Infrastructural Materials

Effect of Boundary Misorientation on Yielding Behavior of
Nanostructured High Purity AluminiumTo understand deformation mechanisms of nanostructured metals, effect of boundary misorientationon yielding behavior of nanostructured aluminium has been investigated. It has been suggested
that in the grain size less than 10 μm an extra hardening mechanism due to lack of dislocation
sources is responsible for the high strength, and that high-angle boundaries are more effective forstrengthening than low-angle boundaries.
Introduction
Nanostructured metals produced by highstrain deformation, with an average grain size
less than 1 μm, are found to show several uniquemechanical properties, such as unexpected
yield point phenomena in pure aluminium,
extra Hall-Petch hardening, annealing-induced
hardening and deformation-induced softening[1-4]. However, detailed deformation mechanismsof nanostructured metals have not yet been
understood. Since nanostructured metals arecharacterized by a high density of grainboundaries, interaction between dislocations
and grain boundaries is expected to be muchsignificant in nanostructured metals, and such
effect may strongly depend on the boundary
misorientation. In the present work, the effectof boundary misorientation on yielding
behavior of nanostructured aluminium has
been investigated.
Experimental
A fully-recrystallized high purity aluminium
was deformed by conventional rolling to 50%
reduction at room temperature, followed byannealing treatment to obtain subgrained
structures composed of low-angle boundaries.
In the annealing treatment, two-step annealing
process [5] was used to change the structuralsize without discontinuous recrystallization. To
produce ultrafine grained samples mainly
surrounded by high-angle boundaries, highstrain deformation followed by two-stepannealing was used, where accumulative rollbonding (ARB) of 6 cycles was applied to givea high strain [3,6].
Microstructures of the obtained samples
were characterized by electron backscatter
diffraction (EBSD) and transmission electron
microscopy (TEM).
Mechanical properties
were determined by tensile test at room
temperature, where an initial strain rate is 8.3 x
10 -4 s-1.
Results and Discussion
The 50% rolled sample showed a typicaldislocation cell structure, containing dislocation
cell boundaries and individual dislocations
between the boundaries. By annealing at
175oC for 24 h, recovery mainly occurred andthe dislocation cell structure was changed into
3

IMR KINKEN Research Highlights 2011(a) CR50+175 oC 24 h(b) 6c ARB+175 oC 6 h

1µm

1µm

Fig. 1 TEM microstructures of the deformed andannealed samples. (a) 50% rolled and annealed
at 175oC for 24 h, (b) 6-cycle ARB processed andannealed at 175oC for 6 h.

a typical subgrained structure with the average
subgrain size of 0.88 μm (Fig. 1a). Additional
high-temperature annealing led to a slight
subgrain coarsening. The EBSD measurements
have proved that boundaries in thesesubgrained structures are mainly low-angle
boundaries with misorientation angle less than
15o. Such samples mainly surrounded by lowangle boundaries are hereafter referred to as“subgrained” samples.
After ARB of 6 cycles, on the other hand,ultrafine grained structure with an averagegrain size of 0.69 μm was obtained, where mostof the boundaries in the sample are high-angle
boundaries with misorientation angle above
15o. Annealing at 175oC for 6 h caused
dislocation annihilation and slight boundary
migration, leading to a quite equiaxed finegrained structure with the grain size of 2.4 μm(Fig. 1b). By the second-step annealing at hightemperature, the grain size can be changed
gradually in the range from sub-micrometer toseveral tens micrometer.
These samples
composed of high-angle boundaries are
termed as “grained” samples, to distinguishfrom “subgrained” samples.
The stress-strain curves for the 50% coldrolled and annealed samples are shown in Fig.
2a. The as-deformed 50% cold-rolled sample
has a tensile strength of approximately 100
MPa, but after annealing of 175oC 24 h thestrength significantly decreases, corresponding

Nominal stress (MPa)to a decrease in the dislocation density. The
strength gradually decreases by further hightemperature annealing, which is due to theslight increase in the subgrain size. It is seenthat in the subgrained samples the stress-strain
curves reveal a conventional continuous
yielding.
The stress-strain curves for the grained
samples, which have been produced by
6-cycle ARB and annealing, are shown in Fig.
2b [3]. The as-deformed 6-cycle ARB sample
shows a tensile strength of 115 MPa and a
relatively large total elongation. Annealing of
175oC for 6 h leads to a significant decrease in
the strength, due to an increase in the grain
size as well as a decrease in the dislocation
density. It should be emphasized here thatsharp yield point phenomena can be observedin the early stage of tensile test. After hightemperature annealing, as well as a yield point
phenomena, a stress-constant region followed
by work-hardening, which corresponds with
the nucleation of Lüders bands and theirpropagation, are observed. However, when
the average grain size is larger than 10 μm,
yield point phenomena disappear and aconventional continuous yielding behavior isobserved.
It is clearly shown from the above resultsthat the yielding behavior of the samples
strongly
dependsontheboundary
misorientation in the fine grained region. It wasalso found that the yield stress is much higher inthe ultrafine grained samples than in thesubgrained samples at the similar grain/
subgrain size range. In other words, high-angle
boundaries are more effective for strengthening
than low-angle boundaries.
Following hypothesis can be applied tounderstand the obtained results. One importantcriterion to reveal yield point phenomena isthat samples have few mobile dislocation in an
initial state. In the ultrafine grained samples,due to the difficulty in propagation of
dislocation sources in the grains [7] or enhanced
recovery of dislocations at high-angle
boundaries, the density of dislocation sources
may be quite low. Therefore, a high stress isrequired to activate alternative dislocation
sources, e.g. at grain boundaries. On the other
hand, in the subgrained samples dislocations
in low-angle boundaries may act as easy
dislocation sources, leading to a lower yield

Nominal stress (MPa)IMR KINKEN Research Highlights 2011(a) CR50+annealed150

As deformed ( d R=1.1 µ m)
175 C 24 h (A) ( d R=0.88 µ m)o100

A + 250 C 0.5 h ( d R=1.1 µ m)o

A + 275 C 0.5 h ( d R=1.3 µ m)
o

50
0

0

10

20

30

Nominal strain (%)40(b) 6c ARB+annealed150

As deformed ( d R=0.69 µ m)
175 oC 6 h (B) ( d R=2.4 µ m)o
B + 225 C 0.5 h ( d R=3.3 µ m)
B + 250 oC 0.5 h ( d R=3.9 µ m)
B + 350 oC 0.5 h ( d R=18 µ m)

100
50
0

0

10

20

30

Nominal strain (%)

40

Fig. 2 Nominal stress-strain curves of the samples
deformed and annealed. (a) 50% cold-rolled andannealed, (b) 6-cycle ARB processed andannealed.

stress. It is therefore suggested that lack of
dislocation sources is responsible for the highstrength in nanostructured metals.
References[1] M.A.Meyers et al., Prog. Mater. Sci. 51 (2006) 427.
[2] N.Tsuji et al., Scripta Mater. 47 (2002) 893.
[3] N.Kamikawa et al., Acta Mater. 57 (2009) 4198.
[4] X.Huang et al., Science 312 (2006) 249.
[5] N.Kamikawa et al., Acta Mater. 54 (2006) 3055.
[6] Y.Saito et al., Acta Mater. 47 (1999) 579.
[7] T.Ohashi et al., Int. J. Plast. 23(2007) 897.

Keywords: nanostructure, deformation, mechanical properties
T.Furuhara, G.Miyamoto, N.Kamikawa (Microstructure Design of Structural Metallic Materials)
E-mail: kamikawa@imr.tohoku.ac.jp
URL:http://www.st-mat.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 20114

Infrastructural Materials

Structural Analysis of a Variety of Approximants for
Quasicrystalline Materials.

S

tructural analyses of o-AlPdMnSi and χ-AlPdRe approximants have been performed by single
crystal X-ray diffraction. The structure of o-AlPdMnSi phase indicates a unique packing of thedecagonal column with a periodicity of 0.72 nm and the pentagonal one with 0.48 nm, which are
different from the typical columnar structures found in Al3Mn and Al3Pd. The structure of χ-AlPdRephase indicates an i3 cluster composed of three icosahedra and a bi-capped pentagonal prism,
and a pseudo Mackay cluster. The pseudo Mackay cluster consists of irregular 1st icosahedral and2nd icosidodecahedral shells with a central heavy metal site.
After the discovery of an icosahedral
quasicrystalline phase in a rapidly solidified
Al 86Mn14 alloy by Shechtman et al. in 1984 [1], avariety of researches on quasicrystals havebeen done in order to understand their unique
physicochemical properties. For providing arealistic structure model for quasicrystals,
structural information on the corresponding
approximants has been frequently used. This
paper demonstrates the structural features of
newly found approximant phases for the
decagonal Al-Pd-Mn and icosahedral Al-PdRe [2,3].
o-AlPdMnSi approximant [2]The orthorhombic approximant phasessuch as Al3Pd, Al73.5Pd22.4Mn4.1, Al3Mn and
Al75Mn20 Pd5 indicate layered structures with 1.2
or 1.6 nm periodicity. The corresponding
columnar structures in such approximants allow
us to estimate atomic arrangements of thedecagonal quasicrystals in the Al-Pd-Mn
system. In the quaternary Al-Mn-Pd-Si system, anew orthorhombic Al72Pd18 Si5Mn5 (here after
o-AlPdMnSi; Al71.7 Pd18.5 Mn5.0 Si4.8) has beenreported nearby 1/1 and 2/1- cubic AlPdMnSi
approximants [4,5].
A needle-like crystal was cut out from anarc-melted alloy ingot (Al69Pd20Mn8 Si3), andthe structure was analyzed by single crystal
X-ray diffraction. The structure (Pnma;
a=1.4396(3), b=2.4003(2), c=7.603(2)nm) can be
represented by the stacking of three types oflayers (A,B,C) perpendicular to the b-axis. The
structure of o-AlPdMnSi consists of uniquestacking of the decagonal column with aperiodicity of 0.72nm and pentagonal one with
0.48nm (Fig.1). This feature is different from thecases of Al3Mn and Al3Pd, where the structure
is composed by the single column produced
by stacking of the corresponding layered
structures. Therefore, the unique stacking
demonstrated in the o-AlPdMnSi is suggested
to produce a new type of the decagonal
columnar structure. The columnar structuresare suggested to introduce unique periodicity
in decagonal quasicrystals.

5

IMR KINKEN Research Highlights 2011(a)[BAb]

a

b

(c)

c

0.48nm0.72nm(b)[C2C1C2i]

a

b

c

b

c2
c1
c2i
Bi
Ai
bi
C2i
C1
C2
b
A
B

Fig. 1: (a) The tiling of layer unit [BAb] and (b) of layer
unit [C2C1C2i]. The double circles indicate theposition of Pd/Al(1) and ovals are split sites Al(24)/
Pd(24). (c) The arrangement of columnar structurewith the periodicity of 0.48 and 0.72 nm is shown.

χ- AlPdRe approximant [3]In the ternary Al-Pd-Re system, a stableicosahedral i-Al70 Pd20 Re10 and a decagonal
D-Al78 Pd8Re14 have been reported.
A single crystal sample of χ-AlPdRe (P31ca=1.23744(4), c=2.7487(1)nm) was cut from the
prepared ingot of Al 80 Pd15Re5 (annealed at
750oC for 10 days).
The structure is isotypic to that of Ir 9Al28 [6].
The heavy metal (Pd and Re) sites of χ-AlPdRecorrespond to Ir sites of Ir 9Al28, suggesting thetypical pseudo-decagonal motifs together
with the pseudo Mackay type icosahedralcluster around a Re site.
Fig. 2 indicates sub-structure units of Pd(8)centred icosahedra (Pd8I) and Pd(5)- centredbi-capped
pentagonal
prism(Pd5BPP)
composedby
heavy
metals.
Pd8I

IMR KINKEN Research Highlights 2011

interpenetrates with each other and produces
an I-type connection [7]. Such I-typeconnections are further connected via acommon vertex Pd(5) so as to form a decagonal
columnar structure along the a-axis. The
decagonal columns run parallel to threeequivalent directions and intersect eachother.

c
(i)
Re(4)

(i-ii-iii-iv)(i-ii)(ii)
7Al[0.245nm]

(iii)

a

Re(4) centered
Pseudo-Mackay Cluster(iii-iv)a

12Pd[0.438nm]

a

(iv)

Pd(8) centred icosahedron
12Pd30Al[0.477nm]

a30Al[0.491nm]

ai3 cluster
Pd(5) bi-capped pendagonal prizm

Re Pd AlPd Re

a

Decagonal columnar structure

Fig.2: Pd(8)-icosahedron, Pd(5)- bi- capped
pentagonal prism, pentagonal column, and i3
cluster are shown.

Fig. 3 shows the psudo-Mackay typeicosahedral cluster (p-Mackay) together with
its network along the c-axis with the Pd(5)triangle junctions. The first shell of p-Mackay isan irregular shaped polyhedron of 7 Al with
Re(4) centre (i-ii). The second shell with a radius
about 0.48nm (iii-iv) may be ranked into two
sub-shells. One is an icosahedron produced by
12 Pd (iii) and each vertex of the Pd-icosahedronis capped by 5 Al, so as to form the Alicosidodecahedron (iv). The second shellstructure well reproduces the atomic
arrangement in the typical Mackay cluster,and is similar to that observed in the structuralmodel of i-AlPdRe.
These structural features of χ-AlPdRe are
well attributed to its chemical composition
nearby the center of those for i-phase andD-phase in the Al-Pd-Re system.

Linkage of 12Pd icosahedrons

Fig.3: Re(4)-pseudo-Mackay cluster (i-ii-iii-iv) can beranked into 4 components (i)-(iv). The values inparentheses indicate the averaged distances of
the consisting atoms from the centre of the cluster.
The linkage of the Pd-icosahedra along the c-axis
is also shown.

References[1] D. Shechtman et al., Phys. Rev. Lett. 53 (1984) 1951.
[2] R.Simura et al., Phil. Mag. 91 (2011) 2603.
[3] S. Suzuki et al., Phil. Mag. 91 (2011) 2610.
[4] K.Sugiyama et al., Z. Kristalogr. 213 (1998) 90.
[5] K.Sugiyama et al., Z. Kristalogr. 213 (1998) 168.
[6] S.Katrych et al., J. Alloy. Compd. 407 (2006) 132.
[7] G. Kreiner, and H.F. Franzen, J. Alloy. Compd. 221 (1995)
15.

Keywords: crystal structure, X-ray diffraction, quasicrystal
K.Sugiyama and R.Simura (Chemical Physics of Non-Crystalline Materials)
URL: http://www.xraylab.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 20116

Infrastructural Materials

Development of Novel Titanium Alloys with Self-adjustable
Young’s Modulus for Biomedical Applications

Young’s modulus of metallic implant device should be sufficiently low to prevent stress shielding forthe patient. However, such the Young’s modulus should be sufficiently high to suppress springback
for the surgeon. In order to satisfy these requirements simultaneously, we proposes a new concept,
self-adjustment of Young’s modulus, for biomedical titanium alloys. This concept can be achieved
using deformation-induced phase transformation; the Young’s modulus of only the deformed part
would increase during operation, while that of the non-deformed part would remain low. In thisstudy, the deformation-induced ω phase transformation was introduced into a β-type titanium alloy.
As a result, the Young’s modulus of the titanium alloy could be increased by deformation.
The stress shielding effect is one of the
problems for metallic implant devices because
it causes excessive bone absorption. It isbelieved that the Young’s modulus equal tothat of bone is ideal for metallic biomaterials to
inhibit the stress shielding effect. Therefore,
extensive efforts have been made to develop
β-type titanium alloys having low Young’s
moduli (40–60 GPa) similar to that of bone (10–30 GPa). One such alloy developed by theauthors [1], Ti-29Nb-13Ta-4.6Zr alloy (TNTZ), is apromising candidate for a next-generation
metallic biomaterial, and attempts have been
made to utilize this alloy in practical applications
such as the production of spinal implant rods.
However, when using TNTZ, some collaborating
surgeons specializing in spinal diseases pointed
out that the amount of springback in the spinal
implant rod should be small to obtain better
handling ability during operations. The amountof springback is considered to depend on boththe strength and the Young’s modulus of the
spinal implant rod. If two spinal implant rods
having the same strength and different Young’s
moduli are used, the spinal implant rod having
a lower Young’s modulus shows greater
springback. The spinal implant rod made of
TNTZ exhibits a low Young’s modulus, resulting in
greater springback.
Thus, a low Young’s
modulus, which is one of the outstanding
features of TNTZ as a metallic biomaterial andobviously a desirable property for patients,
becomes an undesirable property forsurgeons.
With this background, in this study, we
attempted to devise a solution using
metallurgical engineering to satisfy therequirements of both surgeons and patients for
the Young’s modulus of spinal implant rods. Fig.
1 shows the concept of self-adjustment ofYoung's modulus in spinal implant rod [2].
Generally, the Young’s modulus of metals does
not change drastically by deformation.
However, for certain metastable β-type titanium
alloys, non-equilibrium phases such as α′, α″,and ω appear in the β matrix during
deformation. If the Young’s modulus of thedeformation-induced phase is higher than that7

IMR KINKEN Research Highlights 2011

Implant rod
composed of
β phaseMatrix: β phase
Young’s modulus of
the non-deformed
part remains low.

Bending deformation
during operation

High Young’s modulus
is achieved forthe deformed part.

Deformation-induced
phase having highYoung’s modulus

Fig.1: Concept of self-adjustment of Young's modulus
in spinal implant rod.

of the original β phase, the Young’s modulus ofonly the deformed part of the spinal implantrod increases, but that of the non-deformed
part remains low. In orthopaedic operations
for treating spinal diseases, the spinal implantrod is bent by surgeons so that it corresponds tothe curvature of the spine. Therefore, if asuitable titanium alloy is employed as the spinal
implant rod material, springback can besuppressed by deformation-induced phasetransformation occurring during bending in thecourse of the operation, whereas a low Young’s
modulus can be retained for patients.
Among
the
abovementioned
nonequilibrium phases, ω phase is expected toincrease the Young’s modulus because it is wellknown that the presence of the ω phase,formed by aging treatment at a relatively low
temperature below ~673 K, drastically increases
the Young’s modulus of β-type titanium alloys[3]. Further, the occurrence of the deformationinduced ω phase transformation has beenreported in some β-type titanium alloys [4].
However, the effect the deformation-induced
ω phase transformation on the change in the
Young’s modulus has never been focused.
Therefore, in this study, the change in the
Young’s modulus of a metastable β-type
titanium alloy by deformation was examined.

IMR KINKEN Research Highlights 2011100
Young’s modulus (GPa)

According to the previous report[4], Ti-12Cr
alloy (Ti-12Cr) was selected as a candidate
having a chemical composition in which thedeformation-induced ω phase transformation
is likely to occur [2]. In the results of optical
microscopy, the microstructure of Ti-12Cr
subjected to solution treatment (ST) comprised
equiaxial grains. In contrast, after cold rolling(CR), mechanical twins were observed inside
the equiaxial grains of Ti-12Cr. The electron
diffraction patterns of the [110]β zone obtainedby Transmission electron microscopy revealthat only a circular diffuse streak was observedin combination with spots derived from the βphase in Ti-12Cr-ST. However, extra spots inaddition to the reflection from the β phase,
considered to be derived from the ω phase,
could be observed in Ti-12Cr-CR. These results
confirm the occurrence of deformationinduced ω phase transformation accompanying
mechanical twinning in Ti-12Cr during coldrolling.
Fig. 2 shows the results of Young’s modulus
measurement for Ti-12Cr-ST and Ti-12-CR [2]. In
this figure, the Young’s moduli of our previously
developed TNTZ, which does not exhibit
deformation-induced phase transformation
during cold rolling [3], are also shown forcomparison [2]. Ti-12Cr-ST exhibits a low Young’s
modulus of <70 GPa; this value is comparable
to that of TNTZ-ST, which was developed as abiomedical β-type titanium alloy having a low
Young’s modulus. TNTZ-CR also shows a low
Young’s modulus that is almost equal to that of
TNTZ-ST. Thus, cold rolling barely causes any
change in the Young’s modulus of TNTZ.
However, the Young’s modulus of Ti-12Cr
increases by cold rolling and that of Ti-12Cr-CR
exhibits > 80 GPa. After cold rolling, both alloys
have different phase constitutions: no newphase formation was reported in TNTZ, but the
deformation-induced ω phase was detectedin Ti-12Cr. As mentioned earlier, it is well known
that the ω phase drastically increases theYoung’s modulus of β-type titanium alloys [3].
Therefore, the increase in the Young’s modulus
of Ti-12Cr is considered to be derived from the

80
60
40
20
0

ST CR

ST CR

TNTZ
Ti-12Cr[Ref.2](This
study)
Fig. 2: Young's moduli of Ti-12Cr alloy (Ti-12Cr) andTi-29Nb-13Ta-4.6Zr alloy (TNTZ) subjected to solution
treatment (ST) and cold rolling (CR).

deformation-induced ω phase transformation.
For spinal implant rods, the Young’s modulus
should be sufficiently high to suppress
springback for the surgeon, but it should also
be sufficiently low to prevent stress shielding forthe patient. Therefore, this study proposes anovel concept—self-adjustment of Young’s
modulus. In order to achieve this concept,
the change in Young’s modulus of titanium
alloys through deformation-induced phasetransformation was investigated using Ti-12Cr.
This
alloy
successfully
underwent
thedeformation-induced ω phase transformationand then exhibited the increase in Young’s
modulus by deformation.
References[1] D. Kuroda et al., Mater. Sci. Eng. A 243 (1998) 244.
[2] M.nakai et al., Mater. Lett. 65 (2011)688.
[3] T. Akahori et al., Mater. Sci. Eng. C 25 (2005) 248.
[4] S. Hanada et al., J. Mater. Sci. 21 (1986) 4131.

Keywords: alloy, biomedical, phase transformation, Young’s modulus
M.Niinomi (Biomaterials Science Division)
E-mail: niinomi@imr.tohoku.ac.jp
URL: http://biomat.imr.tohoku.ac.jp/

IMR KINKEN Research Highlights 20118

Infrastructural Materials

Fabrication of Novel Cu-Ti Alloys by Aging in Hydrogen
AtmosphereAging Cu-Ti alloys in hydrogen atmosphere significantly improves their electrical conductivity.
Based on this finding, we proposed a model fabrication process of Cu-Ti alloys to improve both oftheir strength and conductivity. When the alloy was severely prior-deformed and then aged at a lowtemperature of 623 K under a high hydrogen pressure of 0.8 MPa, the tensile strength was beyond
1000 MPa and the conductivity was over 30 %IACS, The balance of the mechanical and electricalproperties was comparable with that of commercial Cu-Be alloys.
Copper (Cu)-based alloys for electrical
products such as lead frames and connectors
are required to have high strength and goodelectrical conductivity. Among representative
Cu alloys, Cu-Be alloys have been commerciallyapplied because of their excellent balance ofthe properties [1]. However, health hazards
associated with Be element have posedconsiderable concern, and substitutions for
Cu-Be alloys are sought. As a candidate
materials,
age-hardenableCu-Ti
alloyscontaining approximately 1 to 6 at.% Ti are
attractive
becauseof
theirsuperior
mechanical properties comparable to Cu-Be
alloys [2]. While, the conductivity of the alloys
falls somewhat below that of Cu-Be alloys. Froma practical viewpoint, therefore, a way toenhance the conductivity of the alloys without
degradation of the strength is much needed.
Semboshi et al. has reported that the
conductivity of Cu-Ti alloys increased
significantly when the alloys were aged rather
in hydrogen atmosphere than conventionallyin vacuum [3,4]. In the subsequent studies, theinfluence of aging procedure, such as hydrogenpressure, aging temperature and priortreatment on their strength and conductivity
has been investigated; aging under a highhydrogen-pressure allowed for a rapid
improvement in the conductivity [5]. Aging at a
low temperature provided a reasonablebalance of these properties, although a long
aging-time was required [6]. Cold-rolling prior
to aging is useful to accelerate the improvementof both the strength and conductivity during
the aging [7]. Based on these findings, in thisstudy, we proposed a model fabrication
procedure for novel Cu-Ti alloys with anexcellent balance of the strength andconductivity, comparable to commercial CuBe alloys.
We prepared supersaturated Cu-4.2 at.% Ti
specimens. They were prior-deformed at 298 K
in a reduction rate of 60% and then aged at
623 K under a hydrogen pressure of 0.8 MPa.
Fig. 1 shows the relationship between tensile
strength and conductivity of the specimens.
For comparison, that of Cu-Ti alloys
conventionallyagedinvacuumandcommercial Cu-Be alloys are also shown [8].
9IMR KINKEN Research Highlights 2011

The specimen exhibited a maximum strengthof 1052 MPa with a conductivity of 14% IACS(%IACS: the percentage based on the
conductivity of international annealed copper
standard at 298 K) by the aging for 24 h in thehydrogen atmosphere. On the over-aging
stage for 120 h, the specimen still had a strengthof more than 1000 MPa and high conductivityof 30% IACS. The balance of the strength andconductivity was superior to that for
conventional Cu-Ti alloys, and comparablewith that for some commercial Cu-Be alloys.
The balance of the properties must be
furthermore improved by using the alloys priordeformed more severely and aging at alower temperature under a higher hydrogenpressure.
The microstructural evolution during priordeformation and then aging in hydrogen
atmosphere can be summarized as follows [9];high-dense structural defects such as

Fig. 1: Relationship between the tensile strength andelectrical conductivity of Cu-4.2 at.% Ti specimens
cold-rolled down to a reduction of 60% and thenaged at 623 K in a hydrogen pressure of 0.8 MPa.
The number beside each plot indicates the aging
time in hours. For comparison, the relationships ofCu-Ti alloys conventionally aged in vacuum, andcommercial Cu-Be alloys are also shown [8].

IMR KINKEN Research Highlights 2011

dislocations and deformation twins, withseveral 10 nm in width, were generated byprior-deformation in the specimens. On an
early stage of aging in hydrogen atmosphere,
Cu4Ti particles dispersed finely overall the
specimen, similar to the case of aging in
vacuum. During the subsequent aging, theprecipitation of Cu4Ti proceeds and TiH2
particles are spontaneously formed by the
combination of diffused hydrogen with titanium
in the matrix (see Fig. 2). The nucleation of Cu4Ti
and TiH2 particles is promoted by priordeformation, which is attributed to structuraldefects [7].
Strengthening of the Cu-Ti dilute alloys by
aging in a hydrogen atmosphere is primarily
due to the precipitation of Cu4Ti particles,
according to the Orowan mechanism [4, 10]. In
addition, the improvement in conductivity iscontrolled by the concentration of titanium
dissolved in the matrix, which is efficiently
reduced by the co-formation of Cu4Ti and TiH2[3]. In the specimens severely deformed andthen aged at a low temperature under a highhydrogen pressure, the number density of Cu4Ti
precipitates is effectively increased, which
results in efficient dispersion-strengthening.
Furthermore, the nucleation of TiH2 precipitatesis accelerated in the deformed specimen, sothat the conductivity is increased more rapidly
during the aging. Thus, precipitation of Cu4Ti
and TiH2 particles with a high number density
and a high volume fraction is available for CuTi alloys to possess an excellent balance of thestrength and conductivity.

Fig. 2: Transmission electron microscopy (TEM) imageof Cu-4.2 at.% Ti specimen cold-rolled to a
reduction of 60%, and then aged at 653 K for 48 h
in a hydrogen pressure of 0.8 MPa. Cu4Ti particles,of several nano-meters in size, and TiH2 particles,of approximately 10 nm, between deformation
twin boundaries are found.

References[1] P.J. Rioja and D.E Laughlin, Acta Metall., 28 (1980) 1301.
[2] A. Datta and W.A. Soffa, Acta Metall., 24 (1976) 987.
[3] S. Semboshi and T.J. Konno, J. Mater. Res., 23 (2008)
473.
[4] S. Semboshi et al., Mater. Sci. Eng. A, 517 (2009) 105.
[5] S. Semboshi et al., Mater. Trans., 52 (2011) 605.
[6] S. Semboshi et al., Metall. Mater. Trans. A, 42A (2011)
2136.
[7] S. Semboshi et al., Mater. Sci. Forum, 645-646 (2010)
1315.
[8] S. Nagarjuna et al., Mater. Sci. Eng. A, 259 (1999) 34.
[9] S. Semboshi et al., Ultramicroscopy, 109 (2009) 593.
[10] W.E. Wang, J. Alloys Comp., 238, (1996) 6.

Keywords: mechanical properties, electrical properties, nanostructure
S. Semboshi (Hydrogen Functional Materials Division)
E-mail: semboshi@imr.tohoku.ac.jp
URL:http://www.hydrogen.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201110

Infrastructural Materials

Fracture and Hot Deformation Behavior of Co based Alloys
Governed by Extremely Low Stacking Fault EnergyCobalt based alloys have been investigated extensively to apply for medical implants and structuralcomponents used in aerospace or automobile industries. We have revealed the fracture and hot
deformation behavior of Co-Cr-Mo alloys which are governed by the unique dislocation mechanisms
and nature of phase stability of Co-Cr-Mo alloys. Phase-field simulation has been applied to examine
the Suzuki effect observed in a Co-Ni-Cr-Mo alloy at higher temperatures. It has been establishedthat the Suzuki effect, i.e., solute segregation to a stacking fault bounded by the Shockley partial
dislocations, is essentially concomitant with high temperature deformation of Co based alloys.
Cobalt-chromium-molybdenum (Co-CrMo)- based alloys have been used in various
medical applications because of their excellent
biocompatibility and mechanical properties.
Here, we examined the fracture behavior of
Co-Cr-Mo alloy consisting of γ (fcc) single
phase analyzed by X-ray tomography and thedeformation mode of Co-Cr-Mo alloy consistingof ε (HCP) single phase analyzed by EBSD
method.
The damage process leading to fractureduring tensile testing of a biomedical grade
Co–29Cr–6Mo–0.14N alloy was analyzed on thebasis of three-dimensional damage observation
using X-ray tomography and electronbackscattered diffraction of the fractured
specimen (Fig.1). Initial cracking occurred at
grain and annealing twin boundaries, where
strain concentrates due to impingement of
ε-HCP plates formed through strain-induced
martensitic transformation (SIMT). Cracks
propagated along the interface between the
γ-FCC matrix and SIMTed ε-HCP on {1 1 1} plane,
resulting in a quasi-cleavage fracture [1].
Regarding the deformation mode of CoCr-Mo alloy consisting of ε (HCP) single phase[2], basal <a> and prismatic <a> slips were thedominant deformation modes. The activationof these <a> slips is strongly dependent on
Schmid factor, which indicates that the valuesof critical resolved shear stress (CRSS) for basal<a> slip and prismatic <a> slip are similar. For
tensile strains up to 10%, the frequent
occurrence of deformation twins could not be
confirmed by EBSD analysis.
Further, the article [3] presents the dynamic
recrystallization (DRX) behaviors of Co-29Cr6Mo alloy with additions of both C and N(hereafter CCMCN alloy). The results revealed
that the initial microstructure is a stable γ facecentered cubic (FCC) phase with a large
number of M23C6 precipitates both inside thegrains and at grain boundaries. The DRXed
grains were observed to be uniformly distributed
and to decrease with strain (Fig.2). The highvolume fraction of Σ3 boundaries was observed
after the DRX, indicating a close relation
between the DRX mechanism and Σ3 boundary
formations. In addition, the activation energy
11

IMR KINKEN Research Highlights 2011

Q of CCMCN alloy was observed to be higher
as compared to those of alloys without C or N
addition and that with N addition.
As demonstrated above, the mechanicalproperties of Co-based alloys are closely
related to the stability of HCP-structure andFCC-structure. Besides the above mentioned
detrimental aspect, the SIMTed ε-HCP in Co-CrMo alloys has beneficial aspects as well. It canenhance Young’s modulus, and improvestrength and wear resistance [4]. Co-Ni-based
superalloys with relatively high FCC-phasestability exhibit high ductility by dislocation slip

Fig.1 3D tomographic reconstruction of Co-29Cr-6Mo0.14N alloy at three different relevant damage steps:(a), (b) initial state; (c), (d) crack initiation state and(e), (f) the state just before fracture. Fig. 1 (a), (c), (e)and (b), (d), (f) indicate the outer surface and theinner cracks inside the specimen, respectively.

Fig. 2 (a) Microstructure of deformed sample at 1423
K at 1 s-1 at reduction rate of 60%. (b) Mean grain
size as a function of reduction rate.

IMR KINKEN Research Highlights 2011[5]. The stability of HCP phase affects the natureof stacking faults which governs the plastic
behaviors by both dislocation slip andtwinning.
The phase-stability at stacking fault can besignificantly affected by segregation of soluteatoms, i.e. Suzuki segregation. To date, Suzuki
segregation in Co-alloys has never beenverified experimentally despite many attempts
by researchers all over the world. Atomistic
simulations are commonly applied to studies ofinterfacial segregation, but it is quite difficult toapply atomistic calculation to multinary systems
such as as cobalt-alloys.
Phase-field method is a simulation method
based on thermodynamics and diffuse
interface model, which is applicable tomicrostructural evolution and interfacial
reaction in nanometer and micrometer lengthscale. It has been developed as powerful tool
to bridge the gap between the atomistic andmacroscopic simulations. In the present study,
we have developed a phase-field model ofsegregation in order to examine the Suzukisegregation in Co alloys (Fig.3).
Fig.4 shows the distributions of alloying
elements at and around a stacking fault in Co30Ni-20Cr-10Mo (wt%) alloy (hereafter, denotedas CNCM-alloy) equilibrated at 973 K in asimulation where elastic strain energy is taken
into account [6]. Ni and Nb were strongly
depleted and Cr segregated at the stacking
fault, while Mo exhibits neither segregation nor
depletion. In a simulation in which elastic strainenergy is neglected, Mo segregation wasobserved, implying that the chemicalinteraction for segregation and the size effectis compensating each other. The stacking fault
energy dramatically decreased from 35 mJ/m2
to -937 mJ/m2, as a result of the segregation.
This agree with the experimentally observed
spontaneous formation of very wide stacking
faults in the deformed and subsequently
annealed CNCM-alloy (Fig. 5).
Thus, the segregation at stacking fault in
Co-alloy has been computationally verified forthe first time. This approach can be applied tothe alloy designing for further improvements ofbiomedical and industrial Co-alloys.
References[1] S. Kurosu , H. Matsumoto, A. Chiba, C. Landron, D.
Fabregue and E. Maire: Scripta Materialia, 64 (2010) 367370.
[2] H. Matsumoto, S. Kurosu, B.S. Lee, Y. Li, A. Chiba: Scripta

Fig.3 Model of stacking fault used for the calculation.

Fig.4 Distributions of (a) Cr, (b) Mo, (c) Nb and (d) Ni
at and around a stacking fault in Co-30Ni-20Cr10Mo (wt%) alloy equilibrated at 973 K.

Fig.5 TEM image of stacking fault in Co–Ni-based
superalloy single crystal compressed to 2 % andsubsequently annealed at 973 K for 3.6 ks.
Materialia, 63 (2010) 1092-1095.
[3] Y. Yamashita, Y. Li, E. Onodera, H. Matsumoto and A.
Chiba, Mater.Trans. 51 (2010) 1633-1639.
[4] A. Chiba, K. Kumagai, N. Nomura, S. Miyakawa . Acta
Materialia 55 (2007) 1309-1318.
[5]A. Chiba X. G. Li and M. S. Kim, Philosophical Magazine
A 79 (1999) 1533-1544.
[6] Y. Koizumi, S. Suzuki, T. Otomo, S. Kurosu, Y.P Li, H.
Matsumoto and A. Chiba, Supplemental Proceedings:
vol.2 Materials Fabrication, Properties, Characterizationand Modeling, TMS 2011, pp.409-416.

Keywords: Co-Cr-Mo alloy, fracture behavior, Suzuki effect.
A. Chiba, Y. Koizumi, H. Matsumoto, and Y. Li (Deformation Processing).
E-mail: a.chiba@imr.tohoku.ac.jp.
URL: http://www.chibalab.imr.tohoku.ac.jp/english/index.php

IMR KINKEN Research Highlights 201112

Infrastructural Materials

Excitation Mechanism of Silver Ionic Lines throughCollisions with Metastable Atoms of Plasma Gas in Glow
Discharge Plasmas.

Emission spectra of silver from glow discharge plasmas using argon, krypton, and a krypton-argon

gas mixture were analyzed to elucidate the excitation phenomenon of silver ionic lines occurring incollisions with the plasma gases. The intensities of the ionic lines which were assigned to the 4d95p- 4d95s transition were largely different between argon and krypton; for instance, the intensity of the
Ag II 243.78-nm line emitted form the krypton plasma was 5 times larger than that from the argon
plasma, while the Ag 228.00-nm line yielded much larger intensity in the argon plasma by a factorof 70. This effect can be explained by an excitation process through collisional energy transfer fromthe metastable atom of argon or krypton to a silver ion in the ground state (Penning excitation). This
type of collision requires both matching in the excitation energy between the colliding partners andconservation in the total spin before and after the collision. The 3p54s metastable of argon (11.72 eV
and 11.55 eV) and the 4p55s metastable of krypton (10.56 eV and 9.92 eV) would take part in thecollision, and principally determine the population of the 4d95p excited levels of the silver ion.
Glowdischarge
opticalemission
spectrometry (GD-OES) is extensively employed
for the direct analysis of solid samples, because
the GD source has several benefits such as
rapid sampling through cathode sputtering,
minimal sample pretreatment, and a wide
concentration range in the quantification.
Ionization /excitation phenomena occurring in
the plasma should be investigated to determine
the analytical line as well as the optimum
discharge conditions in GD-OES. Previous
studies have reported that intense emission
lines of the singly-ionized atom are observed inseveral elements, such as copper [1, 2], nickel[3], cobalt [4], and iron [5], of which spectrum
patterns are largely varied by the kind of theplasma gas employed. As similar to all otherplasma sources operating at low pressures, thedifference is not only in their relative intensities
but in the kind of emission lines that can be
excited for each plasma source [6]. This effectpromoted the studies on the excitation process
in GD-OES.
Several papers indicated that the excitationof ionic lines could be caused through a chargetransfer collision in the GD plasma [7]. This typeof collision is generally represented by thefollowing equation: Mg + Ar+ →M+* + Arg + ΔE,
where the superscripts g, +, and * represent a
ground state, an ionic state, and an excitedstate, respectively, and ΔE means the energy
difference before and after the collision.
However, it is considered that the excited levels
belonging to the 4d95p electron configurationof silver ion (the 4d95p excited levels) would be
little populated through the charge transfer
collision with the argon ion because the total
excitation energy of the 4d95p levels is fairly
larger than the internal energy of the argon
ion, whereas the emission lines assigned to the
4d95p-4d95s transition can be clearly observedin the argon GD plasma [8]. Therefore, any
13

IMR KINKEN Research Highlights 2011

different process should be taken into accountfor this excitation. Our study suggests an
excitation mechanism in which a metastable
atom of the plasma gas is involved, represented
by the following equation: Ag+g + Arm → Ag+* +Arg + ΔE, where the superscript m represents a
metastable state [9]. This mechanism well
explains the difference in the emission intensitiesof the silver ionic lines between in an argon
plasma and in a krypton plasma.
Excitation mechanism of Ag ionic lines [9]Fig. 1 shows a detailed energy level
diagram regarding the 4d95p excited state ofsingly-ionized silver when silver ions in the
ground state are in the plasma with atomic
species of argon or krypton. One can findseveral correspondences in the excitationenergy between the 4d95p excited levels andthe metastables of argon, 3p54s [3/2]2 (11.55
eV) and 3p54s [1/2]0 (11.72 eV) [10], and of
krypton, 4p55s [3/2]2 (9.92 eV) and 4p55s [1/2]0(10.56 eV) [11]. It is thus suggested that an
excitation mechanism in which a metastable
atom of the plasma gas is involved, represented
by the following equation: Ag+g + Ar m (or Kr m)→ Ag+* + Arg (or Krg )+ ΔE. This mechanism, socalled Penning excitation, well explains thedifference in the emission intensities of the silver
ionic lines between in an argon plasma and ina krypton plasma.
Fig. 2 shows a plot of the ratio of the intensity
in a pure argon plasma to that in a pure krypton
plasma for 6 Ag I lines and 27 Ag II lines against
the sum of the first ionization potential and theexcitation energy (total excitation energy). The
intensity ratios of the Ag I lines are 1.08 - 1.20.
Their intensities emitted from the argon plasma
are slightly larger than those from the krypton
plasma but they are not so different, thus
indicating that the sampling amounts are
similar between these plasmas. On the otherIMR KINKEN Research Highlights 2011

Fig. 2 Ratios of the emission intensity observed in an
argon plasma to that in a krypton plasma for silver
emission lines as function of their total excitationenergies. Sample: pure silver; plasma gas pressure:670-Pa Ar / 670-Pa Kr; discharge voltage: 500 V(fixed). After Ref. [9].

Fig. 1 Correlation of the energy level diagram
between the 4d95p excited levels of silver ion andthe metastable levels of argon and krypton. After
Ref. [9].

hand, the intensity ratios of the Ag II lines
assigned to the 4d95p-4d95s transition differ
greatly depending on the excitation energy. In
the Ag II 243.78-nm line (from the 4d95p 3 P2level, 17.52 eV), the intensity from the argon
plasma is 5 times smaller than that from thekrypton plasma; however, the emission
intensities from the krypton plasma become
drastically reduced with increasing theexcitation energy whereas those from theargon plasma become elevated, resulting in
the intensity ratios of more than 10 at total
excitation energies of 18.5 - 18.9 eV. Singlet Ag
II lines, such as the Ag II 222.96 nm (from the
4d95p 1D2 level, 18.85 eV) and the Ag II 228.00
nm (from the 4d95p 1P1 level, 18.71 eV), havemuch larger emission intensities in the argon
plasma than in the krypton plasma.
A particular 4d95p level of the silver ion is
selectively excited in the krypton plasma,
whereas all of the 4d95p levels can be

populated in the argon plasma.
A
correspondence in the excitation energy
between the 4d95p excited levels of the silver
ion and the argon or the krypton metastables is
the most important factor for determining thespectrum pattern from these glow discharge
plasmas.
References[1]E.B.M. Steers and R.J. Fielding, J. Anal. Atom. Spectrom.,
2 (1987) 239-244.
[2]K. Wagatsuma and K. Hirokawa, Spectrochim. Acta PartB, 46 (1991) 269-281.
[3]K. Wagatsuma and H. Honda, Spectrochim. Acta Part B,
60 (2005) 1538-1544.
[4]K. Wagatsuma, Anal. Bioanal. Chem., 393 (2009) 20672074.
[5]K. Wagatsuma and K. Hirokawa, Spectrochim. Acta Part
B, 51 (1996) 349-374.
[6]K. Wagatsuma and K. Hirokawa, Anal. Chem., 57 (1985)
2901-2907.
[7]O.S. Duffendach and J.G. Black, Phys. Rev., 34 (1929) 3543.
[8]K. Wagatsuma, Z. Phys., D37 (1996) 231-239.
[9]K. Wagatsuma, Spectrosc. Lett., 44 (2011) 151-157.
[10]C.E. Moore, Atomic Energy Levels Vol. I, NBS Circular
467, Washington DC, (1948).
[11]C.E. Moore, Atomic Energy Levels Vol. II, NBS Circular
467, Washington DC, (1952).

Keywords: atomic emission spectrometry, glow discharge plasma, Penning excitation
K.Wagatsuma, S. Sato and S. Kashiwakura (Analytical Science)
E-mail: wagatuma@imr.tohoku.ac.jp
URL:http://wagatsuma.imr.tohoku.ac.jp

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Infrastructural Materials

Synthesis, Boron Nonstoichiometry and Properties of
Perovskite-type ErRh3Bx Ternary Borides

P

olycrystalline of ErRh3Bx is synthesized by the arc-melting
method. The ErRh3Bx compounds have a_perovskite-type cubic structure (space group: Pm3m). The solid solution range of perovskite-typephase exists for the boron concentration from x = 1.000 to 0.610. The lattice parameter a depends on
x, and varies linearly from a = 0.4151 nm (x = 1.000) to 0.4063 nm (x = 0.610). The micro-Vickers hardnessof the ErRh3Bx increases with increasing their B content. The hardness of stoichiometric ErRh3B1.000 is
8.6 ± 0.3 GPa. The thermo-gravimetric curve indicates that the onset-temperature of oxidation of
ErRh3B1.000 is 1303 K. The weight gain of ErRh3B1.000 during heating in air to 1473 K is 0.7 %.

15

Many studies of simple and complex
perovskite and Cu3Au-type oxides have been
performed due to the interesting features ofsuperconductor, insulator-metallic transition,
ion conduction characteristics, dielectric
properties, and ferroelasticity [1-3]. On theother hand, there have been few studies on
non-oxide
perovskite
andCu3Au-type
compounds such as RERh3Bx (RE: rare earth
metal), where RE and, Rh and B atoms are
positioned at vertices, the face-centered sites
and the body-centered site of the cubic,
respectively, as illustrated in Fig. 1.

(tetragonal system, space group: P42/nmc) also
exist close to ErRh3B. ErRh3B, ErRh3B2 and ErRh4B4
are paramagnetic, ferromagnetic and reentranttype superconducting compounds, respectively[5].
All the synthesized samples had a silvery
metallic luster. According to chemical analyses
using ICP method, the chemical compositionsof the obtained samples almost corresponded
to the atomic ratio of the starting
compositions.

Fig. 1 Schematic representation of crystal structureof ErRh3 B x.

Fig. 2 Tentative phase diagram in the Er-Rh-B system.
Red line indicates ErRh3 B x compounds.

In the present study [4], we introduce thesuccessful synthesis of the perovskite-type
ErRh3Bx ternary borides by an arc-melting
method. The solid solution range of boron in
the ErRh3Bx is clarified by X-ray diffraction (XRD)
analysis, and the change of the lattice constantof the compounds is investigated as a functionof the boron concentration x. Micro-Vickers
hardness and thermochemical properties arestudied for ErRh3Bx in order to investigate
features of the bond.
The compound corresponded to theatomic ratio Er:Rh:B = 1:3:1; stoichiometric ErRh3B
exists at the rhodium-rich portion of the Er-Rh-B
tentative ternary phase diagram as shown in
Fig. 2. Ternary compounds ErRh3B2 (monoclinic
system, space group: C2/m) and ErRh4B4

By varying x in ErRh3Bx, the nonstoichiometryof boron is investigated. In the cases of boron
content x = 0.444 and 0.210, the second phasecoexisted. Fig. 3(a) shows the variation of thelattice constant a as a function of boron
concentration x in the cubic system of ErRh3Bx.
The lattice expansion with increasing boron
concentration x in the compound supports thesolid solution range of boron in ErRh3Bx. That is,
it was found that the nonstoichiometric phase
exists for the boron concentration from x = 1.000
to 0.610. All compounds ErRh3Bx have aperovskite-type cubic structure (space group:Pm3m). On the other hand, the solid solution
range for Sc and Ce-systems exist for the boron
concentration from x = 1.000 to 0.000.

IMR KINKEN Research Highlights 2011

IMR KINKEN Research Highlights 2011

Fig. 3 (a) Lattice constant a and (b) micro-Vickers
hardness of ErRh3 B x as a function of boron content.
Open blue circles represent existence of second
phase.

Fig. 4 TG-DTA curves for stoichiometric ErRh3 B.
Samples are heated in air up to 1473 K at 10 K/min.
Data for ErRh3 B2 and ErRh4B4 are given as
references.

Fig. 3(b) shows the relationship between B
concentration x and micro-Vickers hardness Hv
for ErRh3Bx. The value of Hv is distributed over
4.0～8.6 GPa. Similar to boron nonstoichiometry
ScRh3Bx [6,7] and CeRh3Bx [8,9] compounds,
the anomalous drop of hardness around x =0.5
was noted.
On the basis of TG-DTA between room
temperature and 1473 K, it was found that theoxidation of the stoichiometric compound
ErRh3B starts at 1303 K and the final weight gainis 0.7%. The perovskite-type ErRh3B has very
high resistance for oxidation compare to layerstructure compound ErRh3B2 and clusterincluding compound ErRh4B4 (Fig. 4).
The solid solution range of boron in ErRh3Bx
is clarified as 0.610 ≦ x ≦ 1.000. The microVickers hardness for ErRh3Bx varies with x similar

to those of Sc and Ce-systems. Oxidation of
ErRh3B starts at 1303 K and the final weight gainis 0.7% (room temperature to 1473 K).
References[1] F. Galasso, Perovskites and High-Tc Superconductors,
Gordon and Breach, London, 1990.
[2] H. Iwahara, M. Balkanski, T. Takahashi and H. L. Tuller(eds.), Solid State Ionics, Elsevier, Amsterdam, 1992.
[3] P. Rogl and L. Delong, J. Less-Common Met. 91 (1983)
97.
[4] T. Shishido et al., Pacific Sci. Rev. 12 (2011) 254.
[5] T. Shishido et al., J. Solid State Chem. 133 (1997) 82.
[6] T. Shishido et al., J. Alloys Compd. 309 (2000) 107.
[7] K. Yubuta et al., J. Alloys Compd. 471 (2009) 341.
[8] T. Shishido et al., J. Alloys Compd. 426 (2006) 304.
[9] K. Yubuta et al., J. Alloys Compd. 426 (2006) 308.

Keywords: ternary boride, crystal structure, hardness, oxidation resistanceK. Yubuta and T. Shishido(Research Center of Crystal Growth, Advanced Research Center of Metallic Glasses)
E-mail: yubuta@imr.tohoku.ac.jp
URL:http://www.arcmg.imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201116

Infrastructural Materials

Magnetic Field Effect for the Decomposition Temperaturesof MnBi

Differential thermal analysis were carried out for MnBi in fields up to 45 T. Decomposition temperatures

Tt (MnBi→Mn1.08Bi + liquid Bi) and Tm (Mn1,08Bi→Mn + liquid Bi) rise 84 K and 4 K by applying 45 T,
respectively. Consequently, the equilibrium state of MnBi can be controlled by magnetic fields

Ferromagnetic compound MnBi exhibits
unique magnetic and structural properties.
With increasing temperature from room
temperature, MnBi undergoes the first ordermagnetic
phasetransition
fromtheferromagnetic to the paramagnetic state at Tt~630 K, accompanied by a decomposition
from MnBi to Mn1.08Bi and liquid Bi. On the basesof high field differential thermal analysis (DTA)
experiments for MnBi, Koyama et al. reported
that Tt increases at the rate of 2 K/T in fields up
to 14 T [1]. In order to study the effect ofmagnetic fields on the phase transitions, DTA
experiments under high magnetic fields is oneof the most important experiments. In addition,
it requires high magnetic fields to investigate
the potential of materials at the early stage ofthe development. However, there is no
instrument for thermal analysis under highmagnetic fields over 20 T in the world.
In this work, for investigating the effects ofhigh magnetic fields for the decomposition
process of MnBi, high field DTA experiments for
MnBi were carried out in fields up to 45 T using
the hybrid magnets with 32 mm room
temperature bore by combining a new
developed DTA instrument at HFLSM (IMR,
Tohoku Univ.) and NHMFL (Florida State Univ.).
Fig.1 shows the DTA probe with 32 mm indiameter for DTA experiments using the hybrid
magnets. Two Pt-PtRh type thermocouples
were inserted into the quartz tube with 14 mm
in diameter. The sample (MnBi powder) andreference sample (Al2O3 powder) werewrapped in aluminum film and attached to thetip of the thermocouples. The samples were
heated by the nichrome heater inside thewater cooling jacket. The DTA data wereobtained in fields up to 45 T and in thetemperature range from 300 K to 773 K during
the heating at a rate of 6-7 K/min.
Fig.2 (a) and (b) shows the DTA curves of
MnBi, which are measured at the HFLSM (a)and the NHMFL (b)[2]. Although the vibrationnoises of the water-cooled coils in the hybrid
magnet were detected, the endothermic
peaks due to the decompositions were clearlyobserved even in 45 T. As shown in this figure, Tt
shifts to a higher temperature. Furthermore Tm(Mn1.08Bi→Mn + liquid Bi) slightly shifts to a higher
temperature by applying high magnetic fields.
Fig.3 shows the magnetic fields dependence
17

IMR KINKEN Research Highlights 2011

Fig.1 High field DTA probe with 32 mm in diameter.

Fig.2 DTA curves for MnBi in various magnetic fields
up to 45T, which are measured at the HFLSM,
Japan (a) and the NHMFL, USA (b).

IMR KINKEN Research Highlights 2011of the transformation temperatures Tt and Tm of
the Mn-Bi compound in fields up to 45 T. The
decomposition temperature Tt increases almost
linearly with increasing B up to ~20 T at a rate of~2 K/T, as shown by the broken line in Fig. 3.
However, we can see that the data over ~20 T
deviate from the broken line (~2 K/T) and Tt
reaches very close to Tm at B = 45 T. In addition,
the peritectic temperature Tm increases withincreasing B, which is clearly observed over
40 T.
Fig.4 shows the typical results for themagnetic field dependence of the sample
temperature in the vicinity of 689 K. In thismeasurement, the sample was heated to 689.5
K in B = 11.5 T. In this condition, the sample
contains paramagnetic Mn1.08Bi and liquid
phase. After that, the magnetic field was
generated from 11.5 T to 45 T. As shown in thisfigure,weobserved
exothermic
andendothermic peaks. In other words, this is amagnetocaloric effect. Considering the phase
diagram (fig.3), exothermic and endothermic
peaks are due to the phase transitions fromparamagnetic Mn1.08Bi and liquid phase to
field induced ferromagnetic MnBi for increasing
B and from field induced ferromagnetic MnBi
to paramagnetic Mn1.08Bi and liquid phase,respectively. This is, the field induced
composition and decomposition process occur
in the Mn-Bi system.
The high field DTA measurements in fields
up to 45 T shows that not only the Tt but also Tm
increased by magnetic fields. Furthermore, we
observed magnetpcaloric effect in 11.5-45 T in
the vicinity of 689 K. Consequently, thecomposition,
decomposition temperaturesand equilibrium state of MnBi can be controlled
by high magnetic fields.

Fig. 3 Magnetic fields dependence of Tt and Tm. Blue
points and red points indicate the data taken at
HFLSM and NHFML, respectively.

References[1] K. Koyama, T. Onogi, Y. Mitsui, Y. Nakamori, S. Orimo, andK. Watanabe, Mater. Trans. 48, 2414-2418(2007).
[2] K Koyama, Y. Mitsui, E. S. Choi, Y. Ikehara, E. Palm, and K.
Watanabe, J. Alloy. Compd. 509, L78-80 (2011)

Fig.4 Magnetic field dependence
temperature in the vicinity of 689 K.

ofsampleKeywords: Differential thermal analysis, phase transitions, high magnetic fields
Y. Mitsui1, K. Koyama2, E. S. Choi3, Y. Ikehara1, E. Palm3, and K. Watanabe1
1)HFLSM, IMR, Tohoku Univerisity 2) Graduate school of science and technology, Kagoshima Univ. 3)NHMFL,
Florida State Univ.
E-mail: mitsui@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 201118