Research

Renovation of Ge-Crystal Monochromatorfor Neutron Scattering Spectroscopy

Crystallography
Applied technology

47.8Ei (λ=1.3Å), g-open-120'-B
Ge(331), 17*40*t6 mm3, (h, k, 0)(a) Magnetic Bragg, T = 10 K
Renewed
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Neutron scattering spectroscopy is a powerful probe to
study phonon and magnon excitations in condensed matterphysics. A development of neutron monochromator, which
properly branches monochromatic neutrons from a whitebeam, is therefore one of the instrumental key issues inneutron science; especially, due to the lack of neutron beamflux. Ge and Si single crytals have been one of the typicaldevices for neutron monochromator, because high-quality
and large-size crystals are easily available. However,
because of a mismatch between the incident-beamdivergence and the small mosaicity coming from perfect
crystals, the neutron reflectivity of Ge and Si is severly low
and much efforts have been devoted to overcome it.
We have tried to introduce mosaic crystals into Ge
crystals appropriately by pressing at high temperatures(Fig.1, left). Eventually, an optimal hot-pressing condition
was determined [1,2], so that Ge crystals with a mosaic widthof ~0.3° and a peak reflectivity of ~40% are well reproduced(Fig.1, right). A focusing-type Ge monochromator, which was
renewed for a KINKEN triple-axis spectrometer AKANE, isshown in Fig.1.
We substituted this renewed monochromator for theprevious one on AKANE, and first characterized the beam
size of 24w*32h mm2 at the sample position. This area isabout 1.5 times as large as the previous one, and the neutron
density per area is confirmed to be unchanged. Second, weradiated the monochromatic neutron beam (λ=2.0 Å) onto
relatively large samples and measured scattered-neutron
intensity. Figure 2(a) shows a comparison of reflected
intensity with previous one. The magnetic Bragg reflections
are enhanced in intensity by ~60%, as expected from thebeam size. Further, a newly supplied second-wavelength

19.4Ei (λ=2.0Å), g-open-60’-B
La1.5Ca0.5Co0.97Cr0.03O4, φ6*30 mm

Intensity (k cts/ 30 sec)

A Ge-crystal neutron monochromator was renewed forneutron scattering using hot-pressed Ge pieces of ~40%
peak reflectivity. The neutron beam flux from the improved
one was increased so that the scattering intensity from asample was gained by about 60%. We also succeeded in
extracting monochromatic neutrons of a second wavelengthwith a reasonable beam flux.

10(b) Phonon, T = 300 K
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Fig. 2 Spectra measured by the renovated Ge monochromator.
(a) Magnetic Bragg reflections for renewed and previous
ones. (b) Phonon group measured under the newly equipped
second-wavelength mode of λ = 1.3 Å.

mode (λ=1.3 Å) works well for high-energy excitation
measurements, as seen in Fig. 2(b). Present renovation not
only vitalizes the scientific research on AKANE, but also
opens new fundamental techniques on neutron scattering.
References[1] Y. Miyake, H. Hiraka, K. Ohoyama, Y. Yamaguchi, K.
Yamada; J. Phys. (Conf. Ser.) 200, 112006 (2010).
[2] Y. Miyake, Master Thesis of Science (Tohoku University,
Feb. in 2010).
Key Words
Neutron Scattering, Ge Monochromator, Hot Press
Contact toHaruhiro Hiraka (Metal Physics with Quantum Beam Spectroscopy
Division)
E-mail: hiraka@imr.tohoku.ac.jp
URL: http://www.yamada-lab.imr.tohoku.ac.jp/jp/index.html

Neutron Science

Neutron device

Reflectivity(Hot press)

20 cm

Mosaic crystal
dislocation

Ge(113) Monochromator

Intensity (arb. unit)

Perfect crystal

Crystal angle, θ(°)

Fig. 1 Neutron crystal monochromator is a product of applied
crystallography and promotes neutron science as afundamental device.

2

Infrastructural Materials

IMR KINKEN Research Highlights 2010

Research

Strain Accommodation in Martensitic Transformationin Ferrous Alloys
Lath martensite structure becomes more important inrecent steel technology due to increasing demands for highstrength structural steels for various applications. Grain
refinement in lath martensite is needed for both strengthening
and toughening. During growth of martensite, elastic strain isaccumulated because of the natures of a displacivetransformation. Thus, strain accommodation affects strongly
the final crystal sizes of martensite.
Here we report our recent study on mechanisms of strain
accommodation during martensitic transformation in Fe-Ni
based alloys with advanced crystallographic analyses [1-3].
Optical micrograph of Fig. 1(a) shows a representative
morphology of ferrous martensite; ‘lenticular martensite’. Thismorphology is developed in the following sequence; (1) formationof thin-plate martensite seen as ‘midrib’ which is a straight line
inside the martensite; (2) further growth of martensite into a
lenticular shape [1]. In the first stage, transformation strain isaccommodated mostly by twinning of martensite. In the second
stage, possibly due to adiabatic heating, the strain accomodation
mode of martensite changes from twinning to dislocation slip.
Furthermore, plastic deformation in austenite plays a part of strain
accomodation. Thus, high densities of dislocations are introduced
inside and outside of the martensite. We have recently developed
an advanced method to analyze strain accumulation precisely by
image analysis of electron backscatters diffraction (EBSD) patterns[2]. Figure. 1(b) shows the misorientation map of the austenite
surrounding the martensite caused by such plastic accommodation.
Figure. 1(c) shows variation of the rotation components of thestrain tensor measured along the line A-B in (b). There is strain
accumulation causing lattice rotation and its distribution isasymmetric. Since it is found that elastic components are negligibly
small, it can be concluded that the observed rotations are resulted
from the plastic accommodation in austenite. This kind of technique
can be applicable to other fields such as an analysis of crack
propagation in the brittle fracture.
Plastic accommodation of strain also introduces many
dislocations inside the martensite. However, those dislocations
are a residue of lattice invariant deformation because dislocation
slip needs to penetrate to the austenite/martensite interphaseboundary to fully accommodate the strain accumulated. Thus,
there must be a high density of dislocations at the boundary.
High-resolution TEM image of Fig. 2(a) shows the atomic
structure of an austenite/martensite interphase boundary for thelenticular martensite. There are many lattice dislocations, oneof which are enlarged in (b), at the boundary. Quantitative
analysis of those dislocations has revealed that there are two
sets of dislocations at the boundary in order to accommodate
transformation strain effectively.
It has been long considered that growth of martensite occurs
by migration of a glisile boundary. A conventional TEM analysis
indicated that each set of dislocations is in the form of glissile shear
loops. However, once the two sets interact each other or one ofthose sets interacts with dislocations in the surrounding austenite,sessile dislocation nodes are introduced to the boundary and makeFig. 1 (a) Optical micrograph of lenticular martensite (M) formedin Fe–33mass%Ni. (b) Misorientation map of austenite (γ)
matrix measured by EBSD analysis. (c) Rotation components ofthe strain tensor in γ along a line A-B.
it more sessile. It is well known that plastic deformation processingof austenite before the martensitic transformation (ausforming) is
effective for improvement of strength-toughness balance ofmartensitic steels. Clearly, the present results provide important
information for the effect of plastic accommodation in on the growthof martensite and its final grain size.
References[1] A. Shibata, S. Morito, T. Furuhara and T. Maki, Acta
Materialia, 57, 483(2009).
[2] G. Miyamoto, A. Shibata, T. Maki and T. Furuhara: Acta
Materialia, 57, 1120(2009).
[3] A. Shibata, T. Furuhara and T. Maki: Acta Materialia, 58,3477(2009).
Key Words
High Strength Steel, Phase Transformation, Deformation
Contact to
Tadashi Furuhara (Microstructure Design of Structural Metallic
Materials Division)
E-mail: furuhara@imr.tohoku.ac.jp

Fig. 2 HREM image showing the interphase boundary of lenticular martensite in Fe–33mass%Ni, viewed along [-1 1 0]A//[-1 0 0]M. (A:austenite, M: martensite), (b) Enlarged atomic structure near the interfacial dislocation in (a). (c) Schematic illustration showing thearray of interfacial dislocations in (a).

KINKEN Research Highlights 2010 IMRInfrastructural Materials

3ResearchHigh Resolution Stress Mapping of the Directionally Solidified
Oxide Eutectics by Scanning Near-Field Optical Microscopy
Thermal and transformation-induced strains aresuggested to modify the mechanical behavior, which can bestudied by means of the piezospectroscopic effect of theluminescence of Cr3+ in Al2O3. The recently developed
scanning near-field optical microscopy (SNOM) techniqueallows us to obtain the nanometer-scaled residual stressmapping. This new analystical method has been applied for
directionally solidified Al2O3/ZrO2 and Al2O3/cubic-ZrO2(Y2O3) eutectics, whch are ranked as candidates for highstrength materials at high temperature.

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Directionally solidified eutectic (DSE) samples of Al2O3/
ZrO2 and Al2O3/cubic-ZrO2(Y2O3) were prepared by the micro
pulling down (μ-PD) method using a radio frequency heating
system with a pulling rate 1.0 mm/min. Laser-induced
fluorescence spectra from the sample surface were collectedusing the SNOM system which was built into a commerciallyavailable AFM (SPM9500J2; Shimadzu Corp.). Excitation
laser was a 488 nm line of Ar+ laser and was focused using a
SNOM fiber probe, whose aperture size was 300 nm. It may
be added that trace amount of Cr2O3 was added to the parent
samples in order to improve fluorescence efficiency of Al2O3,
The further analysis for the position and FWHM of fluorescenceline allowed us to discuss the residual stress distribution andits anisotropy.
Very intense red fluorescence is shown by Cr-doped
Al2O3 (ruby) and its spectrum is widely used as a pressure
scale. The fluorescence spectrum of ruby shows two intense
sharp peaks in the region of 690–700 nm . The peak at thelonger wavelength is called R1; and the other peak at the
shorter wavelength is called R2. Figure 1 shows the twodimensional contoured maps of intensity, position and FWHM
of R1. The peak position was converted to stress values. An
intense fluorescence region inside the circle in Fig. 1(a)indicates that an Al2O3 grain exists inside this circle. On theother hand, a low-intensity region corresponds to the ZrO2

FWHM (nm)Stress (GPa)

(a)

(b)

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Fig. 2 A fluorescence image of Al2O3/cubic-ZrO2 (Y2O3) DSE

grains because the fluorescence intensity of ZrO2 is much
weaker than that of Al2O3. Stress mapping shown in Fig. 1(b)suggests that tensile stress is increasing toward the center ofthis circle, thereby indicating that tensile stress is stronger at
the center of the Al2O3 grain than at the grain boundary
between Al2O3 and ZrO2. Furthermore, Fig. 1(c) indicatesthat the FWHM is larger at the outer part of this circle than theinside. This result suggests that anisotropic stress exists
around the Al2O3 grain and that it is relaxed near the centerof grain. Figure 2 shows the result of the fluorescenceintensity measurements in the Al2O3/cubic-ZrO2 (Y2O3) DSE.
Although an intense fluorescence region again corresponds
to an Al2O3 grain, the Al2O3 grain serves compressive andanisotropic stress and the observed anisotropy was clearlyenhanced at the grain boundary between Al2O3 and cubicZrO2(Y2O3).
References[1] T. Nakagawa, S.Fukura, M.Nakai, K.Sugiyama,R.Kohara
and H.Kagi. Optical Mater., 13, 269(2006).
[2] S.Fukura, H.Kagi, M.Nakai, K.Sugiyama and T.Fukuda.
J.Cryst. Growth, 311, 998(2009).
Key Words
Al2O3/ZrO2 DSE, SNOM, Stress Mapping
Contact toKazumasa Sugiyama (Chemical Physics of Non-Crystalline
Materials Division)(c)E-mail: kazumasa@imr.tohoku.ac.jp

Fig. 1 A fluorescence image of Al2O3/ZrO2 DSE. (a) Peakheight of R1 line; (b) peak position of R1 line(stress); (c)FWHM of the R1 line (stress anisotropy).

4

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IMR KINKEN Research Highlights 2010

Research

Improvement in Fatigue Strength of Ti-29Nb-13Ta-4.6Zr
with Low Young’s Modulus
In order to achieve the improvement in fatigue strengthof Ti-29Nb-13Ta-4.6Zr (TNTZ) with low Young’s modulus, athermomechanical treatment was optimized in this study[1].
The ω phase, which is likely to increase mechanical strength
significantly but decrease ductility drastically in comparison
with the α phase, was employed for precipitation
strengthening. Further, a short time aging treatment was
examined because a small amount of precipitated phasecan suppress the increase in Young’s modulus. As a result,the fatigue strength of TNTZ could be improved through ashort time aging treatment using ω phase precipitation while
keeping relatively low Young’s modulus.
A biomedical β-type titanium alloy, Ti-29Nb-13Ta-4.6Zr(TNTZ), that is composed of non-toxic and allergy-freeelements and exhibits low Young’s modulus of around 60GPa
after solution treatment, has been developed by our group. In
practice, fatigue strength is one of the most important
properties for metallic biomaterials. Therefore, in order toimprove fatigue strength of TNTZ, various thermomechanicaltreatments were conducted previously [2]. According to ourprevious studies, the fatigue strength of TNTZ could be
improved up to the same level as a practical titanium alloy,
Ti-6Al-4V ELI, by a thermomechanical treatment. However,the Young’s modulus of TNTZ was increased above 80GPaby this treatment. In this case, a short time aging is expected
to be one of the effective techniques to improve fatigue
strength and keep low Young’s modulus simultaneously
because the amount of precipitate is smaller than the abovementioned treatment. Further, the ω phase is appropriate asa precipitated phase because this phase increases tensile
strength more greatly than the α phase, although too much
precipitation of this phase drastically decreases elongation.
Therefore, a short time aging at 573K after cold rolling was
conducted to introduce a suitable amount of ω phase
precipitation, and then the effects of the short time aging onYoung’s modulus, tensile and fatigue properties of TNTZ
were investigated in this study.
The forged TNTZ samples were prepared, and then thesamples were subjected to the solution treatment in vacuum
at 1063 K for 3.6 ks, followed by water quenching (ST). Afterthe solution treatment, the samples were subjected to thecold rolling in air at room temperature with the reduction ratioof around 90% (CR). Further, the samples were subjected tothe aging treatment in vacuum at 573 K for various durations
up to 86.4 ks, followed by water quenching (AT).
The Young’s modulus decreases slightly after cold rollingand decreases further after the aging treatment up to 0.6 ks.
After the Young’s modulus decreases once, it increases withaging time. The Young’s modulus of the TNTZ samples aged
up to 10.8 ks can be kept below the value of 80 GPa.
With regard to tensile properties, the tensile strength
increases slightly, but the elongation decreases during coldrolling. Then, the tensile properties do not change much until
an aging time of 1.8 ks. However, the tensile strength starts
to increase over an aging time of 3.6 ks. The tensile strengthand 0.2% proof stress also increase, but the elongation
drastically decreases over an aging time of 10.8 ks. At an
aging time of 86.4 ks, the tensile strength increases, but theelongation becomes almost zero, and, as a result, the 0.2%
proof stress can not be measured.
Figure 1 shows the maximum cyclic stress-fatigue life
curves (the number of cycles to failure), i.e., S-N curves, for
the TNTZ samples aged for 3.6 ks and 10.8 ks. In this figure,the fatigue limits, which are maximum cyclic stresses atwhich the samples do not fail at 107 cycles, of the solutionized
and cold-rolled samples [2] are also shown for comparison.
The fatigue limit of the sample aged for 3.6 ks is similar tothat of the solutionized and cold-rolled samples. However, inthe case of the sample aged for 10.8 ks, the fatigue limit
increases remarkably as compared to that of the solutionized
and cold-rolled samples.
References[1] M. Nakai, T. Akahori, M. Niinomi, H. Tsutsumi and M.
Ogawa: Processing and Fabrication of Advanced MaterialsXVIII, 2, 787-794 (2009).
[2] T. Akahori, M. Niinomi, H. Fukui, M. Ogawa, and H. Toda,
Mater. Sci. Eng. C, 25, 248-254 (2005).
Key Words
Metallic Biomaterial, Thermomechanical Treatment,
Mechanical Properties
Contact to
Mitsuo Niinomi (Biomaterials Science Division)
E-mail: niinomi@imr.tohoku.ac.jp
URL: http://biomat.imr.tohoku.ac.jp/

Fig. 1 Fatigue properties of TNTZ aged at 573 K for 3.6 ks(AT3.6) and 10.8 ks (AT10.8) with fatigue limit of cold-rolled
and solutionized TNTZ (CR [2] and ST [2], respectively).

KINKEN Research Highlights 2010 IMRInfrastructural Materials

5Research

Remarkable Effect of Minor Boron Dopingon the Formation of the Largest Size Ni-rich Bulk Metallic Glasses
We reported the formation of a Ni-rich bulk metallic glasswith a diameter of 15 mm, the largest size known in Ni-richalloys. It was found that the minor addition of boron plays acritical role in the dramatic improvement of glass-forming
ability of Ni-rich Ni-Pd-P alloy. Thermal and chemical
analyses suggest that the improved glass-forming ability is
associated with the decrease of fragility rather than impurity
scavenging [1].
A series of cylindrical Ni60Pd20P20 and Ni60Pd20P17B3
samples with diameters ranging from 6 to 15 mm wereprepared by water-quenching (W.Q.) after a B2O3 flux
treatment. Figure. 1 shows the outer appearances of a φ6
W.Q. Ni60Pd20P20 sample and a φ15 W.Q. Ni60Pd20P17B3
sample. Unlike the φ6 Ni60Pd20P20 sample, the φ15
Ni60Pd20P17B3 sample exhibits a lustrous and smooth surface
without any concaves or cavities caused by crystalline phases.
XRD and TEM characterization further confirms the amorphousnature of the φ15 Ni60Pd20P17B3 sample. The φ15 Ni60Pd20P17B3
BMG was the largest Ni-rich BMG reported so far.
For the remarkable effect of boron in the GFA of Ni-rich
Ni-Pd-P, at first glance, one may consider that the boron
effect comes from oxygen impurity scavenging. To verify thisassumption, the oxygen concentration was quantitatively
measured. The results show that the difference between the
φ15 Ni60Pd20P17B3 sample (11 mass ppm) and the φ6
Ni60Pd20P20 sample (40 mass ppm) is very slight. Thus “selfflux” effect of the added boron in the Ni60Pd20P17B3 alloy ismost likely not the main reason that leads to the dramatic
improvement in GFA, which is consistent with the previous
observation in Pd40Ni40B10P10 alloy [2].
The fragility of the Ni60Pd20P20 and Ni60Pd20P17B3 alloys
were also characterized. From the plots shown in Fig. 2, thefragility parameter (D*) is determined to be 9 for Ni60Pd20P20
and 19 for Ni60Pd20P17B3, respectively. According to theclassification of Angell [4], good glass formers are the stronger
liquid having larger D*. The larger D* is corresponding to

Fig. 2 Fragility plots of the inverse heating rate as a functionof Tg normalized to Tg* measured for glassy Ni60Pd20P20 and
Ni60Pd20P17B3 at a heating rate of 0.067 K/s. The inset shows
the range of D* observed for different materials [3].

higher viscosity and lower diffusivity. If steady-state nucleation
is assumed, the nucleation rate is determined by effective
diffusivity and activation barrier for nucleation at a determined
temperature. The decrease of diffusivity will cause a moresluggish kinetics and result in a lower nucleation and growthrate of crystals which are favorable for the stabilization of thesupercooled liquid region and the improvement of the GFA.
The most likely cause for the liquid behavior change is themuch smaller atomic size of boron. The increased mismatchin atomic size makes it possible to produce denser atomic
packing, which is related to harder diffusion and higher
viscosity in the liquid.
References[1] Y. Zeng, A. Inoue, N. Nishiyama and M. W. Chen, Scripta
Mater. 60, 925 (2009).
[2] Q. Li, D. Greig, S.H. Kilcoyne, P.J. Hine, J.A.D. Matthew
and G. Beamson, Mater. Sci. Eng. A 408, 154 (2005).
[3] L. Shadowspeaker and R. Busch, Appl. Phys. Lett. 85, 13(2004).
[4] C. A. Angell, Sci. 267, 1924 (1995).
Key Words
Glass-forming Ability, Impurity Scavenging, Fragility
Contact toYuqiao Zeng (Non-Equilibrium Materials Division)
E-mail: zyuqiao@imr.tohoku.ac.jp
URL: http://www.nem.imr.tohoku.ac.jp

Fig. 1 Outer appearance of the W.Q. Ni60Pd20P20 (d = 6 mm)
and W.Q. Ni60Pd20P17B3 (d = 15 mm) bulk samples.

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IMR KINKEN Research Highlights 2010

Research

Apatite Formation Ability of Ca-P-O System Bioceramic Coatings
Prepared by Laser Chemical Vapor Deposition
Ca-P-O system bioceramic coatings, e.g., hydroxyapatite(HAp) and α-, β-tricalcium phosphate, were first prepared bymetal organic chemical vapor deposition (MOCVD) andlaser chemical vapor deposition (LCVD). (002) oriented
HAp films prepared by MOCVD and LCVD exhibited thehighest HAp regeneration rate in a pseudo body fluid among
all reported values.
Calcium phosphate ceramic materials, such as α-,β-tricalcium phosphate (Ca3(PO4)2, TCP) and hydroxyapatite(Ca10(PO4)6(OH)2, HAp), are candidate artificial bones andteeth because of their similar structure with human bones
and high HAp regeneration rate. However, the mechanicalstrength and ductility of these materials are too weak forpractical usages. Metallic materials, typically Ti-base alloys,
are also another candidate possessing high strength andmoderate osteoconductivity, whereas the regeneration rateof HAp on Ti–base alloys is insufficient. Therefore, calcium
phosphate coated Ti-based alloy can be promising for thebiomedical applications.
In the present study, Ca-P-O films were prepared on Ti
substrate by MOCVD and LCVD using Ca(dpm)2 and(C6H5O)3PO. Hereafter, the results of LCVD are mainlypresented. A high power Nd: YAG laser with a beam diameterof 15 mm in a continuous wave mode was irradiated to
substrate. The effects of deposition temperature (Tdep), laser
power (PL) and Ca to P molar ratio (RCa/P) on crystal phaseand microstructure of Ca-P-O films were investigated. The
Ca-P-O films were immersed in Hanks’ solution at 37 oC for
6 h to 14 d.
β-TCP films in a single phase were obtained on Ti
substrate at PL = 50 W, Tdep > 900 K and RCa/P < 0.4, while
HAp films in a single phase were obtained at PL = 50 W, Tdep
> 900 K, RCa/P > 1.0 and Tdep < 900 K, RCa/P > 0.45. The
β-TCP films had a (220) orientation with an elongated andangular roof-shaped surface texture, whereas HAp films hada (300) orientation with a granular surface texture [1]. β-TCP
and HAp films had a dense cross section. At PL = 100 W and
Tdep > 950 K, HAp coating in a single phase with (002)orientation was obtained at RCa/P = 0.4 to 0.5. At PL = 200 W
and Tdep = 1100 to 1400 K, α-TCP coating in a single phase
was obtained at RCa/P < 0.6.
Figure 1 summarizes the apatite formation time on various
bioceramic coatings in pseudo body fluid reported in literaturescomparing with those of present study. An apatite phase wasformed on the MOCVD α-TCP coating after 1 day, and covered
the whole specimen surface after 14 d [2]. On the MOCVDHAp coating, an apatite phase was formed after 1 h andcovered the whole surface after 6 h [3]. Particle-shaped
precipitates were observed on the LCVD α- and β- TCP films
after 3 d immersion. However, the precipitate has not wholly
covered the films after 14 d immersion. On the LCVD (300)oriented HAp coating, needle-shaped precipitates were formed
after 12 h and covered the whole surface after 24 h immersion(Fig. 2 left). On the other hand, on the (002) oriented HAp
coating, precipitates were formed after 3 h and wholly covered

KINKEN Research Highlights 2010 IMR

Fig. 1 Apatite formation time on bioceramic coatings preparedby MOCVD and LCVD in a Hanks’ solution.

Fig. 2 Surface microstructure of apatite formed on LCVD
HAp coatings with (300) (left) and (002) orientation (right)
after immersion in Hanks’ solution for 24 and 6 h,
respectively.
the surface after 6 h (Fig. 2 right). (002) oriented HAp preparedby MOCVD and LCVD showed the highest regeneration rateof HAp among literature data reported so far.
References[1] M. Sato, R. Tu, T. Goto, K. Ueda and T. Narushima:
Mater. Trans. 50, 2455(2009).
[2] M. Sato, R. Tu, T. Goto, K. Ueda and T. Narushima: J.
Ceram. Soc. Jpn. 117, 461(2009).
[3] M. Sato, R. Tu, T. Goto, K. Ueda and T. Narushima:
Mater. Sci. Forum 631-632, 193(2010).
Key Words
Bioceramic Coating, Laser CVD, Hydroxyapatite
Contact toTakashi Goto (Multi-functional Materials Science Division)
E-mail: goto@imr.tohoku.ac.jp
URL: http://www.goto.imr.tohoku.ac.jp

Infrastructural Materials

7Research

Novel Grain Refining Technique in Biomedical Co-Cr-Mo Alloy
Using Reverse Transformation
Advanced grain refinement of a biomedical Ni-free Co27Cr-5Mo-0.16N alloy without plastic deformation was
successfully achieved by a reverse transformation fromlamellar (hcp+M2N) phases to fcc phase. Our results willcontribute to the development of biomedical Ni-free Co-CrMo alloy with refined grain size and good mechanicalproperties, without requiring any hot workings.
Currently, improved reliability and safety are required for
hip and knee joint replacements, since these replacements
are being applied in younger patients, greatly prolonging theimplants’ period of service. Grain refinement is one of themost effective methods of achieving higher mechanicalreliability. So far, conventional grain refinement techniques
for medical Co-Cr-Mo alloys have involved thermomechanicaltreatment with recrystallization, such as hot forging, swaging,
and rolling, and do result in greatly improved mechanicalproperties. However, when inhomogeneous strains are
introduced during the hot working process, non-uniformly
distributed grains are produced. Thus, the mechanicalproperties inside the part vary from area to area, owing to the
non-uniform microstructure. In addition, the conventional
techniques are hardly applied to the products fabricated by
precise casting, or sintering. Therefore, the advanced grain
refinement technique without any plastic deformation process
should be considered.
We have advantages for developing advanced grain
refining technique in biomedical Co-Cr-Mo alloys since theirphase transformation behaviors have been revealed
systematicaly in our previous research [1]. Improving themechanical properties of biomedical Co-Cr-Mo alloy can beachieved using reverse transformation; it is convenient andsimple heat treatment.
Figure 1 (a)-(c) show microstructures after solution
treatment (the initial state), aging treatment and reverse
transformation treatment, respectively [2]. The solutiontreated specimen was subjected to aging treatment at 800
oC, forming a lamellar structure of hcp and M2N phases.
Then, the aged specimen having a completely lamellar
microstructure was reverse-treated at 1000 oC, where the fcc
phase is stable. The resultant grains could be refined from200 μm to 20 μm, namely approximately 1/10 of their initialsize. In addition, the obtained grain obtained by reverse
transformation at 1000 oC could be remained constant size
(20-25 μm) regardless of holding time. This is the fact that theresidual fine M2N particles could suppress grain coarsening
by pinning the grain boundaries. Thus, from the viewpoint ofFig. 2 Nominal stress-strain curves of specimens after
solution treatment (ST), aging treatment (AT) and reverse
transformation treatment (RT). The tensile properties ofreverse-transformed specimen and ISO standard (ISO 583212) are shown in the table within the figure.
practical application, this slower grain growth is an attractive
feature for obtaining homogeneous microstructures in
products with complicated shapes or varying thicknesses.
Therefore, this technique is widely applicable to biomedical
products such as hip and/or knee joints, which are produced
by investment casting into various shapes and sizes.
Surprisingly, it was found that the tensile properties of
specimens after reverse transformation were superior tothose required by the ISO standard for biomedical hot-worked
Co-Cr-Mo alloy (ISO 5832-12). Figure 2 show nominal stressstrain curves of specimens after solution treatment (ST),aging treatment (AT), and reverse transformation treatment(RT) [2]. The tensile properties of the reverse transformation
specimen and above-mentioned ISO standard are shown in
the table inserted in the figure. The tensile properties after
reverse transformation exhibited an excellent balance
between strength and elongation (ultimate tensile strength:>
1200MPa, elongation: >30%) compared with the alloy without
reverse transformation treatment.
The technique presented here, based on heat treatment
utilizing reverse transformation, is an effective method forgrain size refining without any deformation. Moreover, thetechnique can be innovated to conventional thermomechanicalprocess since it consists of only heat treatment and hence
we will establish the method for further grain refinement andimprovement of mechanical reliability in biomedical Co-CrMo alloys at the next step.
References[1] S. Kurosu, H. Matsumoto and A. Chiba, Proceeding of
TMS2010 Meeting.
[2] S. Kurosu, H. Matsumoto and A. Chiba, Mater. Let. 64, 49(2010).
Key Words
Grain Refinement, Metallic Biomaterial, Reverse
Transformation

Fig. 1 Microstructure after (a) solution treatment, (b) aging
treatment and (c) reverse transformation treatment.

8

Infrastructural Materials

Contact to
Akihiko Chiba (Deformation Processing Division)
E-mail: a.chiba@imr.tohoku.ac.jp
URL: http://www.chibalab.imr.tohoku.ac.jp/english/index.php

IMR KINKEN Research Highlights 2010

Research

Bias-current Modulation Technique of a Radio-Frequency Glow Discharge Plasma
for Atomic Emission Analysis Associated with a Fast Fourier Transform Analyzer

We have reported that the intensities of particular emission
lines are largely elevated when a dc current, which is driven by the
self-bias voltage, is introduced into the plasma in RF GD-OES [1].
We call this current a dc bias current. The introduction of the dc
bias current could enhance their emission intensities by a factor of
10-20 when atomic emission lines having low excitation energies(3-4 eV) were employed as the analytical line [2]. Furthermore, thedc bias current could be easily pulsated so that a cyclic variation in
the emission signal could be selectively detected by using
modulation spectroscopic methods. Our previous paper describeda detection technique using a lock-in amplifier, which enabled only
the modulated component to be detected selectively [3]. In thispaper, we report on the analytical characteristics of anotherdetection method using an FFT analyzer, indicating that theemission signals can be measured at very low noise levels andthat the analytical precision can be improved greatly [4].
Figure 1 shows the block diagram of the measuring system
employed. A glow discharge lamp was made in our laboratory
according to the original model of Grimm, where the inner diameterof the hollow anode was 8.0 mm. An RF glow discharge plasma
was produced by an RF generator through a matching circuit. A dc
self-bias voltage induced in the RF plasma was separated with an
RF filter and monitored on a digital voltmeter. By connecting a load
resistor with the glow discharge lamp, a dc bias current was
conducted through the overall circuit including the plasma body [1].
The current was adjusted by using a variable resistor and it wasmonitored on a digital ammeter. The bias current could be pulsatedwith a switching circuit using a photoMOS-relay device, whichgave a modulated component of the emission signal at thefrequency of switching. Parameters for the switching: the waveform,
the frequency and the duty ratio, were pre-determined with a digital
function generator. Because the timing pulse is transmitted withthe photo-coupling device, leak of the RF power can be completely
isolated and thus the switching can be performed with very small
distortion. The emission signal was dispersed and detected on ascanning spectrometer, comprising a modified Czerny-Turnermounting monochromator and a photomultiplier tube. An FFT

0.12

Amplitude / arb. unit

A measuring method using a fast Fourier transform (FFT)
analyzer is suggested to estimate the emission intensity from aradio-frequency (RF)-powered glow discharge plasma for atomic
emission analysis. The FFT analyzer has an ability to disperse thecomponents by frequency from an overall signal, and thus works
as a selective detector in modulation spectroscopy. In the RF
glow discharge plasma, a dc bias current can be introduced byconnecting an external electric circuit with the discharge lamp,
which predominantly enhances the emission intensities. Further,the bias current can be pulsated with a switching device tomodulate the emission intensities, and then the modulated
component was selectively detected with the FFT analyzer. Thismethod greatly improved the data precision. The emission
intensity of the Cu I 324.75-nm line in an Fe-based alloy sample
containing 0.043 mass % Cu could be estimated with a relative
standard deviation of 0.20 %. The 3-sigma detection limit of Cu in
Fe-based alloys could be obtained to be 2.3 x 10-6 mass % Cu for
Cu I 324.75 nm.

Data 1
Data 2
Data 3
Data 4

0.10
0.08
0.06
0.04
0.02
0.00

Frequency / Hz
Fig. 2 Reproducibility of the FFT power spectrum in the Cu I
327.45-nm line. Sample: a low-alloyed steel containing 0.043
mass % Cu; rf forward power: 120 W; modulation frequency: 77 Hz(duty ratio of 50%); bias current: 22.6 mA (effective value); thenumber of averaging: 200; data overlapping on calculation: 60%.
analyzer was employed to detect signal components which weremodulated at the frequency of the timing pulse and to remove
noise components having any other frequencies. Based on analgorism of the fast Fourier transform, the FFT analyzer yields a
power spectrum, where the intensity components are separated byfrequency from the overall emission signal. Therefore, only themodulated component can be obtained from the power spectrum.
Because the FFT calculation is very fast, the calculator can use apart of the previous data for estimating the next spectrum in thedata accumulation. Such overlapping of the data enables a power
spectrum to be averaged many times for a short measuring time,
and thus averaging of the power spectrum can contribute to an
accurate estimation of the intensity data in addition to the removalof noise components having any other frequencies.
Figure 2 shows four-fold measurements of a power spectrum
for the emission signal of the Cu I 324.74-nm line when a steel
standard sample containing 0.043 mass % Cu is employed. In thiscase, each spectrum was recorded at a fundamental frequency of
77 Hz and accumulated 200 times at a data overlapping of 60%.
These FFT profiles are almost the same as each other, giving an
average value of 0.098171 with a standard deviation of 0.000195(RSD 0.20 %). This result implies that the FFT detection technique
can lead to a great improvement in the data precision of theemission intensity when the bias current is modulated. Based on
the sensitvity and the signal fluctuation in a blank sample, we could
estimate 3-sigma detection limits of 2.3 x 10-6 mass % Cu for Cu I
324.75 nm and 6.8 x 10-6 mass % Cu for Cu I 327.40 nm.
Furthermore, the background equivalent concentration was
estimated to be 8.4 x 10-4 mass % Cu for Cu I 324.75 nm and 2.6
x 10-3 mass % Cu for Cu I 327.40 nm [4].
References[1] K. Wagatsuma: Spectrochim. Acta, 54B, 527 (1999).
[2] K. Kodama, K. Wagatsuma: Spectrochim. Acta, 59B, 429(2004).
[3] K. Wagatsuma: Anal. Sci., 17, i65 (2001).
[4] K. Wagatsuma, S. Urushibata: Microchem. J., 95, 107 (2010).
Key Words
Modulation Spectroscopy, Atomic Emission Spectrometry, Glow
Discharge Plasma
Contact toKazuaki Wagatsuma (Analytical Science Division)
E-mail: wagatuma@imr.tohoku.ac.jp

Fig. 1 Block diagram of the measuring system.

KINKEN Research Highlights 2010 IMRInfrastructural Materials

9Research

Development of Mass Production System for
Zr-based Bulk Metallic Glasses
In order to realize the constant quality of bulk metallic glasses(BMGs), mass production system of Zr-based BMGs was newly
developed. This mass production system is composed of threemachines; weighing-, master alloying- and casting-machines.
By making full use of developed mass production system,
standardization of BMGs will be accelerated.
Figure 1 shows the outer appearance of automatic weighing
machine, which was designed by the combination of shearing
and weighing processes. The machine weigh 4-metals i.e. Zr,
Cu, Ni and Al, and each metal was cut by shearing to be an
objective weight. In actual operation, since it is difficult to cut the
Zr metal by its own high strength and high ductility, we adjust theweights of other three metals; Cu, Ni and Al to be an objective
alloy composition.

(a)

(b)

Fig. 2 Outer appearance of automatic master alloy fabrication
machine (a)[1], photo of fabricated Zr55Cu30Ni5Al10 masteralloys using the automatic master alloy fabrication machine(b).

Fig. 1 Outer appearance of automatic weighing machine,
right: computer and controller board, left; shearing andweighing machines with turn table for 15 samples, back;materials feeder of three metals, Cu, Ni and Al.
The weight deviation of automatically measure system isabout 1 % in this process, we are trying much higher accuracyof this value to be less than 0.5 %. By utilizing the automaticweighing machine, we can suppress the human error andwide weighing deviation.
Figure 2(a) shows the outer appearance of automatic
master alloy fabrication machine. There was a semi-automaticarc furnace proposed by DIAVAC ltd. [1], the problem to arrest
the development of full-automatic system is the flipping ofmaster alloys. The flipping is an important process tohomogenize the deviated composition in master alloy using
an arc furnace, and we usually repeat the flipping&remelting
process at least 4 times for every master alloys. During thecollaborative research between the YKK co. ltd. and Tohoku
university, we developed the automatic flipping system [2],and we went into production of full-automatic arc-furnace. We
will be able to control the homogenization of master alloys byusing this machine. Before the development of this machine,
the homogenization of master alloy probably depends on the
preparer’s skill. In order to achieve the standardization ofBMGs, we have to keep on the quality of BMGs and we also
standardize the fabrication equipments and those conditions.
To achieve the homogenization of master alloy, we
reduce the total weight of the master alloy less than 30 g. We

10

Infrastructural Materials

can prepare 30 master alloys for one batch, even though thelarge quantity of a master alloy is a common practice formass production. This point is completely different from theprevious mass production process for BMGs. Figure 2(b)shows photo of Zr55Cu30Ni5Al10 master alloys prepared bythis machine. The shining surface as like a mirror points outthat the surface region had already vitrified during thesolidification on water-cooled Cu-hearth after arc melting [3].
The clear vitrification of almost all master alloys implies thehigh reproducibility of developed automatic alloying process.
Recently, we are developing an automatic casting process
as a next step of described above two processes for theachievement of mass production system for standardizationof BMGs. It will be published in elsewhere near future.
References[1] http://www.diavac.co.jp/[2] Japanese Patent Application Number: 2009-118159[3] Y. Yokoyama, H. Fredriksson, H. Yasuda, M. Nishijima
and A. Inoue, Mater. Trans., 48, No6, 1363 (2007).
Key Words
Automatic Weighing Machine, Automatic Master Alloy
Fabrication Machine, Standardization
Contact toYoshihiko Yokoyama (Advanced Research Center of Metallic
Glasses)
E-mail: yy@imr.tohoku.ac.jp

IMR KINKEN Research Highlights 2010