnetic core losses
Akihiro Makino (Advanced Research Center of Metallic Glasses)
E-mail: amakino@imr.tohoku.ac.jp
URL: http://www.arcmg.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 201142