Synthesis of β-FeSi2 with Co or Cu addition by molten salt method
Wen Li
Graduate School of Science and Technology, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8011, Japan
TEL&FAX: +81 53 478 1099
e-mail: f5045024@ipc.shizuoka.ac.jp

Abstract
β-FeSi2 bulk crystals were grown by molten salt method with Co or Cu addition. The effectof Co or Cu addition on structural and thermoelectric property was investigated. The diameterof the domains close to the crystal surface was increased as Cu addition increased, and thevoids are formed with 3.3% Cu addition. As increasing Co addition to 5.0 mol%, α-FexCo1-xSi2
layer was obtained. β-FeSi2 without any impurity addition was performed p-type. On the other
hand, n-type β-FeSi2 crystals were obtained by Co doping with 1.5–4 mol % addition.

1. Introduction
β-FeSi2 is one of the attractive
thermoelectric materials that can be used athigh temperature in air. The silicide growthtechnique using the molten salt method has
been improved to fabricate large sized bulk
β-FeSi2 crystals required for industrial use[1-3]. However, the β-FeSi2 has highresistivity and only the p-type conduction isrealized at this moment. It is important tocontrol the conduction type and the carrier
density of the β-FeSi2. The efforts to obtain ahigh value of the dimensionless figure of
merit, ZT of β-FeSi2 by sintering and otherconventional bulk crystal growth procedures
have been performed by adding Co and otherimpurities [4-7]. However, it is not clear thatelectrical control can be realized by Co
addition.
In this paper, Co addition by molten salt
method to fabricate electrically controlled
β-FeSi2 substrates has been investigated. The
structure property of β-FeSi2 with Cu
addition is also examined.

2. Experiments
β-FeSi2 bulk crystals were prepared by
the molten salt method with Co or Cu
addition using FeSi. The salt is composed of
73.16 mol% NaCl - 21.95 mol% NaF - 4.89
mol% Na2SiF6 [1–3]. The Co or Cu powders
were weighed in the molar ratio of 0-6.0

mol% for the salt mixture. The salt mixture,silicon powder (21.85 mol% for the salt
mixture) and Co (Cu) powders were placedin a SiO2-Al2O3 crucible with the FeSi. They
were thermally treated at a temperature of
900 °C for 120 h in an N2 atmosphere. Afterthe heat treatment, the salt was removed
from the sample using deionized H2O.
The resulting bulk crystals werecharacterized by X-ray diffraction (XRD)
and scanning electron microscopy (SEM).
Seebeck coefficient was measured by using
the temperature differential method at room
temperature.

3. Results and Discussion
Figure 1 shows the cross-sectional SEM
images of β-FeSi2 crystals treated at 900 ˚Cfor 120h with different amount of Cu
addition. The crystals are composed ofcolumnar shaped domains having a diameterof several to 10 µm, which is roughly similar
to the structure of the β-FeSi2 without any
impurity addition. The diameter of thedomain closes to the crystal surface isincreased compared with that of the β-FeSi2
without any impurity addition. It is reported
that the rate of β-FeSi2 formation isenhanced by Cu addition [8], which lead toincreasing the diameter. As increasing Cu
addition to 3.3 mol%, voids are formed. It
has been reported that the Cu atoms are very
mobile and easily move through silicidedomains [9]. As large amount of Cu atoms
are incorporated into Fe-silicide, Cu atoms
diffuse through the Fe-silicide domains into
the crystal and Fe, Si atoms are diffused out[10]. As a result, the voids are formed at thecenter of the crystal.

unintentionally doped impurity. As the Co
addition more than 1.5 mol%, n-typeconduction of appears due to the acceptors
are compensated by Co doping. The small
Seebeck coefficient at Co addition of 5.0
mol% is due to the formation of α-FeSi2
phase. Kamabayashi et al. found inunintentionally doped crystals, a switch of
the sign of the thermopower from p to n withrising silicon content [12].

Fig. 1 Cross-sectional SEM images of β-FeSi2
bulk crystals treated at 900 ˚C for 120 h with Cu
addition of (a) 0.11, (b) 0.33 and (c) 3.3 mol %Figure 2 shows cross-sectional SEM
images of the β-FeSi2 crystals treated at 900
˚C for 120 h with different amount of Co
addition. For the crystals (a) and (b), obvious
difference is not observed between β-FeSi2
without any impurity addition and β-FeSi2
with Co addition less than 5.0 mol%.
However, for the case of Co addition of mol
5.0 %, the silicide layers with the thicknessof 10 m are formed at the surface region of
the crystal, as shown in the Fig.2 (c). The
layer is confirmed to be α-phase by XRD
analysis, and about 1.0 atm % Co atoms areincorporated in the α-FeSi2 region. It isreported that the transition temperature fromβ- to α-phase of Fe1-xCoxSi2 is lowered to850 ˚C [11].

Fig. 1 Seebeck coefficient of the β-FeSi2
crystals with Co addition.

Although the higher Seebeck coefficient
can be obtained, it is difficult to obtain large
sized crystals of β-FeSi2. The β-FeSi2
crystals obtained by molten salt method here
shows equivalent Seebeck coefficient withthat of single crystals. Moreover, thedimension of the crystals is limited by thesize of containers, it is possible to obtainlarger sized crystals by molten salt methodwith Co doping.

4. Conclusion

Fig. 2 Cross-sectional SEM images of the
β-FeSi2 crystals treated at 900 ˚C for 120 h with
Co addition of (a) 1.0, (b) 3.0 and (c) 5.0 mol %.

Figure 3 shows Seebeck coefficient of
the β-FeSi2 crystals with Co addition. It isnoted that the β-FeSi2 crystals grown by themolten salts without any impurity additionshows p-type conduction, due to low Si/Fe
content, formation of domain boundary and

β-FeSi2 bulk crystals were grown bymolten salt method with Co or Cu addition.
It is found that the impurity compensationand n-type conduction appear by Co dopingwith 1.5 – 4 % addition. It is also found that
too much addition of the impurity to the
crystals caused the degradation of thecrystalline quality of the β-FeSi2 crystals,
such as α-phase formation for too much Co
addition. The result leads us to expect thatthe appropriate impurity addition for themolten salt method causes the structuralmodification and electrical control of thegrown crystals.

Acknowledgements
I would like to thank Prof. H.
Tatsuoka, Dr. C.L. Wen, M. Yamashita, T.
Nonomura for their fruitful discussions.
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Fabrication and characterization of modified hydroxyapatite coatings on
Mg-Zn-Ca alloy
Meng Erchao* and Guan Shaokang
Materials Research Center, School of Materials Science and Engineering, Zhengzhou University,
Wenhua Road, No.97, Zhengzhou 450002, P.R. China
TEL&FAX: +81-53-478-1099*E-mail: f5145034@ipc.shizuoka.ac.jp

Abstract:
In this work, dense and uniform fluorine doped hydroxyapatite coating was prepared on thesurface of Mg alloy by pulse electrodeposition. The morphology and chemical composition of
the coating was characterized by SEM and EDS. The corrosion property of the coating wasevaluated by immersion test. And the result indicated that fluorine doped hydroxyapatite was apromising candidate for bio-coating on the surface of Mg alloy.

1. Introduction
Currently used metallic biomaterials
include stainless steels, titanium andcobalt-chromium based alloys. A limitationof these current metallic biomaterials is thepossible release of toxic metallic ions andparticles through corrosion or wear processes
which reduces biocompatibility. And these
metallic biomaterials are essentially neutral
in vivo, remaining as permanent fixtures,such as plates, screws and pins used to secure
serious fractures, which must be removed bya second surgical procedure after the tissue
has healed sufficiently. Repeat surgery
increases costs to the health care system andfurther morbidity to the patient [1]. So it isnecessary to develop safe and biodegradable
metallic biomaterials.
Magnesiumis
anexceptionallylightweight metal. It has similar mechanicalproperties to natural bone, and shows
biocompatibility and biodegradation in thephysiological environment. Simultaneously
in the biodegradation process, the implants ofmagnesium alloy do not release toxic ions.
But the biodegradation rate is high in the
biological environment. Therefore to improvethe corrosion resistance of magnesium alloys,
an effective approach might be theapplication of a surface modified withhydroxyapatite. Some researchs havedemonstrated that pure hydroxyapatite
suffers high dissolution rate in the biologicalenvironment, which is unfavorable for longterm stability of the implants [2]. However,
fluorine doped hydroxyapatite has lower

dissolution rate in biological environment.
Therefore, many researchers focused on theapplication of fluorine doped hydroxyapatite
as bioactive coating.
Fluorine doped hydroxyapatite coatings
have been obtained on other metallic
materials such as titanium [2]. However,reports on the fabrication of fluorine dopedhydroxyapatite coating on the surface ofmagnesium alloy are scarce at present.
Therefore, the purpose of this study is todeposit fluorine doped hydroxyapatite on
magnesium alloy surface by electrochemical
deposition.

2. Experimental procedures
The substrate material was magnesium
alloy with a size of 25mm×10mm×4mm. The
substrates were polished with silicon carbide
papersof
100-1000 grits,cleanedultrasonically in acetone and ethylalcohol
mixed solution until the surface was bright.
Finally, the substrates were activated with
40% HF for 10min and dried using a drying
apparatus.
The electrodeposition was carried out at
65℃ in a two-electrode cell equipped with agraphite plate serving as the counterelectrode. The Mg alloy substrate was usedas the working electrode. An electric heater
was used to maintain the temperature ofelectrolyte, and magnetic agitation was used
to stir the electrolyte at an appropriate
constant speed to keep the concentration
uniform of electrolyte. The average current
density was 1mA/cm2.

The electrolyte was prepared bydissolving 0.15 mol/L NaNO3, 0.025 mol/L
NH4H2PO4 and 0.042 mol/L Ca(NO3)2. For
fabricating fluorine doped hydroxyapatite
coatings, 2 mmol/L NaF was added into theelectrolyte. The pH value of the electrolyte
was adjusted to 5.0 by HNO3 solution and(CH2OH)3CNH2 at room temperature [3].
These reagents were all in analytic grade.
After deposition, the specimens were rinsed
in distilled water and then were dried.

3. Results and discussion
3.1 Effect of fluorine on the surface
characteristics of coatings

rod-like nano-crystals structure, and thediameter of crystal particles reached nearnano grade. It was known that bone wascomposed of nano-crystals of apatite incollagen matrix, therefore, coating material
was closer to bone mineral in crystal
structure, size and morphology, bioactivity ofcoating were better [4].
Trace amount of fluorine element was
detected in fluorine doped hydroxyapatite
coating according to analyses of EDS.
Fluorine played an important role in
influencing the physical and biologicalproperties of the coating. It was known thatfluorine was an essential trace element inbone tissues and teeth, which could promote
the mineralization and crystallization of
calcium phosphate in the new bone forming
process. Therefore, trace fluorine, which was
released in the degeneration process of
coating, was beneficial to the healing of bone
tissue.
3.2 Effect of electrodeposition modes on thesurface characteristics of coatings

Fig. 1 SEM and EDS of coatings: hydroxyapatite(a, b) and fluorine doped hydroxyapatite (c, d).

The coatings of hydroxyapatite coating
and fluorine doped hydroxyapatite coatingwere characterized by SEM and EDS, aspresented in Fig. 1. Hydroxyapatite coating
comprised plate-like crystals, and the size of
crystals was larger than that of the bone
tissue, which is nano-crystal. More important,the distribution of crystals was disordered,
which was unfavorable to the healing of bone
tissue. Fluorine doped hydroxyapatite coating
was obviously different from hydroxyapatite
coating, which took on a closely packed

Fig. 2 SEM of coatings: (a) constant current and(b) pulse current electrodeposition.

SEM images of constant current
electrodeposition
and
pulse
current
electrodeposition coatings were shown in
fig.2. It clearly showed that constant current
electrodeposition coating had defects such as
cracks and pores, and the surface of coatings
was uneven. In biological environment, body
fluid could easily infiltrate into the interfaceof the coating and Mg alloy substrate throughthese cracks and pores. The phenomena not
only induced fast corrosion of the substrates,
but also resulted in undesirable debris andeven delamination, which eventually leaded
to the failure of the implant. It could be seenfrom Fig. 2 (b) that pulse current
electrodeposition coating was dense, uniform
and smooth, and covered the substratecompletely. Thus, it could be supposed thatthe coating prepared by pulse current
electrodeposition could more efficiently
protect Mg alloy substrate from corrosion.
3.3 Immersion testing

crystals of fluorine doped hydroxyapatite
coatingwere
nanograde,
andbetter-organized than hydroxyapatite coating.
The surface of fluorine doped hydroxyapatite
coating prepared by pulse current
electrodeposition was denser and more
uniform than that by constant current
electrodeposition. And immersion testing
demonstrated
that
fluorine
dopedhydroxyapatite coating could effectively
protect Mg alloy substrate from corrosion insimulated body fluid.

Acknowledgements
I would like to thank Guan Shaokang for
his fruitful discussions.

Reference

Fig. 3 SEM of specimens without coating (a) andwith coating (b) immersed in SBF for two weeks.

In order to evaluate the protection
capacity of fluorine doped hydroxyapatite
coating to Mg alloy substrate, immersion
testing was performed. Fig.3 showed that
corrosion had occurred on the surface of thespecimens, but the degree of corrosion attack
was different. It was clear that the surface ofmagnesium alloy substrate had been corroded
seriously, while the specimen with coating
almostretained
surfacemorphology
unchanged. The result indicated that fluorine
dopedhydroxyapatite
coating
could
effectively protect magnesium alloy substratefrom corrosion by simulated body fluid, andcontrolled corrosion rate of the samples,
thereby ensuring long-term stability of theimplants.

4. Conclusion
Fluorine doped hydroxyapatite coating
was successfully prepared on the surface of
Mg alloy. The result indicated that the[1] M. P. Staiger, A. M. Pietak, J. Huadmai
and G. Dias, Biomaterials. 27 (2006)
1728-1734.
[2] J. Wang, Y.L. Chao, Q.B. Wan, Z.M. Zhu,
H.Y. Yu, Acta Biomaterials, 5(2009)
1798-1807.
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Wang, S.J. Zhu, J.H. Hu, C.X. Ren, J.H. Gao,
Y.S. Feng, Appl. Surf. Sci. 257 (2011)
4811-4816.
[4] R. Narayanan, T.Y. Kwon, K.H. Kim,
Mater. Sci. Eng. C28(2008) 1265-1270.
1) Present address: Graduate School of
Science and Technology, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu,
Shizuoka, 432-8011, JapanSynthesis of hexagonal MoSi2 powder by mechanical alloying
Daisuke Ishikawa
Faculty of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0811007@ipc.shizuoika.ac.jp

Abstract
Mechanical alloying (MA) of Mo-Si powders (1 : 2.2 molar ratio) was carried out under
different milling conditions and size of Si powder. The mechanically alloyed powders werecharacterized by X-ray diffraction (XRD), scanning electron microscopy (SEM) and energy
dispersive X-ray spectroscopy(EDS). The preparation conditions of the Mo-Si system by the
MA method are described. The results showed that hexagonal MoSi2 was formed without
impurity.

1. Introduction
Since the melting point is high, thesilicides system metallic compound isapplied as high-temperature structural
material. However, the silicides system
metallic compounds contain large amountsof high melting point active metal. Therefore,
in order that a reaction with a crucible ordissolved material may occur in the synthetic
method by the dissolution, it is difficult toobtain an alloy without impurities.
On the other hand, MA has beenshown to be a technique for producing
homogeneous alloys with controlled
structures or unique properties. MA is a highenergy ball milling process consisting ofrepeated mechanical mixing, fracturing cold
welding, and rewelding of urtrafine alloy
powder during the collision of the ball andthe powder [1]. MA allows materials
scientist to circumvent manufacture alloyand material limitations that are impossible
or difficult to produce by casting techniquesand conventional melting.
The MA of molybdenum and silicon has
been studied drastically with mostresearchers reporting a naturally-arising hightemperature synthesis reaction after only a
few hours of milling to form the tetragonal
C11b structure (t-MoSi2). Continued highenergy deformation induced by milling has
also been shown to fovour the formation of
the hexagonal C40 polymorph (h-MoSi2) [2].
In this study, the result of synthesis of

h-MoSi2 powder by mechanical alloying is
described.

2. Experiments
The
Mo-silicide
powders
weresynthesized by MA method. As the starting
materials, Mo (99.97%up, 0.7 µm) powders
and #200 mesh Si powder (74µm,99.9999% ) or Si powder (0.05µm 0.1µm, 99.8%) were used. The Mo and Si
powders were weighed in the molar ratio of
Mo : Si =1:2.2. The powders and steel
milling balls were placed in a steel vial
under an argon atmosphere, then the vial wassealed by an O-ring. The milling was carriedout using a SPEX 8000 vibratory mill for
50h. Details of milling conditions aredescribed in Table1. The powders structuralproperties were characterized by X-ray
diffraction (XRD) measurements using a Cu
Kα1 source, scanning electron microscopy(SEM) along with energy dispersion
spectroscopy (EDS).

Table 1 MA conditions
MA parameters
Milling balls
(A)
(B)
(C)

Φ5/32×16
Φ5/32×16
Φ1/4 ×10

3. Results and Discussion

Figure 1 XRD patterns of the Mo and Si
powder mixture synthesized by MA under thecondition shown in Table 1.

Figure 1 shows the X-ray diffraction
patterns for the mechanical alloying of Mo
and Si under milling intensities (A), (B), and(C), respectively as defined in Table 1. It is
confirmed that the MoSi2 and Mo5Si3 phases
are formed. Figure 1 (A), (B) and (C) shows
strong and broad h-MoSi2 peaks, strong Mo
peaks and weak MoSi2 and Mo5Si3 peaks,
respectively. The intensity of Mo-silicides
peaks of (B) greatly reduce compare to thatof (C). The intensity of h-MoSi2 of (a)greatly enlarge compare to that of (B) and(C). It was found that Mo-silicides weremore formed by using bigger milling balls
and Si-rich compounds were more formedby using bigger particle size of Si powder
than that of Mo powder.
Fig.2 shows the SEM micrographs of thepowders milled under condition A. The
average grain size is about 1 µm, which areabout the same as the original Mo powder
size. However it was found that the larger Si
powders have been crushed and are notobserved in the images. The table shown in
Figure 2 shows components analysis. It isobserved not only Mo and Si but also Cr, Fe,
C and O, however impurities and oxidation
which arose from mechanical alloying

The particle sizeof Mo powder[µm]
0.7
0.7
0.7

The particle sizeof Si powder[µm]740.05~0.1
0.05~0.1

process are deleted from table shown infigure 2. In addition, the components
analysis suggest the composition of thepowders in the images of Mo : Si ≈ 1 : 2.

1
2
3
4

Mo
34.66
26.80
31.85
31.75

Si
65.34
73.20
68.15
68.25

Figure 2 The SEM micrograph and thecorresponding EDS analysis for the resulting
MoSi2 powders synthesized under condition A
shown in Table 1.

4. Conclusion
Mo-silicides were synthesized bymechanical alloying under three conditions.
It was found that Mo-silicides were moreformed by using bigger milling balls andSi-rich compound were more formed byusing bigger particle size of Si powder thanthat of Mo powder. It was observed thecomposition of the powders in the SEM
images of Mo : Si ≈ 1 : 2 under condition A.

Acknowledgements
I would like to thank T.Nonomura and K.
Shirai and K.Nakane and H. Tatsuoka for
their fruitful discussions.

Reference[1]

Pee-Yew Lee, Tzeng-Rong Chen,
Ju-Lung Yang and Tsung-Shune Chin,
Mater. Sci. Eng 192-193(1995) 556.

[2] A.J. Heron and G.B. Schaffer, Mater. Sci.
Eng. A352(2003) 105.

Synthesis of Ca2Si Powder by Mechanical Alloying
Shingo Oda
Graduate School of Engineering, Shizuoka University
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0130098@ipc.shizuoika.ac.jp

Abstract
Powder of Ca-silicide compounds were grown by mechanical alloying (MA). Ca-silicidecompounds powder were grown in all conditions. The effect of milling energy was investigated.
Some of Ca and Si powders were not reacted by MA using 5/32 inch × 16 stainless steel
milling balls.

1. Introduction
Recently, semiconducting silicides haveattracted much attention for their potential tocreate new classes of environmentallyconscious
electronics[1].
Because,semiconducting silicides is non- toxic andabundant materials. Alkaline-earth metal(AEM) silicides are categorized in thismaterials group. The growth of AEM
silicides and their applications to solar celland thermoelectric materials have been
investigated [2]. AEM silicides based
compounds, such as Mg2Si1-xGex [3] andMg2Si1-xSnx [4], could be promisingthermoelectric semiconductors in the middle
temperature range. However, it is generally
difficult to obtain highly efficient
thermoelectric p-type semiconductors.
Ca2Si is one of the AEM silicides, andthe electronic structures of Ca2Si have been
calculated with the expectation of a directtransition [5, 6].
It was also reported that theorthorhombic Ca2Si is a stable phase with adirect energy gap of about 0.35 eV [7].
However, its semiconducting property has
not yet been experimentally determined [8].
Matsui et al. grew Ca2Si layers by vapor
phase growth (VPG) method using Mg2Si/Si
substrates [9].
Mechanical alloying (MA) is a powder
processing technique that allows productionof homogeneous materials starting fromblended elemental powder mixtures. MA
method is simple compared with VPG
method. It is also possible to synthesize

metastable phases with an unknown structure
by MA technique [10]. It is expected that the
MA method combined with the recent
silicide technology would provide further
development of new semiconducting silicide
technologies.
In this paper, it was attempted thatCa-silicides were synthesized by MA. In
addition,resultantpowders
werecharacterized.

2. Experiments
Ca and Si powders were synthesized bymechanical alloying. As the starting
materials, fine Si powder (0.07-0.10
micrometer APS, 98%) and Ca granules (-16mesh, 99.5%) were used. Ca and Si powders
were weighed in the molar ratio of Ca:Si =7:3. The powders and 1/8-1/4 stainless steel
milling balls were placed in a steel vial
under an argon atmosphere. Then, the vial
was sealed by an O-ring. The milling was
carried out in a SPEX 8000 vibratory mill.
Detail of milling conditions are shown in
Table 1.
The powders structure properties werecharacterized by X-ray diffraction (XRD)
measurements using a Cu Kα1 source andscanning electron microscopy (SEM) alongwith energy dispersion spectroscopy (EDS).

3. Results and Discussion
Figure 1 shows X-ray diffraction patterns

Table 1 MA conditions

MA parameters Milling balls The particle size

The particle sizeof Ca powder[mm] of Si powder[µm](a)Φ1/4×4~1.0

0.07~0.1(b)Φ1/8×16~1.0

0.07~0.1(c)Φ5/32 ×16

~1.0

0.07~0.1

Fig. 1 XRD patterns of the Ca and Si
powders mechanically alloyed under thecondition shown in Table 1. respectively.

Fig. 2 the SEM micrograph and thecorresponding EDS analysis for the Ca and Si
powders mechanically alloyed under condition
shown in Table 1.

of the powders mechanically alloyed of Ca
and Si powders under milling conditions of(a), (b) and (c), respectively defined in Table
1. The diffraction peaks of Ca-silicides andSi were observed for the powder synthesizedunder the all milling conditions. Moreover,the diffraction peak of Ca was observed forthe powder synthesized under the milling
condition of (b). Some of Ca and Si powders
were not reacted because milling conditionof (b) was the lowest energy in all conditions.
However, most of Ca powders were oxidized
and then CaO were changed into Ca(OH)2 in
the air. The diffraction peaks of Ca2Si withthe cubic structure were observed undercondition (b). However, PDF data of Ca2Si
with cubic structure is not determined.
Hence, the cubic structure of Ca2Si isdiscussed.
Figure 2 shows SEM images and thecorresponding EDS analyses of the powder
synthesized under the all milling conditions.
The composition ratio of Ca and O wereabout 1:1.6 in all milling conditions.
According to XRD spectra and EDS
analyses, it was confirmed that Ca powders
were changed into Ca-silicides and Ca(OH)2.
Fe and/or Cr atoms were existed in allconditions. It was considered that this result
was caused by stainless milling balls. It isobserved that Ca atoms were homogeneously
distributed corresponding to the Si. In otherwords, it was considered that Ca-silicides, Si
and
Ca(OH)2
were
homogeneously
distributed.

4. Conclusion
Powder of Ca-silicide compounds weregrown by MA. The diffraction peaks of
Ca-silicides, Ca(OH)2 and Si were observedin all milling conditions. The diffraction
peak of Ca was also observed for the powder
synthesized under the milling condition of(b). Some of Ca and Si powders were notreacted for the powder synthesized under the
milling condition of (b). Ca atoms were
homogeneously distributed corresponding to

the Si and the composition ratio of Ca and O
were about 1:1.6 for the powder synthesizedunder the all milling conditions.

Acknowledgements
I would like to thank T. Nonomura and D.
Ishikawa and H. Tatsuoka for their fruitful
discussions.

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Thermoelectrics (2006) 535.
[9] H. Matsui, M. Kuramoto, T. Ono, H.
Nose, H. Tatsuoka, H. Kuwabara, J.
Cryst. Growth 237/239 (2002) 2121.
[10] Y. Warashina, Y. Ito, T. Nakamura, H.
Tatsuoka, J. Snyder, M. Tanaka, T.
Suemasu, Y. Anma, M. Shimomura, Y.
Hayakawa, e-J. Surf. Sci. Nanotech. 7(2009) 129.

Synthesis and characterization of NaSi
Kengo Shirai
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0030124@ipc.shizuoika.ac.jp

Abstract
The single phase NaSi was obtained by grinding the smaller grains of the sample prepared byheating Na and Si at 700ºC for 24 h. Two layers were formed on the surface of the material
they are sodium oxide, sodium silicate and NaSi from the surface sequentially, which werecharacterized by XPS. From XPS results, semiconductor-like behavior of NaSi is shown.

1. Introduction

2. Experiments

Silicide semiconductors have attracted
much attention as one of the environmental
conscious semiconductors. There are somephases including clathrate structure in Na Si system [1]. Some phase has beencharacterized as air and moisture sensitive[2].
This study showed that the clatherate
structure of Na8Si46 is non-reactive with air
and moisture. However, it is agreed that thereaction of NaSi with water is rapid.
Recently, first-principle calculation showedthat alkali metal monosilicides, for example,
NaSi (1.24 eV), have larger band gaps
compared with other conventional silicides,
such as β-FeSi2 (0.85 eV), MnSi1.7 (0.68
eV) and Mg2Si (0.78 eV) [3-7].
However, the detailed properties of alkali
metal silicides have not been clarified. NaSi
crystals were prepared from a Na-Si melt byMorito et al. [8]. However, complete phase
diagram has not been clarified yet [8].
Instability of NaSi in the air degrades its
crystalline quality itself, and makes it
difficult to characterize the intrinsic propertyof the NaSi.
In this paper, the surface states of thegrown NaSi were clarified, and thestabilization of the NaSi surface was
examined.

The NaSi were prepared with starting
materials of Si grains (High Purity
Chemicals, 99.999% purity) and Na lumps(High Purity Chemicals, 99% purity), whichwere weighed in the molar ratio of
Na:Si=1.2:1.0 in an Argon gas-filled glove
box and charged in boron nitride (BN)crucibles. The crucible was sealed in astainless-steel tube to prevent thevaporization of Na and the reaction with H20of the silicides during the heat treatment.
The sealed tube was heated at 700 oC for 24
h in an electric furnace.
The structural properties of the resultant
compounds were characterized using X-ray
diffraction measurement (XRD), scanningelectron microscopy (SEM) with an energy
dispersion spectroscopy (EDS) analyzer. In
addition, the surface condition wascharacterized by X-ray photoelectron
spectroscopy (XPS) using a VG, ESCA-LAB
Mk2 with a nonmonochromatized Al Kαsource (hν=1486.6 eV).

3. Results and Discussion
Figure 1 shows XRD spectra of (a) the
resultant powder prepared at 700 ºC (b)NaSi PDF -card data [9]. As shown in
the figure the spectrum of (a) agrees withthe powder diffraction file (PDF) of NaSi,
which shows that single phase NaSi isformed without any predominant
byproducts.

Fig. 1 XRD spectra of (a) the resultant powder prepared at 700 ºC and (b) NaSi PDF -card data (PDF
card No.01-089-2625).

Figure 2 shows SEM image and EDS
spectrum of NaSi grains. As shown SEM
image, faceted surfaces were observed on thegrain, which shows the grain growth of thecrystals. In addition, the surface change wasobserved. Based on the results of EDS
analysis, it is observed that just O, Na, andSi elements are detected on the surface. On
many regions, the ratio of Na to Si atoms isaround 1: 1. On other regions, More O
elements are detected. Those regions are
regarded as sodium oxide and sodium
silicate.

Fig. 2 SEM image and EDS spectrum of
NaSi grains.

Figure 3 shows XPS spectra related to
Si2p , Na1s and NaKLL of the NaSi. Sample
was etched by Ar+ ion beam sputtering from0 min to 150 min. The binding energy shift
originated from a charge-up was correctedby referring to the C1s peaks ofhydrocarbons at 285.2 eV. From results, it
was found that the binding energy of both
Si 2p and NaKLL peaks change withincreasing sputtering time. The peak of
150min deserves attention in particular. It isshown that two spectra of Si2p core-level
ware detected. The higher was related tosodium silicate. Another one was not sodium
silicate, which was lower than elemental Si.
This phenomenon suggests that a Si atom
receives an electron from neighboring atoms.
In addition , It is shown that the Auger
energy difference between those of oxideand metallic Na compounds is about 4.5 eV[10][11][12]which agree the experimental
results in Figure 3 (d). Furthermore, theplasmon intensity at the lower energy side of
the peak is not observed. From the results
mentioned above, it is suggested that Si and
Na atoms are bound each other and thecompoundconsistingthemaresemiconducting. Also, this result shows thatthe sample surface consists of three layers. It
is sodium oxide, sodium silicate and NaSi
from the surface sequentially.s formed
without any predominant byproducts. On theother hand, for the spectrum (b), thediffraction peaks due to Si are overlapped inaddition to a series of that of NaSi. It isconsidered that for larger sized grain, thereaction to form NaSi is not completed, andSi as a one of the starting materials remains
in the grains. However, it is certain from theFig. 3 XPS spectra of Si2p, Na1s and NaKLL before stabilization was etched by Ar+ ion beam
sputtering from 0min to 150min.

spectra that NaSi is the dominant product,even though that the residual Si is observedin the silicides.

4. Conclusion
The powder of the single phase NaSi was
obtained by grinding the smaller grains of
the sample prepared by heating Na and Si at
700ºC for 24 h. The results of XPS
measurement is suggested that Si and Na
atoms are bound each other and thecompoundconsisting
of
them
is
semiconducting. Moreover, compounds
surface consists of three layers. They aresodium oxide, sodium silicate and NaSi fromthe surface sequentially.

Acknowledgements
I would like to thank M. Shimomura, K.
Nakane and H. Tatsuoka for their fruitful
discussions.

Reference[1] J. Sangster and A.D. Pelton, J. Phase
Equilibr. 13, 67 (1992).
[2] He, D.D. Klug, K. Uehara, K.F. Preston,
C.I. Ratcliffe and J.S. Tse. J. Phys. Chem.
B, 105,3465 (2001).
[3] Y. Imai and A. Watanabe, J. Alloys
Compd. 478, 754 (2009).
[4] N.E. Christensen. Phys. Rev. B, 42, 7148
(1990).
[5] M.C.Bost and J.E.mahan, Electro.Mater.
16, 389 (1987).
[6] G. Busch and U. Winkler, Helv.Phys.Acta.
26, 359 (1953).
[7] G. Busch and U. Winkler, Physica. 20,1067 (1953).
[8] H. Morito, T. Yamada, T. Ikeda and H.
Yamane, J. Alloys Compd. 480, 723(2009).
[9] J. Witte, H. G. von Schnering, Z. Anorg.
Allg. Chem. 327, 260 (1964).
[10] J. F. Moulder, W. F. Stickle, P. E. Sobol,
K. D. Bomben, Handbook of X-ray

Photoelectron.
Spectroscopy,(Perkin-Elmer, Minnesota, 1992)[11] A. Barrie and F. J. Street, J. Electron.
Spectrosc. Relat. Phenom. 7, 1 (1975)[12] C.D. Wagner, W.M. Riggs, L.E. Davis,
J.F. Moulder and G.E. Mullerberg,
Handbook of X-ray Photoelectron
Spectroscopy, (Perkin-Elmer, Phys.
Electron. Div., Eden Prairie, MN 55344,
1979)

Synthesis of Tungsten Oxide Nano-Particles by Thermal Oxidation Technique
Yuki Usuda
Faculty of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka, 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0811014@ipc.shizuoka.ac.jp

Abstract
The tungsten trioxide nano-particles on tungsten wire were grown by the thermal oxidationin the air. The dependence temperature of structural properties of the nano-particles isinvestigated. The WO2 nano-particles with the diameter less than 300 nm are formed whengrown at 500 °C with the heat treatment time of 4 hours. The WO3 with certain thickness andhigh density are formed as the reaction temperature increased to 750 °C. This simple growth
procedure could be used to synthesize various metal oxide nanostructures.
1. Introduction
Transition metal oxides have various
interesting properties such as colossal
magnetoresistance, superconductivity, andpiezoelectricity. Since nanosized transition
metal oxides have huge surface areas, they
are promising as a material for a vast range ofapplications including lithium-ion batteries,
catalysts, and electro-chromic materials [1].
WOx is an excellent material forelectro-chromic or photo-chromic materials
and could be used as fluorescent material of
X-ray screen, semiconductor gas sensors, andmanufacture of smart or electro-chromic
window. Moreover, the WO3 powders areused for pigment ink. Studies on nanoscaled
WOx, however, are relatively new andeffective preparation methods for suchmaterials are not yet reported.
This paper aim is to exemplify a simple
synthesis of WO3 nano-particles throughdirect thermal oxidation in the air. The
growth condition dependence of thenano-particles is shown on the dynamic
evolution of WO3 nano-particles.

2. Experiment
WO3 nano-particles were grown by thethermal oxidization technique of tungsten
wire (Φ 1.0 mm × 10 mm). The tungsten
wires were loaded into the quartz tube which
is open to the air. The nano-particles

synthesis was performed by exposure of thetungsten wire to the air, and thenano-particles were synthesized for 4 h at thetemperature of 500 °C and 750 °C,
respectively.
The structure properties of nano-particles
were characterized by field emission
scanning electron microscopy (FE-SEM,
JSM-7001F) and X-ray diffraction (XRD,
RINT Ultima III).

3. Results and Discussion
Figure 1 shows the SEM images of thetungsten wire surface grown at 500 °C withthe heat treatment time of 4 h. It is observedthat the feature of the wire surface roughly
conform to the unreacted wire surface. A few
nano-particles with the diameter less than
300 nm are distributed on the surface.
Figure 2 shows the SEM images of across-section and the surface of tungsten wire
treated at 750 °C with the heat treatment timeof 4 h. The layers are formed around the wire,as shown in Fig. 2 (a). In Fig.2 (b), the layer
consist of nano-particles with high density,
while the diameter of the nano-particles isincreased compared with that of thenano-particles synthesized at 500 °C.
Figure 3 (a) and (b) shows the X-ray
diffraction spectra of tungsten wires treatedat 500 ºC and 750 ºC, respectively. It is
confirmed that the nano-particles synthesized
at 500 ºC are composed of WO2. As shown in
Fig. 3(b), the peaks can be completely

indicated to diffraction of WO3, which
reveals that the nano-particles with highdensity consist of WO3.

Fig. 1 SEM images of the tungsten wire grown at 500 ºC with the heat treatment time of 4 h.

Fig. 2 SEM images of (a) cross-section and (b) surface of the tungsten wire grown at 750 ºC
with the heat treatment time of 4 h.

Fig. 3 XRD spectra of the tungsten wires treated at (a) 500 ºC and (b) 750 ºC.

4. Conclusion
The tungsten trioxide nano-particles ontungsten wire were synthesized by thethermal oxidation technique in the air. WO3
nano-particles and high density aresynthesized at 750 ºC with the heat treatment
time of 4h. The experimental temperature
affects the structural properties of thenano-particles. This simple growth procedureencourages us to synthesize various metaloxidesnanostructures
using
the
compositional conversion of the starting
materials.

Acknowledgments
I would like to thank H. Tatsuoka, W. Li,
S. Cai, T. Nonomura, T. Matsushita, and H.
Sato for their fruitful discussions.

Reference[1] Gang Gu, et al. Nano Lett. 2(2002) 849.

Growth of Ti-Oxide Nanorods by Thermal Oxidation Technique
Tomoji Matsushita
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0130151@ipc.shizuoika.ac.jp

Abstract
The titanium oxide nanorods on titanium wire were grown by the thermal oxidation
technique in the air. The growth time dependence of nanorods was investigated. The length andthe diameter of nanorods grown at 800 ºC for 4 hours were approximately 840 nm and 160 nm,
respectively. Layer and nanorods were composed of Ti-oxide such as TiO 2 rutile.

It isclarified that the length and diameter of nanorods depend on growth time.

1.

Introductionhydrothermal

process[10],

electrochemical oxidation method [6] andTiO2 is an n-type semiconductor withthermal oxidation technique [11]. Amongwide band gap energy of 3.0-3.2 eV and hasof them, thermal oxidation technique isbeen widely used for various applications in

simple and cheap. TiO2 nanostructures

dye-sensitized solar cells [1], gas sensors [2],grown by this technique was carried out

photocatalyst [3] and negative electrode forunder carrier gas such as acetone [7] andlithium rechargeable batteries [4].

ethanol [11]. However, almost none of thisNanostructured materials have attracted
much attention for its promising potential forelectronics,

optoelectronics

technique is able to grow the TiO 2nanostructures in the static air.

applications

In this study, we reported the growth of

because they have many Merits such as a

Ti-oxide nanorods by thermal oxidationhigh surface to volume ratio and shorttechnique of titanium substrates in theelectron pathway [5]. They have beenstatic air. In addition, the growth timereported in various forms such as nanotubes

dependence[6], nanowires [7] and nanoparticles [8].

investigated.

TiO2 nanostructures are synthesized bymany techniques such as sol-gel method [9],

of

the

nanorods

was2.

Experiments

respectively. Plate-like crystals in fig. 1(d)were formed larger than that in fig. 1(b).

The TiO2 nanorods were grown by

Figure 2 shows XRD spectra of TiO2

thermal oxidation technique of titanium

nanostructures grown for 0.5 and 4 hours

substrates in the static air. Titanium wires

at the temperature of 800 ºC. TiO 2 exist inwith diameter of 1 mm were used as thethree forms of tetragonal rutile, tetragonalsource material. Each of the substrates wasanatase and orthorhombic brookite. In fig.

polished with #400 and #1500 sandpapers,

2, some diffraction peaks of TiO2 rutile

then ultrasonic cleaned in acetone for 10

were identified. Overlapped diffraction

min. Substrates were placed at the center ofpeaks due to other TiO2 structures and Ti-

electric furnace. The temperatures wereoxide materials such as TiO, Ti 2O3 andraised to 800-900 ºC. Substrates wereTi3O5 were observed. Hence, layer andannealed in air for 0.5-4 hour.

Thenanorods in fig. 1(a)-(d) were composed

temperatures

to

room

of Ti-oxide such as TiO2 rutile.

nanorods

were

were

reducedtemperature in atmosphere.
The

Ti-oxidecharacterized by field emission scanningelectron microscopy (FE-SEM, JSM-7001F),
and X-ray diffraction (XRD, RINT Ultima
III).

3.

Results and Discussion

Figure 1 shows SEM images of TiO2
nanostructures grown at the temperatureof 800 ºC for 0.5 and 4 hours. In Fig. 1(a)and (b), nanorods were not grown, butlayer was formed on the substrate at thetemperature of 800 ºC for 0.5 hour. Layer
was consisted of plate-like crystals. In Fig.
1(c) and (d), nanorods were formed andlaid randomly on layer at the temperatureof 800 ºC for 4 hours. In Fig. 1(d), thelength and the diameter of nanorods wereapproximately 840 nm and 160 nm,

Fig. 4 XRD spectra of Ti-oxide layer andnanorods at 900 ºC for 4 h (a) and 1 h (b). R
represents TiO2 rutile structure.

710 nm and 180 nm, respectively. The
length and diameter of nanorods in Fig.
3(d) were formed the larger than that in
Figure 3(b). According to fig. 3(b) and (d),the length and diameter of nanorods
depend on growth time.
Figure 4 shows XRD spectra of TiO2
Fig. 1 (a) Low and (b) high magnification FESEM image of Ti-oxide layer and nanorods at
800 ºC for 0.5 h and (c) low and (d) highmagnification FE-SEM image at 800 ºC for 4h.

Figure 3 shows SEM images of TiO2
nanorods grown at the temperature of 900 ºC
for 1 and 4 hours. In Fig. 3, nanorods wereformed and laid randomly on layer at thetemperature of 900 ºC. In Fig. 3(b), thelength and the diameter of nanorods wereapproximately

490

nm

and

70

nm,

Fig. 3 (a) Low and (b) high magnification FESEM image of Ti-oxide layer and nanorods at
900 ºC for 1 h and (c) low and (d) highmagnification FE-SEM image at 900 ºC for 4h.

nanostructures grown at the temperatureof 900 ºC for 1 and 4 hours. In fig. 4,
some diffraction peaks of TiO2 rutile werealso identified. Overlapped diffraction
peaks due to other TiO2 structures and Tioxide materials such as TiO, Ti 2O3 andTi3O5 were observed. Hence, layer andnanorods in fig. 3(a)-(d) were composedof Ti-oxide such as TiO2 rutile.
Fig. 2 XRD spectra of Ti-oxide layer andnanorods at 800 ºC for 4 h (a) and 0.5 h (b). R
represents TiO2 rutile structure.

4.

Conclusions

respectively. In Fig. 3(d), the length and theTi-oxide nanorods were grown by

diameter of nanorods were approximately

thermal oxidation technique in the staticair. Nanorods were grown at the temperature

178 (2005) 2179.

of 800 ºC for 4 hour and at the temperature[11] S. Daothong et al. Scripta Mater. 57

of 900 ºC for 1-4 hours. Layer and nanorods(2007) 567.

were composed of Ti-oxide such as TiO 2,
TiO, Ti2O3 and Ti3O5. The size of nanorods
and layer was dependent of the growth time.

Acknowledgements
I would like to thank S. Cai and S. Oda
and Y. Usuda and S. Sato and H. Tatsuoka
for their fruitful discussions.

Reference[1] S. Dai et al. Sol. Energy Mater. Sol.
Cells 84 (2004) 125.
[2] L. Francioso et al. Sens. Actuat. B 130(2008) 70.
[3] M. Muruganandham et al. Sol. Energy
Mater. Sol. Cells 81 (2004) 439.
[4] C. Natarajan et al. Electrochim. Acta.
43 (1998) 3371.
[5] Y. Xia et al. Adv. Mater. 15 (2003)
353.
[6] X. Zeng et al. J. Alloys Compd. 509(2011) L221.
[7] K. Huo et al. J. Nanosci. Nanotechnol.
9 (2009) 3341.
[8] H.-D. Jang et al. Ultramicroscopy 108(2008) 1241.
[9] Lei et al. Appl. Phys. Lett. 78 (2001)
1125.
[10] R. Yoshida et al. J. Solid State Chem.

Synthesis of Sodium – Tin system composites by Thermal Annealing
Kaito Nakane
Faculty of Engineering, Shizuoka University
TEL&FAX : 053-478-1099
E-mail : f0811091@ipc.shizuoka.ac.jp

Abstract
Nanostructured sodium – tin system composites were synthesized by thermal annealing.
These shapes of nanostructures grown on surface of composites were different each other. The
structural properties of the resultant composites were characterized by scanning electron
microscopy (SEM), and X-ray diffraction (XRD).

1. Introduction
Recently, a tin and tin compound attracts
much attention in various fields. It wasreported that tin catalyst was useful forsynthesis silicon nanowires [1]. In addition,
it was reported that tin oxide nanorods wereused as a Li storage compound for Li-ion
batteries [2]. Furthermore, it is reported that
amorphous zinc tin oxide thin films were
examined for thin film transistor application[3].
On the other hand, the intermetallic
compounds of the alkali metals withgroup-IV B elements (Si, Ge, Sn, Pb) from aninteresting class of materials, exhibiting quite
peculiar crystal structures associated with avariety of chemical-bond types [4]. Amongthem, Na – Sn phase system has intensively
been investigated for more than 80 years [5],the discovery of binary phases and theirstructural characterization occurred only
recently. After the structural characterizationof NaSn in 1977 [6] it took 20 years until thestructures of other binary phases on the
tin-rich side were detected and characterized,
namely Na5Sn13(1997) [7], NaSn5(1998) [8],Na7Sn12(2003) [9], and NaSn2(2005) [10].
In this study, Na-Sn system composites
were synthesized by simple thermal
annealing method, and their structure
property is characterized.

2. Experimental
Sn

The composites were synthesized frompowder (Alfa Aesar, -100 mesh,99.999%) and Na lumps (High Purity
Materials, 99%), which were weighted in an
Ar filled glove bag, and charged in boron
nitride (BN) crucibles (DENKA; inside
diameter, 6 mm; depth, 17 mm). The detailsof experimental conditions are described inTable 1. Each crucible was sealed in astainless-steel tube (inside diameter, 13mm;
length, 80 mm) to prevent the vaporizationof Na. The sealed tubes were heated at
600ºC for 24 hours in an electric furnace.
The structural properties of the resultant
composites were characterized by scanningelectron microscopy (SEM), and X-ray
diffraction (XRD).
Table 1 Preparation conditions

Composite#1#2

Molar ratio(Na : Sn)
1.0 : 1.1
1.0 : 2.0

3. Results and Discussion
Figure 1 shows the XRD spectra of (a)composite #1 and (b) composite #2 shown in
Table1. As shown in Fig.1 (a) and (b), it isconsidered that Na15Sn4, Na2(Sn(OH)6) andNa2CO3-H2O were synthesized respectively.
By comparing diffraction peaks of bothcomposites, it shows that the tin of
non-reaction exists in both composites.
Figure 2 shows the SEM image ofComposite #1. As shown in Fig.2 (a), a large
number of wires were formed on thecomposite surface. Furthermore, it isinteresting that the wires tend to grow up

Fig.2 SEM images of Composite #1 shown in
Table1.

Fig.1 XRD spectra of (a) the composite #1 and(b) the composite #2.

Fig.3 SEM images of Composite #2 shown in
Table1.

from a lumps formed on the composite
surface, which is illustrated by the image in
Fig.2(b). Additionally, it is observed that anindividual wire is thin, straight, and has a
sharp top end. Meanwhile, it is found thatthe average diameter of the wires is 400 nm
and their lengths are in the range of about
6.7-11µm.
Figure 3 shows the SEM images ofComposite #2. As shown in Fig.3 (a), a large
number of rods were formed on thecomposite surface. The average diameter ofthe rods is 1.76 µm and their lengths are in
the range of about 7.7-10µm. Additionally, it
is observed that an individual rod has an
unhomogeneous shape and has a rugged
surface, which is illustrated by the image in
Fig.3 (b).
Though composite #2 was synthesizedunder the Sn-rich condition, the peaks of tinof composite #2 is weaker than that ofcomposite #1 from the XRD spectra.
However, the XRD spectrum of composite#2 shows surface. On the other hand, the
XRD measurement of composite #1 wascarried out powder of composite #1.
According to XRD spectra, it is consideredthat most of Sn existed in center ofcomposite #1 and #2.

annealing method. The different-shaped
nanostructure was formed on eachcomposites surface. The wires, which arethin, straight, and have sharp top end, weregrown in composite #1. On the other hand,
the rods, which have an unhomogeneous
shape and a rugged surface, were grown incomposite #2.
The synthesis condition of Na-Sn systemcomposites is clarified by characterization of
the cross-section of composite #1 and #2.

4. Conclusion
Nanostructured sodium – tin systemcomposites were obtained by simple thermal

Acknowledgments
I would like to thank T. Nonomura, S.
Cai, K. Shirai, T. Matsushita, D. Ishikawa, Y.
Usuda and H. Tatsuoka for their fruitful
discussions.

Reference[1] M. Jeon, K. Kamisako, Mater. Let 63(2009) 777[2] Y. Wang, J. Y. Lee, J. Phys. Chem. B 46
(2004) 17832[3] M. K. Jayaraj, K. J. Saji, K. Nomura, T.
Kamiya, H. Hosono, J.Vac.Technol. B 26
(2008) 495[4] H. Schaefer, B. Eisenmann, W. Muller,
Angew. Chem. 85 (1973) 742[5] W. muller, K. Volk, Z. Maturforsch. 32b(1977) 709[6] W. Hume-Rothery, J. Chem. Soc. 131

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36 (1997) 4316
[8] T. F. Fassler, C. Kronseder, Angew.
Chem. 110 (1998) 1641; Angew. Cem.
Int. Ed. 37 (1998) 1571
[9] T. F. Fassler, S. Hoffmann, Inorg. Chem.
42 (2003) 5474
[10] F. Dubois, M. Schreyer, T. F. Fassler,
Inorg. Chem. 44 (2005) 477

Synthesis of MoSi2 by Molten Salt Method
Tomomi Nonomura
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
e-mail: f0030154@ipc.shizuoika.ac.jp

Abstract
MoSi2 was grown on Mo substrate using the molten salt method. The effect of thetemperature on the structure was investigated. Hexagonal phase of silicide was formed at low
temperature. It was also revealed that tetragonal phase of the silicides was formed on thesubstrate at high temperature. It was found that MoSi2 layer, which thickness was about 4 µm,
was formed, and the Si atoms were distributed throughout the substrate.

1. Introduction
Semiconducting silicides have attracted
much attention due to their potential tocreate new classes of environmentallyconscious electronics. For the group VI
metal silicides, many studies of CrSi2 havebeen reported and considered to be one ofthermoelectric materials in the past decade[1][2]. MoSi2 and WSi2 have two phases. One
is stable tetragonal (t-) phase, which ismetallic, and used as gate materials in large
scale integrated circuit. On the other hand,another hexagonal (h-) phase MoSi2 andWSi2 epitaxial layers have been grown on Si
substrates, and the h-MoSi2 and h-WSi2 areconsider to be narrow gap semiconductors [3].
It have been investigated the electronic
structures and thermoelectric properties of
the h-MoSi2 and h-WSi2, as well as the CrSi2
bulk crystals by using the first-principles
total-energy
calculation
programin
pseudopotential schemes with plane-wave
basis functions. The calculated results showthat h-MoSi2 and h-WSi2 are expected tohave improved thermoelectric properties [4].
However, the electrical property of the bulk
crystal h- MoSi2 and h-WSi2 are not wellunderstood.
However, it is generally difficult to grow
h-MoSi2 and h-WSi2 bulk crystals from theliquid phase for the existence of the hightemperature phase of t-MoSi2 and t-WSi2,
respectively. The semiconducting phases of
h-MoSi2 and h-WSi2 are grown by themechanical alloying method and the molten

salt method for the bulk growth techniques. The successful growth of the lowtemperature phase of β-FeSi2 bulk crystal
was previously investigated, and thestructural and electrical properties werereported [8].
In this study, the Mo-silicides layer weresynthesized by the molten salt method byusing Mo substrate, and the detail structuralproperty of the silicides were examined.
[4][5][6][7]2. Experiments
Mo- and W-silicide layers were grownon Mo (99.9%up) and W (99.9%) substrates,
respectively, using a molten salt, comprisedof 36.58 mol%NaCl - 36.58 mol%KCl 21.95 mol%NaF - 4.89 mol%Na2SiF6. The
salt mixture and silicon powder (99.999% 74µm, 21.85 mol% for the salt mixture) wereplaced in an SiO2–Al2O3 crucible with thesubstrates. They were thermally treated at
700 ,750 and 800 ºC. After the heat treatment,
the salt was removed from the sample using
deionized H2O.
The structural properties of the layers
and compacts were characterized by X-ray
diffraction (XRD) measurements using a Cu
Kα1 source and scanning electron microscopy(SEM) along with energy dispersion
spectroscopy (EDS).

3. Results and Discussion

Figure 1 shows XRD spectra of thelayers on the Mo substrate growth by using
molten salt method at 700 ,750 and 800 ºC,
respectively. It is observed that MoSi2 wasformed on the substrate surface. On the other
hand, the t-phase was formed as well as theh-phase. Especially h-phase was muchgrown at lower temperature than higher one.
The stable t-phase was formed at higher
temperature.

Fig. 1 XRD spectra of the layers on the Mo
substrate growth at 700 ºC, 750 ºC and 800 ºC,
respectively.

Figure 2 shows cross-sectional SEM
image of the substrate with corresponding
EDS mapping of the MoSi2 layer growth at
700 ºC. The figure shows the half size of thesubstrate from the surface to the center. The
surface of substrate had irregular structure.
The result of EDS mapping shows that
MoSi2 layer was formed around the substratesurface. And the Si atoms were also
distributed throughout the substrate.

Fig. 2 Cross-sectional SEM image of thesubstrate with the corresponding EDS mappingof the MoSi2 layer growth at 700 ºC.

Figure 3 shows cross-sectional enlarged
SEM image of the substrate with thecorresponding EDS line analysis, EDS
surface mapping and chemical compositionof MoSi2 layer growth at 700 ºC. The resultof the EDS line analysis shows, though it isobserved that the thickness of the MoSi2
layer was about 4 µm, it was not observed
the boundary between the MoSi2 layer andMo substrate. Moreover it was observed thatthe composition ratio of the layer was Mo:Si≈ 1:2.

Fig. 3 Cross-sectional enlarged SEM imageof the substrate with the corresponding EDS
line analysis , EDS surface mapping andchemical composition of the MoSi2 layer
growth at 700 ºC.

Figure 4 shows cross-sectional enlarged
SEM image of the substrate with thecorresponding EDS inner mapping andchemical composition of the Mo substrate
growth at 700 ºC. The result of the chemicalcomposition shows that the Si-richcompound was formed. The size of the Siatom is about 2 µm. It is considered that thevoid was existed throughout the Mo
substrate before the reaction.
Nonomura et al. reported the structure
property and the structure property for theMo-silicide synthesized by MA method andmolten salt method. It was considered it isdifficult to grow bulk h- MoSi2 with highquality crystalline. This result shows hMoSi2 layer growth. It is expected that bulk
h- MoSi2 would be obtained by the longthermal treatment using h- MoSi2 layer bymolten salt method.

[6] T. Nonomura, C. Wen, K. Shirai, K.
Isobe, A. Kato, Y. Kubota, T. Nakamura,
Y. Hayakawa and H. Tatsuoka, Mater.
Sci. Eng. 18 (2011) 142010.
[7] R. O. Suzuki, M. Ishikawa, K. Ono, J.
Alloys Compd. 306 (2000) 285.
[8] Okubo M, Ohishi T, Mishina A,
Yamauchi I, Udono H, Suemasu T,
Matsuyama T and Tatsuoka H, Thin
Solid Films 515 (2007) 8268.
Fig. 4 Cross-sectional enlarged SEM image ofthe substrate with the corresponding EDS inner
mapping and chemical composition of the MoSi2
layer growth at 700 ºC.

4. Conclusion
MoSi2 was grown on Mo substrate using
the molten salt method. It was revealed that
tetragonal phase of the silicides was formedon the substrate. It was observed that MoSi2
layer, which thickness was about 4 µm, wasformed, and the Si atoms were distributed
throughout the substrate. It is expected thatbulk h- MoSi2 would be obtained by the longthermal treatment.

Acknowledgements
I’d like to thank the sponsor of SSSV
program for giving me the chance tocomplete this special issue. On earlier
iterations of this paper, I would like toacknowledge the help of H. Tatsuoka, W.Li
and D.Ishikawa. This article was improvedby the suggestions of them. I thank them fortheir time and thoughtful reflections.

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Y. Kubota, T. Nakamura, M. Yamashita,
Y. Hayakawa and H. Tatsuoka, Phys.
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Synthesis of Molybdenum Trioxide Nanostructuresby Thermal Oxidation Technique
Hiroto Sato
Faculty of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka, 432-8561, Japan,
Phone/FAX.:+81-53-478-1099
f0811057@ipc.shizuoka.ac.jp

Abstract
Molybdenum trioxide (MoO3) nanostructures were grown by the thermal oxidation
technique of molybdenum wires in an ambient atmosphere. Those nanostructures wereobserved on the inner wall of quartz tube just at the center and the place of about 5 cm awayfrom the center of the tube furnace. Those nanostructures were characterized by X-ray
Diffraction (XRD) and Field Emission Scanning Electron Microscopy (FE-SEM). The
structures of synthesized nanostructures are orthorhombic MoO3.

1. Introduction
Quasi-one-dimensional nanostructureswith different morphologies such as
nanotubes, nanowires, and nanobelts haveattracted much attention over the past decadedue to their remarkable physical/chemical
properties and wide range of applications [1,2]. Layered crystal MoO3 has been growing
in recent years for its multifaced functional
properties. MoO3 has been widely used as
charge-density wave conductors [3, 4],
optical materials [5−7], display devices [8,9], and catalysts [10−14].
In this study, we report the simple
synthesis of MoO3 belt and nanostructuresby direct thermal oxidation in the air.

2. Experiments
This paper conducted two experiments.
First of all, the fabrication of MoO3
nanostructures was based on thermaloxidation of molybdenum metal in an
ambient condition. In this experiment, theMolybdenum wires (diameter: 1 mm length
1 cm) were polished with sandpapers inorder of #400 to #1500, and ultrasoniccleaning with Acetone, then a molybdenum
wire was heated in electric furnace fromroom temperature to set temperature in air.
After reaching preset temperature, it washeated for 1-12 hours. After heat treatment
the wire was cooled to room temperature in
the atmosphere.
Secondly, after the electric furnace being
raised to 750ºC, silicon (111) substrates wereinserted into it for oxidation reaction. Afterthe thermal oxidation processing for 2.5 - 10
minutes, their silicon substrates were downout from the furnace and cooled down toroom temperature naturally.
The as-grown nanostructures werecharacterized by FE-SEM (JSM-7001F) andXRD (RINT Ultima III).

Fig. 1, SEM images of molybdenum nanostructures for (a) 1h, (b) 4h, and (c) 12h.

3. Results and Discussion
Figure 1 shows SEM images of the
prepared at 600ºC. These figures show thatsize of MoO3 nanostructures depend on
growth time.
Fig.
2
shows
molybdenum
nanostructures were grown by the thermaloxidation at 750 ºC. It could be observed on
the inner wall of quartz tube just at the centerand the place of about 5 cm away from thecenter of the tube furnace. Theirnanostructures are colorless and plate-likestructures. Their nanostructures were formed

Fig. 2, Image and SEM image of nanostructures
for 6h.

on Si substrate.
Figure 3 shows the XRD spectrum.
Molybdenum oxide nanostructures wereobserved on the Si substrate.
XRD
spectrumshows
thatnanostructures
synthesized on Si substrate are orthorhombic
MoO3. It is known that lattice parameters of
MoO3 are a. 3.962 Å, b. 13.85 Å and c.
3.697 Å (JCPDS 05-0508).
Figure
4
shows
that
MoO3
nanostructures were uniformly distributed
over the entire substrate on Si substrate. The
sizes of nanostructures are from about 100
nm to 4 µm.

Fig. 3, XRD spectra of nanostructures on Si
substrate at 750 ºC for 10 minutes.

Fig. 4, SEM image of nanostructures on Si
substrate at 750 ºC 10 minutes.

4. Conclusion
In summary, MoO3 nanostructures weresynthesized by direct oxidation ofmolybdenum metal in air.
The difference in the size of MoO3
nanostructures becomes small as growthtime becomes long.
MoO3 nanostructures were grownthermal oxidation technique of Mo wires on
the Si substrate. XRD spectrum indicatesthat the MoO3 structures are orthorhombic
structure.

Reference[1] M. Law, J. Goldberger, P.D. Yang, Annu.
Rev. Mater. Res. 34 (2004) 83.
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Non-Cryst. Solids 185 (1995) 135.
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Rabalais, J. Appl. Phys. 49 (1978) 409.
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J. Electrochem. Soc. 141 (1994) 876.
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S.-Y. Peng, in: Y.A. Attia (Ed.), Sol−Gel
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N. Shigemoto, Ind. Eng. Chem. Res. 34(1995) 137.

[13] M. Fournier, A. Aouissi,
C. Rocchiccioli-Deltcheff, J. Chem. Soc.
Chem. Commun. 3 (1994) 307.
[14] P. Benard, L. Seguin, D. Louer,
M. Figlarz, J. Solid State Chem. 108(1994) 170.

Synthesis of Copper Oxide Nanowires andTheir Application for Solar Cells
Shanming Cai*
Graduate school of Engineering, Shizuoka University
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka,
432-8561, Japan
TEL&FAX: +81-53-478-1092*
e-mail: f0030118@ipc.shizuoka.ac.jp

Abstract
This paper introduces synthesis of copper oxide nanowires using thermal oxidation routeand their application for solar cells. Thermal oxidation method is a simple and cheap route tosynthesize metal oxide nanowires. The morphology of as-grown nanowires were characterizedby X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron
microscopy (TEM) with selected area electron diffraction (SAED). The compositional analysis
was made using energy dispersive X-ray spectroscopy (EDS). Furthermore, the nanostructureof copper oxide nanowires is discussed.

1. Introduction

2. Experiment

Nanostructured materials have attracted
much scientific attention due to theirinteresting size-dependent chemical andphysicalproperties
andpotential
technological applications.
Cuprous oxide (Cu2O), a direct-gap
semiconductor with a band-gap energy of 2.0
eV, has been regarded as one of the mostpromising materials for application tophotovoltaic cells (Pollack and Trivich,1975). The attractiveness of Cu2O as aphotovoltaic material lies in the fact that theconstituent materials are nontoxic andabundantly available on the earth, and that
the Cu2O has a high absorption coefficient invisibleregionsandlow-cost
producibility. Cu-oxide nanowires have beengrown on Cu substrates by thermal oxidationin the air. However, the detailed structure
and the growth mechanism of the Cu-oxide
nanowires have not been clarified.[1]The present work reports the detailed
structure and growth of heterostructured
CuO/Cu2O nanowires prepared by thermaloxidation in air. The as grown nanostructures
have been observed in detail. Growthmechanism of the nanowires has been
discussed.

Synthesis of copper oxide Nanowires iscarried out on a copper foil (99.9% purity,
1.0mm thickness) of about 10×10 mm2 usedas the starting material. The copper foils
were grinded by #400 and #1500 sandpapers
to remove the oxide layer, followed byultrasonic cleaning in ethanol and acetone,
successively. The copper foils were loadedin the center of a furnace. The growthoccured under atmospheric pressure in air.
The nanowires were grown for 0.5–4 h at

the temperatures of 773–1073 K.
The resulting nanostructures werecharacterizedbyscanning
electronmicroscopy (SEM), transmission electron
microscopy (TEM) with selected area
electrondiffraction(SAED)
andhigh-resolution
transmission
electronmicroscopy (HRTEM). The compositional
analysis was made using energy dispersive
X-ray spectroscopy (EDS).

3. Results and Discussion
Figure1 shows the XRD spectrum of thesample grown at 873K for 1h in the air. It isnoted that both of CuO and Cu2O phases
were formed simultaneously.

Fig.1 XRD spectrum of the sample grown at
873K for 1 h.

Figure 2(a) shows SEM images of asgrown nanowires of the same sample inFig.1. It is observed that the average lengthof the nanowires is about 9.3 µm. Figure 2(b)shows an enlarged SEM image of the
nanowires shown in Fig.2(a). We find thatthe diameter of bottom part of one nanowire
is distinctly different from the top part. The
diameter of bottom part and top part areabout 10 nm and several µm, respectively.

Fig.2 (a) SEM image of as grown Nanowiresof the same sample in Fig.1. (b) An enlarged
SEM image of the Nanowires shown in Fig.2(a).

Figure 3 shows the STEM image of one
as grown nanowires shown in Fig.2. We findthat the distribution of oxygen element on
the top part of the nanowire is much higher
than that of the bottom part. Table 1 shows
the point analysis of chemical elements

Fig.3 STEM image of one of the as grown
nanowires with the corresponding EDS
mapping for Cu and O. Point EDS results areshown in Fig.3(d)shown in Fig.3. From the results of point
analysis, it is found that there are two phases
in one nanowire. It is considered that CuOphase is formed at the top part of the
nanowire, and Cu2O phase is formed at thebottom part of the nanowire.
Fig.4(a),(b) show the HRTEM image of
top part and bottom part of the nanowires
with the corresponding fast Fouriertransformation, respectively. CuO has amonoclinic structure (a= 0.4668 nm,
b=0.3442 nm and c=0.5131 nm, =99.506
degrees). On the other hand, Cu2O has a
cubic structure with the lattice constant of
0.4217 nm. As shown in Fig.4(a),(b), theinterplanar spacings are corresponding to themonoclinic CuO and cubic Cu2O,
respectively.

Fig.4 (a) HRTEM image of top part of the
nanowires with the corresponding fast Fouriertransformation. (b) HRTEM image of bottom
part of the nanowires with the corresponding
fast Fourier transformation

The growth mechanism of the as grown
nanowires is discussed as below. At theinitial growth stage of the nanowires, single
phase CuO nanowires are formed. With theincreasing of the temperature, Cu atoms
diffusion is enhanced. However, Oxygen
atoms are consumed during the formation of
the CuO domains. Then, Cu diffused into
CuO nanowires to form the Cu2O domains
because of the insufficiency of oxygen.
Consequently, CuO/Cu2O heterostructures
are formed.

4. Conclusion
CuO/Cu2O heterostructures nanowires
are grown by thermal oxidation in air

successfully. The as grown nanowires are
divided by two regions of CuO on the top
part and Cu2O on the bottom part. The
growth mechanism of the CuO/Cu2O
heterostructures is discussed.

Acknowledgments
I would like to thank H. Fujii and T.
Matsushita and H. Tatsuoka for theirfruitful discussion.
Reference[1] K. Akimoto , , S. Ishizuka, M. Yanagita,
Y. Nawa, Goutam K. Paula and T. Sakurai,
Solar Energy 80 (2006) 715.

Composite coating made by micro-arc oxidation followed by LPD technology on
Mg-Zn-Ca alloy and anticorrosion property
Bai Kuifeng, Guan Shaokang*, Zhu Shijie, Wang Liguo, Hu Junhua
Material Research Center, School of Materials Science and Engineering, Zhengzhou University,
Zhenzhou, 450002, China*E-mail: skguan@zzu.edu.cn

Abstract:
Recent studies indicate that Mg alloy have many good biocompatibility with living boneand gradual degradation in human physiological environment. Therefore, it is widely used as
biomaterials in human body. But the poor corrosion resistance of magnesium alloys in a humanbody fluid or blood plasma limited their clinical application. In this study, a composite coating
was prepared on Mg-Zn-Ca alloy by micro-arc oxidation (MAO) method at the working
voltage of 120-140V. The next research indicated that the MAO coating was porous. In order
to diminish the contact with SBF, a TiO2 layer was coated on the porous MAO layer by liquid
phase deposition (LPD) followed by a annealing treatment. The coating were characterized by
X-ray diffractometry (XRD), Fourier transform infrared spectroscopy and scanning electron
microscopy (SEM). XRD analysis indicated that the MAO coating was mainly composed of
Mg3(PO4)2. SEM observation showed that a dense and uniform coating was formed. The
electrochemical performance of the MAO and TiO2/MAO coated alloy was evaluated byanodic polarization measurements. The electrochemical polarization tests showed that theMAO coating and composite coating improved the anticorrosion properties observably. The
results showed that, compared with the as-cast Mg-Zn-Ca alloy, the corrosion potential of the
MAO coating and the composite coating increased from -1.771V to -1.6478V to -1.601V. In
the same way, the corrosion current density decreased from 1.43×10-4 to 6.35×10-5 to 5.69×10-7.
Thus Mg alloy coated with MAO and composite coating is a better method to the developmentof biomaterials.

Key words: Mg-Zn-Ca, Mg phosphate, TiO2, LPD, anticorrosion resistance
1. Introduction
Magnesium alloys, as a new kind of
bio-implant materials, have attracted great
attention because they possess significant
advantages, such as high specific strength,
perfect
biocompatibility
andnaturalabundance [1-3]. Unfortunately, magnesium
and its alloys are highly susceptible tocorrosion in physiological environment of
human body, including human body fluid or
blood plasma, which has limited theirbiomedical applications [4, 5]. For thisreason, it is important to find ways toimprove the corrosion resistance ofmagnesium alloys. To date, various surfaces
coating technology was developed to reducethe magnesium alloy's degradation rate andimprove the bioactivity. For example, TiO2
coatings, phosphate, hydrotalcite andcarbonate et al. fall in this category ofmaterials [6-10]. In this paper, we used ahome-developed Mg-Zn-Ca alloy was usedas substrate. A new Composite technology
which combines the MAO and LPD wasdeveloped. The process of MAO canfabricate ceramic coating with high adhesion
to the substrate, but it is porous. The LPD
technology could seal the porous of the
MAO coating with TiO2 particle. The goalof this article is that the composite coating
made by the composite technology could
improve the corrosion property greatly.

2. Experiment2.1 Preparation of MAO coating
In this paper, the Mg-Zn-Ca alloy
substrate was prepared in our research group.
Before MAO experiment, the as-cast
Mg-Zn-Ca alloy was first cut into
dimensions of 25 mm×10 mm×4 mm, after
that, the sample were ground and polishedwith SiC papers progressively to 800 grits,
followed by ultrasonically cleaning in
acetone and ethanol for 15 min, respectively.
During the process of MAO, the sample wasused as the working electrode. The
electrolyte was prepared by dissolving 0.1
mol/L Na3PO4.12H2O and 0.05 mol/l NaOH
into distilled water. All the reagents
mentioned were analytic grade. The MAO
coating was prepared under a constantvoltage mode and the cell potential wasincreased gradually till-180V, then theworking electrode was oxidized for 20-30
min.
2.2 Preparation of LPD coating
In this method, TiO2 layer was fabricatedby hydrolysis of [NH4]2TiF6. The details of
the liquid phase deposition were discribed inreferences [16]. After the mother solution
was prepared, the MAO coating substrate
were immersed into the mother solution andsuspended vertically in the mother liquor for
40h at the room temperature. After that, thesample was took out from the solution andrinsed in the deionized water and dried atroom temperature. After deposition, thesample was annealed at 250 in the vacuum
furnace and the annealing time was 2h.
2.3 Dynamic potential test
A three-electrode cell was used forthe in vitro potentiodynamic polarization
measurements. The counter electrode wasmade of platinum and the reference electrode
was a saturated calomel electrode. The areaof the working electrode exposed to thesolution was 1cm2. The measurements werecarried out on an electrochemical station(RST 5000). The Simulation body
fluids(SBF) we were used was preparedaccording to Ref[17].
The phase composition of coating wereidentified by X-ray diffraction (XRD, Philips,
PW-1700X) with Cu Ka=0.15406 nm in thescanning angles from 20°to 80°. The surface
morphology was observed using electron
microscopy (SEM). The functional groups ofthe composite coatings can be determined by
Fourier transform infrared spectroscopy(FTIR).

3. Result and discussion

Fig (1) (a) XRD pattern of MAO treated Mg alloy; (b) XRD patterns of bare Mg-Zn-Ca alloy

From picture (a) we can see that only
magnesium phosphate and Mg’s peaks wereobserved. This shows that MAO coating was
mainly composed by the magnesium
phosphate. The diffraction peaks of alloy
substrate were quite strong because that thecoating was porous and the X-ray canpenetrate through the coating. Magnesium

phosphate is a biological material with good
flexibility.
In order to further determine theFunctional groups of MAO coating, Fouriertransform infrared spectroscopy (FTIR) wasused to test the powder which were scraped
off from the coating above.

Fig (2) The Fourier transform infrared spectra

From the infrared spectra, hydroxyl andphosphate absorption peak were shown on
that. The absorption at 1417cm-1 was due tothe substitution of CO32- ions in the positionsof PO43- ions. The absorptions at 1037,605
and 566 cm-1 can be attributed to the PO43group, which indirectly proves the existenceof magnesium phosphate. The absorption atabout 3435cm-1 was the absorbance peak ofhydroxy. The reason was due to the presenceof OH- ion. The resultant MAO layer surface
was a hydroxyl group terminated. The
dehydration reaction can also occur betweenhydroxyl group absorbed on the surface andTiO(OH)x group. In such a case, the TiO2
layer can chemically anchor on the porous
MAO surface. On the other hand, theinteraction of TiO2 and magnesium
phosphate would also benefit the
improvement in interface condition andanticorrosion properties. They can anchor
strongly with each other via Mg2+ cations
that bridge O from TiO2 with phosphate O.
The earlier studies on the interaction of TiO2
and phosphate (or sulfate) proved primarilyon the improvement of the thermal stability
and interface boning. The heating after
deposition was expected to strengthen thisbonding on the interface. It was concluded
that the interface condition is also an
important factor to influence its degradation.
A chemically bonded interface would benefit
the improvement of biocorrosion resistance.
The last infrared absorption peaks at about
1639cm-1 which reason is not clear. Maybe
the purity of the potassium bromide was not
sufficient.

Fig.3 (a) and (c) shows the surface morphology of the alloy after MAO treated at the voltage of 120V
and 140V for 20 min, respectively.
Fig.3 (b) and (d) shows the surface SEM images of the MAO samples after coated by TiO2 using LPD
method respectively

During the MAO treatment, to maintainthe stable oxidation process and avoid theformation of crack and large-sized pores, thevoltage was kept within a scope of 120-140V.
The surface was crack-free and characterizedwith uniformly distributed pores. The poresare well separated. The diameter of the poresvaries from 1 to several microns withincreasing of the applied voltage from 120 to
140 V. The density and size of pores on the
MAO coating were influence by the applied
voltage. The formation of the pores on the
MAO coating would permit some ions in
SBF to penetrate through the coating, which
would cause the corrosion and mechanicalintegrity loss of substrate in the initial
implantation stage. It is necessary to seal thepores by other biocompatible materials.
Figure 3 (b) and (d) shows the surface
SEM images of the samples after coated by
TiO2 using LPD method correspondingly. It
was clearly shown that the surface was
nonporous after coated by TiO2 layers.
Cracks could be observed on the TiO2
coatings. Cracking formation could be

attributed to a large number of causes whichlead to building up of stresses. For
Mg/MAO/TiO2, cracking could be mainly
ascribed to the thickness of the TiO2 coating
required to cover the pores in the MAO layer
and mismatch between MAO and TiO2 layer.
In Fig.1 (d), some pores were left unclosed.
The inset showed the enlarged image ofpores. It can be observed that the TiO2 was
uniformly deposited inside the pores. By
comparing Fig.1 (b) and (d), it can beconcluded the pores size and density in
MAO layer can affect the up TiO2 layer
evidently.
It was clearly seen that the as-cast
Mg-Zn-Ca alloys has a corrosion potential(Ecorr) of about -1.771V. When the MAO
coatings were coated on Mg-Zn-Ca alloys
substrate, the obviously positive offset ofpolarization curves can be observed. With
the working voltage of the MAO coating
increased from 120V to 140V, thecorresponding corrosion potentials increased
from -1.687 V to -1.648 V, respectively. And
the corrosion current density was decrease

from 1.43 ×10-4 to 1.57 ×10-5 and 6.36 ×10-5,
respectively. Thus, it is necessary to seal thepores in order to further diminish the contactof substrate with the fluid environment. Aftercoated by the TiO2/MAO composite coating,the corrosion potential was improved from-1.63V to -1.601V, respectively. As shown in
Fig4 (b), when the working voltage was
120V, the corrosion potential of thecomposite coating was -1.630V. Similarly, inthe working voltage of 140V, the corrosion
potential of the composite coating was-1.601V. In the meantime, the corrosion
current density of the composite coating
obtained from Tafel curves was decreased to2.35 ×10-6 and 5.69×10-7. It can beconcluded that the upper TiO2 layer played
significant role in the improvement ofcorrosion resistance of the samples.

Fig.4 (a) Tafel curves for MAO coated alloys at the working voltage of 120V and 140V; (b) Tafel curvesof alloys coated by TiO2/MAO composite coatings at the working voltage of 120V and 140V,
respectively.

4. Conclusion
In this work, a process of MAO
combined with LPD was used to preparecomposite coatings on Mg-Zn-Ca alloy.
Compared with the Mg-Zn-Ca alloy
substrate, when the working voltage of MAO
process was 120V and 140V, the corrosion

potential was improved by about 84mv and123mv, respectively. At the same time, thecorrosion current density was decreased
from 1.43×10-4 to 1.57×10-5 and 6.36×10-5.
After coated by the TiO2/MAO composite
coating, the corrosion potential wasimproved by 141mv and 170mv comparedwith the Mg-Zn-Ca alloy substrate.
Meanwhile, the corrosion current density
was decreased from 1.43×10-4 to 2.35 × 10-6
and 5.69 × 10-7. Therefore, the process of
MAO combined with LPD was proved abetter way to fabricate composite coating on
magnesium alloys.

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Huang, H.Q. Liu, Appl. Surf. Sci. 254(2008) 2712-2719.
[9]L.Xu, F.Pan, G.Yu, L.Yang, E.Zhang,
K.Yang, Biomaterials 30(2009)1512-1523.
[10]J.Hu, C.Zhang, S.Guan, L.Wang, S. Zhu,
Mater. Lett. 64(2010)2569-2571.
[11] J.H. Gao, S.K. Guan, J. Chen, L.G.
Wang, S.J. Zhu, J.H. Hu, Z.W. Ren,
Applied Surface Science 257 (2011)
2231—2237.
[12] Junhua Hu, Guan Shaokang, Caili
Zhang, Chenxing Ren, Cuilian Wen,
Zhaoqin Zeng, Li Peng, Surface &Coating Technology 203 (2009)2017—2020.
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Sun, S.J. Zhu, B. Wang, Materials Letters
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[14] Zunjie Wei, Hui Du and Erlin Zhang,Surf. Interface Anal. 2011,43, 791—794.
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L.X. Guo, Y. Zhou, J.H. Ouyang, J.M.
Dai, Current Applied Physics 11 (2011)
1405—1409.
[16] Bo Lin, Ting Li, Yong Zhao, Fang Ke
Huang, Lin Guo, Yu Qi Feng, J.
Chromatogr. A 1192 (2008) 95.
[17] Tadashi Kokubo, Hiroaki Takadama.
Biomaterials 27(2006) 2907-2915.

The Effect of Microscale Mn and RE Addition to AZ61 Cast Tissue
Liu Zhidan, Qu Ruixiao, Wan Yuanyuan, Li Qingkui, Guan Shaokang*
Materials Research Center, School of Materials Science and Engineering, Zhengzhou
University, 97# Wenhua Road, Zhenghzhou 450002, P.R. China*E-mail: skguan@zzu.edu.cn

Abstract
The destination of this paper is to research changes of AZ61 cast microstructure under thecondition adding small amount of manganese and mixed rare earth element. Throughmetallographic picture inspection and SEM test of etched sample, alloys of different
composition are compared range from macroscopic homogeneity, phase distribution andmorphology changes and segregation of each element in the alloy. With the increase of Mnelement, Al8Mn5 has a slight increasing in the number and size. As the addition of rare earth,
Al4RE phase's number and size increase in the grain boundary and intragranular significantly.
Component 4 (Mn = 0.68wt %, RE = 0.67wt %) has the best cast structure.

1.

Introduction

2.

Magnesiumhasahexagonal
close-packed structure and poor deformation
ability. Typically, most of extruded
magnesium alloy products are not important
force components, because of its poor
mechanical properties compared to othermetal structural material. Therefore, it isnecessary to adopt some methods to improvetheir mechanical properties. In this paper, thecommercial magnesium alloy AZ61 is added
both manganese [1] and mixed rare earth
elements trying to change the homogeneityof its microstructure, in order to improve
deformation coordination in the extrusion
process and uniform mechanical propertiesof products.

Experimental

The alloy is composed of AZ61, Mn and
RE. Mn / RE ratios range from 0.305 to 1.87.
In the design of this component, Mn
compositioninterval:
0.78-0.3,
RE
composition interval: 0.39-1.0. Following
components are detected by direct reading
spectrometer from the front of furnace.
Ce-rich mixed rare earth, composed of:
cerium (Ce) about 65%, neodymium (Nd)about 10%, lanthanum (La) about 20%, also
contains a small amount of praseodymium(Pr) about 5%.
In the 2/3 cross-sectional diameter ofcast sample, Φ15 × 15mm metallographic
specimens are intercepted, grinded out theplane and chamfer followed by 180 #, 400 #,600 #, 800 # and 1000 # metallographic
sandpaper, then was mechanical polished

Table 1 The composition of alloying elements detected by direct reading spectrometer
Component

Vol.

Al

Zn

Mn

Fe

Si

Cu

Ni

Be

RE

1(RE=0)5.8354 0.8145 0.3004 0.0042 0.0067 0.0012 0.0005 0.00006 02(Mn=0.7230)

5.8261 0.8421 0.7830 0.0021 0.0073 0.0011 0.0005 0.00007 0

3(Mn/RE=1.87)

5.6014 0.9358 0.7312 0.0012 0.0093 0.0012 0.0008 0.00008 0.394(Mn/RE=1)

5.6874 0.8487 0.6807 0.0014 0.0075 0.0014 0.0006 0.00007 0.67

5(Mn/RE=0.775)

5.6581 0.9256 0.6201 0.0010 0.0065 0.0013 0.0011 0.00005 0.8

6(Mn/RE=0.593)

5.6428 0.9254 0.5395 0.0013 0.0077 0.0015 0.0008 0.00009 0.91

7(Mn/RE= 0.305)

5.7432 0.9821 0.3047 0.0014 0.0094 0.0012 0.0006 0.00009 1.0
Table 2

Corrosive components

Caustic

Picric acidAcetic acid

Alcohol

Water

Cast Sample

3g

20ml

50ml

20ml

using the size of 1µm diamond polishingpaste.
And
finallythecorrosion
metallographic photos are shot. Etchingsolution formulations shown in Table 2.
Metallurgical microscope is OLYMPUS
BH2-UMA. We use the Scanning Electron
Microscope (SEM) to observe the state of
the high-powered microstructure. Using
electron spectroscopy, electron beam throughthe inner atomic electron excitation of thesample, the qualitative and quantitative
analysis is taken. In this study, we takeadvantage of JSM-6490LV scanning electron
microscope.

3.

Results and Discussion

AZ61A is the Mg-Al-Zn series ofcommercial wrought magnesium alloy, andthe main alloying elements are Al and Zn. As
the Zn content is less than one percent (massfraction), which generally do not forming
new compounds [2], and mainly exist in the
α-Mg solid solution and Mg17Al12 binary
phase. According to the Mg-Al binary phase
diagram, equilibrium eutectic occurs in
437 ℃ , the eutectic reaction occurs as
follows: L
α-Mg+γ-Mg17Al12. The
aluminum content at eutectic point is 32.3%,

aand Al in the magnesium alloy achieves the
maximum solubility of 12.7%, then thesolubility decreases with temperature until
100 ℃, 2%. In the actual casting process,due to the characteristics of rapid
solidification rate, the large size ofmagnesium atoms, great diffusion activationenergy in the solid state and small latent heat
during the crystallization, magnesium alloy
solidification occurs in the non-equilibrium
solidification,
generating
α-Mgand
γ-Mg17Al12 eutectic.
In figure 1, those are the low times AZ61
magnesium alloy microstructures, which arecooled with the mold. In the pictures, itwrapped in dark phase is the pro eutectic
α-Mg solid solution, and the white phasewith a black outline in the boundaries oflarge pro eutectic grain is phase ofγ-Mg17Al12 due to non-equilibrium eutectic
solidification. In the alloy, the pro eutectic
α-Mg organization is more than the eutecticγ- Mg17Al12, so when the eutectic reaction
occurs, the α-Mg phase attaches to the pro
eutectic α-Mg phase and grows. It pushes theγ - Mg17Al12 to the final solidification place,the α-Mg grain boundaries, leading toirregular distribution of γ-Mg17Al12 on thegrain boundaries like blocks.

b

c

d

e

f

g

Fig. 1 The metallographic of each cast alloy prepared by water-cooled casting and cooling with mold

As shown in Figure a to g, in the alloy,
thick dendritic γ-Mg17Al12 phase gradually
be thinner and changes from continuous
thick dendritic morphology into a continuous
thin dendritic, semi-continuous network andthe short island morphology. Figure d to
Figure g, γ-Mg17Al12 phase graduallydispersed, showing a gradually smaller,
scattered trend, and has needle-like phaseformation, then the size and number ofneedle-like phase gradually larger and more.

It can be seen from the figure that all thealloys with different composition show: With
the cooling rate decreases, the number of
second phase become more in the structure;
pro eutectic α-Mg phase and γ-Mg17Al12
phase size decreases, and the microstructure
becomes more dense; the number of Al8Mn5
phase and needle-like phase increases,
however the size has no significant change
uniformly.
We can learn from the above chart thatthe above alloys, composition 4 and 5, havethe best cast structure, and that is to say
when the Mn / RE ratio range from 0.775 to
1. The γ-Mg17Al12 in alloy is scattered small,and not with dendritic characteristic,
however the other second phases have lowamount, so it is conducive to subsequent hot
extrusion. Figure 2 shows high-powered
scanning photographs of alloys 1, 2, 7.

a

b

c

Fig. 2 The SEM photograph of cast alloy 1, 2, 7

From the figure a and b we can see thatγ-Mg17Al12 in figures is thinner than in a, andthe only distinction between these two is Mnelement content. By Mg-Mn binary phase
diagram shows, Mn solubility in Mg alloy
dropped significantly as the temperature
decreases, from 650 ℃, 2.2wt% down to
300 ℃, 0.1wt%. The reason is rapid cooling
speed of water-cooled cast mold. In thesolidification process, part of the Mnelement in the alloy react with Al forming
AlMn binary phase or AlMnFe ternary
phase; the last part of the Mn supersaturated
in Mg is unable to form a new phase before
solidification takes place. In the subsequentcooling process it precipitates in the α-Mggrain boundaries, thus preventing the grainand γ-Mg17Al12 phase of growing up.
Contrasting figure a and b, by adding
1.0wt% of the mixed rare earth, the alloy
microstructure of γ-Mg17Al12 phase
precipitation becomes less, while there has
been needle-like rare-earth phase.
The
addition of mixed rare earth refined alloy
from two aspects. For the first, mixed rareearth can refine the γ-Mg17Al12 phase, tomake it small and to be fragmentation.
According to the theory of electronegativity
elements, the difficulty degree of forming
different compounds is determined by eachelement's electronegativity value. The more
electronegativity differences between twoelements, the bigger their binding force, andthe easier to form compound. The element
electronegativity values of Mg, Al, Ce, La
are shown in Table 3.
From the table 3, we can see theelectronegative difference between Ce and
Al or La and Al is greater than it between
Mg and Al. Therefore, in alloy solidification,
the rare earth elements is prior to form a newphase with Al, which cost part of the Al
element that should form γ-Mg17Al12 phase,then a lot of rare earth elements are
consumed, so it is hard to find the Mg-RE
phase in magnesium alloy. In addition, from
the Mg-RE phase diagram, we can know thatthesoluteequilibrium
distribution
coefficient K is less than 1(k<1), so the rareearth elements will be enriched in theforefront of solid-liquid interface and thecompounds exist in the branch crystal
network. When the alloy containing RE
curdles,thesolidusandliquidus
temperatures decrease. The enrichment of
rare-earth in the solid-liquid interface leads
to undercooling trend increase, and promotethe nucleation of α-Mg which increases the
crystal growth, secondary crystal branch
growing , and decreases branch arm spacing ,eventually makes the grain refinement. In
addition, as the surface activity elements, it

can lead to enrichment in the solid-liquidinterface, and then plays a role in grain
refinement. Figure 3 is SEM images andenergy spectra of the component 4(Mn =0.68wt%, RE = 0.67%).

Table 3 The electronegativity values of major elements
Element

Mg

Al

Electronegativity

1.31

1.61

From the table 3, we can see theelectronegative difference between Ce and
Al or La and Al is greater than it between
Mg and Al. Therefore, in alloy solidification,
the rare earth elements is prior to form a newphase with Al, which cost part of the Al
element that should form γ-Mg17Al12 phase,then a lot of rare earth elements are
consumed, so it is hard to find the Mg-RE
phase in magnesium alloy. In addition, from
the Mg-RE phase diagram, we can know thatthesoluteequilibrium
distribution
coefficient K is less than 1(k<1), so the rareearth elements will be enriched in theforefront of solid-liquid interface and thecompounds exist in the branch crystal

La

1.12

1.10

network. When the alloy containing RE
curdles,thesolidusandliquidus
temperatures decrease. The enrichment of
rare-earth in the solid-liquid interface leads
to undercooling trend increase, and promote
the nucleation of α-Mg which increases the
crystal growth, secondary crystal branch
growing , and decreases branch arm spacing ,eventually makes the grain refinement. In
addition, as the surface activity elements, it
can lead to enrichment in the solid-liquidinterface, and then plays a role in grain
refinement. Figure 3 is SEM images andenergy spectra of the component 4(Mn =0.68wt%, RE = 0.67%).

a

b

c

d

Fig. 3Ce

The SEM images and energy spectra of the component 4(Mn = 0.68wt%, RE = 0.67%)

In figure 3, the bright field shows
precipitate, and dark field shows Mg-basedsolid solution. Precipitates for the alloy arewith different sizes and shapes of needle-likeor particle-like. Through the spectrum
analysis of three precipitate points in the
figure, the massive particles are mainlycomposed of Al and Mn elements in figure b
and the Mg and O elements diffraction peaksin the spectrum may caused by thediffraction of magnesium and oxidation ofmagnesium. The Al, Mn elements ratio is
close to 8:5 in massive particle precipitation
phase. According to reports in the literature,
this phase is seen as Al8Mn5 preliminary.
In figure c, the phase is mainlycomposed of Mg and Al elements, of which
the O element is caused by oxidation of thealloy. From the atomic ratio of Mg and Al,
and SEM images can determine this initial
phase which is Mg17Al12. In figure d, themain elements are Al, Se and La. The ratioof Al and Ce+La is 4:1. Combined with theliterature [3], such needle-like precipitates
are maybe Al4La and Al4Ce in the alloy.

4.

Conclusions

Ordinary
commercial
AZ61A-castmagnesiumalloy
microstructureiscomposed of α-Mg matrix, γ-Mg17Al12 phaseand a small amount of Al8Mn5 phase, inwhich γ-Mg17Al12 with irregular shape
distributes in the grain boundaries. By
adding Mn, RE elements, as the Mn / RE
ratio in the alloy decreases, of γ-Mg17Al12

phase in cast structure gradually changed:
Large-size dendritic distribution likenetwork
Small dendritic distribution
like network
Semi-continuous network
distribution
Uniform dispersion ofpoint-like. With the increase of Mn element,
Al8Mn5 has a slight increasing in the number
and size. As the addition of rare earth, Al4RE
phase's number and size increase in the grainboundary and intragranular significantly.
Component 4 (Mn = 0.68wt%, RE =0.67wt%) has the best cast structure.

Acknowledgments
The partners acknowledge National
Key Technology R&D Program (Project
Number:
2011BAE22B04-1-3)
andMaterial molding process and mold Key
Laboratory of Ministry of Education forsupporting this research.
References[1] Yasuhiro, Hiramitsu. Tomoyuki, Homma
and Shigeharu Kamado. Improvement ofthemechanicalproperties
ofMg-Gd-Y-Zn alloy castings by grain
refinement[J]. S-GME2011.
[2] 高 岩 , 王 渠 东 , 赵 阳 , 童 炎 , 郑 江 .
Mg97ZnY2 合金中 LPSO 结 构的研究现状[J]. 材料导 报 , 2008, 22(1): 94~97[3] 温 景 林 ． 金 属 挤 压 与 拉 拔 工 艺 学[M]．沈阳：东北大学出版社．1996：
31～33

Enhancement in mechanical and degradation properties of Mg-Zn-Y-Nd alloy bycyclic extrusion compression
Wu Qiong, Guan Shaokang*, Zhu Shijie, Wang Liguo, Hu Junhua
School of Materials Science and Engineering, Zhengzhou University,
Zhenzhou, 450002, China
Tel & Fax: +86 371 67780051*E-mail: skguan@zzu.edu.cn

Abstract
This paper introduces the effect of cyclic extrusion compression(CEC) process on
the mechanical and degradation properties of Mg-Zn-Y-Nd alloy. Results show that thegrain size was greatly refined and the mechanical properties were greatly improved. At
the same time, the homogeneous corrosion were observed in CEC treated samples.
1. Introduction

2. Experimental

Metallic materials including stainlesssteel, titanium alloys, and cobalt-based
alloys have been used as materials for
vascular stent due to their high strength,
ductibility and good corrosion resistance[1].
However, these metallic materials are notbiodegradable in the human body and cause
long term complications. Magnesium alloys
have been attracting growing attention inbiodegradable stent materials because of
their good biocompatibility and degradation
property. But studies show that the early lose
in mechanical property occurred tomagnesium implants along with rapid
degradation in human body. Therefore, it isnecessary to improve the mechanicalproperty and corrosion resistance. Methods
to improve the mechanical properties involve
purification[2], microalloying[3], fluoride
conversion coatings[4], alloying with otherelements and anodizing[5]. In this study,
Mg–Zn–Y–Nd alloy with ultra-fine grain
structuresandnano-sizedparticles
homogeneous distribution was prepared by
CEC. The mechanical property andcorrosion behavior of the CEC treatedmagnesium alloy were investigated. In orderto clearly understand the corrosion behaviorof the CEC treated Mg–Zn–Y–Nd alloy, theas-cast and extruded samples were alsostudied as control groups.

Mg–Zn–Y–Nd alloy ingot was preparedwith high purity Mg, high purity Zn,
Mg–25Y(wt.%, 99.99% in purity) and
Mg–25Nd(wt.%, 99.99% in purity) masteralloys through induction of mild steel
crucible at approximately 740ºC under
CO2/SF6 (volume fraction rate, 3000:1)
atmosphere in an electronic resistancefurnace. The cast bar(Ø60 × 75 mm) cut
from the ingot was extruded at 543 K withan extrusion ratio of 25 and then cooled inair. The samples(Ø30mm×55 mm) used for
CEC were machined from ingot.
The schematic illustration of CEC
facility was shown in Fig.1[6]. The numberof extrusion passes was defined as thenumber of specimen passes through the crosssection neck die. In the present study, thesample was extruded 2 and 4 passes. D0 andd are 30mm and 18mm, respectively. The die
was lubricated using supramoly andpreheated to 523K before processing. The
extrusion temperature was 543K and the ram
speed was 4 mm/s. During the final
extrusion pass, the opposite ram B was
removed in order to release the rod. The
specimen of CEC treated is shown in Fig.2.

were improved significantly while theelongation was not improved apparently
after hot extrusion. However, the elongationof CEC treated samples was about 2 times
higher than that of as-cast alloy and wasabout 1 time higher than that of extruded
alloy. This is related to the grain size, second
phase distribution, dislocation density, andgrain boundary structure of the alloy.
Tab.1 The tensile properties of the samples
UTS

YS

Elongation

(MPa)

(MPa)

(%)

As-cast

209

105

10.6

Hot extrusion

316

183

15.6

CEC-2-CT

303

185

30.2

CEC-4-CT

345

193

20.7

Samples

Fig1 Schematic illustration of CEC facility andprocedure: (a) Initial state; (b) Extruding by ram
A; (c) End of ram A involvement; (d) Reverse
extruding employing ram B; (e) End of ram B
involvement

Fig.3 shows the corrosion potential
curves (a) and the polarization curves (b) of
the specimens after immersion in SBF. It canbe seen that the corrosion potential of CEC
treated alloys shifted to a more positive
value than that of as-cast and extruded ones,
and the corrosion current density of the alloy
decreasedfrom
2.8×10-4A/cm2
to2
6.6×10-5A/cm . The reason is that thegrain size of alloys was strongly refined
during the CEC process, and the small
grain size creates more grain boundaries
that act as a corrosion barrier to increase
the corrosion resistance.

Fig2 The specimen of CEC treated alloys

3. Results and Discussion
The obtained values of ultimate tensile
strength(UTS), average yield strength(YS)
and elongation(δ) are summarized in Table1.
It shows that the UTS and YS of the alloy

Fig.4 shows the morphology of the
Mg–Zn–Y–Nd alloy after immersion in SBF
for 48 h. As shown in Fig.4(a) and (b),as-cast and extruded alloys suffered fromapparent pitting corrosion. However CEC
treated alloys showed the trend of uniform
corrosion. Uniform film layer weregenerated on the surface of the alloys. The
corrosion cracks were fewer than that of as
cast and conventional extrusion, fewer
cracks means the film layer is more compact
which results in better corrosion resistance.

(a)

-1.62

-1.60

As-castHot extrusion
CEC-2-CT
CEC-4-CT

-1.64
-1.66
-1.68

-1.62

-1.70

Potential / V

(b)

-1.64

Potential (V vs.SCE)

-1.60

-1.72
-1.74
-1.76
-1.78
-1.80
-1.82
-1.84

-1.66
-1.68

As-castHot extrusion
CEC-2-CT
CEC-4-CT

-1.70
-1.72
-1.74
-1.76
-1.78
-1.80
-1.82

-1.86

-1.84

-1.88

-1.86
-1.88

-1.90
0

500

1000

1500

2000

Time/s

2500

3000

3500

400

-1.90
-7

-6

-5

-4

log current density ( A/cm2)-3-2

Fig 3 (a) The corrosion potential of samples via time; (b) Potentiodynamic polarization curves ofdifferent samples in SBF solution at 36.5 ± 0.5℃.

(a)

(b)

(c)

(d)

Fig 4 SEM images of (a) as-cast, (b) hot extrusion, (c) CEC-2-CT and (d) CEC-4-CT samples after
immersion in SBF solution at 36.5 ± 0.5℃ for 24h

4. Conclusion
After CEC treatment the mechanicalproperties of Mg–Zn–Y–Nd alloy weresignificantly improved. The UTS wereachieved 303 MPa after CEC 2 passes
treatment, and the elongation of alloys
reached to 30.7%, which can be ascribedto grain refinement and precipitation
strengthening. Mostly, the CEC treatedsample exhibited uniform corrosion due tothe grain refinement, grid second phase
distributed at grain boundaries andnano-sized particles distribution in grains,
and the internal stress were relieved for GBS.
However, the as cast and extruded alloy
suffered from pitting corrosion. So the CEC
processing will be a promise candidate
process for the vascular stent.

Reference[1] Mani, G., Feldman, M.D., Patel, D.,
Agrawal, C.M., 2007. Coronary stents: amaterials perspective. Biomat. 28,
1689–710.
[2] Song, G.L., 2007. Control ofbiodegradationof
biocompatible
magnesium alloys. Corro. Sci. 49,
1696–701.
[3] Hänzi, A.C., Dalla Torre, F.H.,
Sologubenko,
A.S.,
Gunde,
P.,
Schmid-Fetzer, R., Kuehlein, M., Löffler,
J.F., Uggowitze, P.J., 2009. Design
strategy for microalloyed ultra-ductile
magnesium alloys. Philos.Mag. Lett. 89,
377–390.
[4] Hänzi, A.C., Sologubenko, A.S.,
Uggowitzer, P.J., 2009. Design strategy
for new bioabsorbable Mg–Y–Zn alloys
for medical applications. Int J Mater Res.
100, 1127–1136.
[5] Chiu, K.Y., Wong, M.H., Cheng, F.T.,
Man, H.C., 2007. Characterization andcorrosion studies of fluoride conversion
coating on degradable Mg implants. Surf
Coat Technol. 202, 590–8.
[6] Wu, C.S., Zhang, Z., Cao, F.H., Zhang,
L.J., Zhang, J.Q., Cao, C.N., 2007. Studyon the anodizing of AZ31 magnesium
alloys in alkaline borate solutions. Appl
Surf Sci. 253, 3893–8.

The morphology and property of the PPY coated electrode after silane coupling
agent treatment
ChenGuang1, GuanShaokang1*, WangLiguo1, ZhuShijie1, RenChenxing1
School of Materials Science and Engineering, Zhengzhou University,
Zhenzhou, 450002, China
Tel & Fax: +86 371 67780051*E-mail: skguan@zzu.edu.cn

Abstract
In this paper, silane coupling agent was applied to pre-treat on stainless steel wire, then
PPY/TSO- film was fabricated on it by electrode position process. SEM, EIS, and CV, wereused to present the morphology and electrochemical properties of coating. The results indicatedthat the morphology of the coating changed after pretreatment in different current density,
resistance decreased to 1 kHz, and QH increased significantly.

1. Introduction
In recent years, polypyrrole (ppy)preparation techniques have been greatly
developed [1-5]. As a conductive polymer
material
inbattery
industry[6],
supercapacitors [7-8], pollution control [9],ppy has broad application prospects. In
addition, ppy as a new biological material in
biological sensors [10-11] has also
been extensively studied. But the surface
morphology of ppy which prepared byelectrochemical method looks loose andmessy, and the ppy coating showed poor
adhesion strength with the stainless steel
substrate [12]. Moreover, electrochemical
impedance(in
1KHZ)
and catholic
charge reservesare are also
generally not satisfactory. Therefore, it isnecessary to improve the electrochemical
property and morphology. Typically, the ppy
produced, with PH = 1 and the current
density of about 1mA/cm2, showed best
performance.
In this study, SS wire was used as
substrate material, with its surface pretreatedby silane coupling. Then ppy was preparedon the surface at current density of 0.5
mA/cm2, 1 mA/cm2 and 1.5 mA/cm2
respectively in py/Tso- solution. The
electrochemical property and morphology of
the SS wire with ppy coating were
investigated. In order to illustrate theinfluence on electrochemical property andmorphology after silane pretreatment, thesamples without pretreatment were alsostudied as control groups.

2. Experimental
Stainless steel wire (Ø127µm) withinsulating barrier was used as the substratematerial. Both ends of the electrode wire
insulating barrier was stripped; working
parts's stripping length is 5 mm. The otherend of the electrode wire welding together
with copper wire. Finally, with silicon rubber
sealing up the solder joints. The specimen isshown in Fig.1

Fig. 1 Sample structure schemes

0.5 ml of silane and 9.5 ml of water weremixed and stir 1-2 h to make silane fullyhydrolysis. The working part was immersed
in silane coupling agent solution, for about
two minutes. This process was repeated twoor three times, to get even silane film. In theFig.2 Schematic diagram of ppy film preparation.

electrochemical tests, the samples wereusedasworking
electrode, withplatinum plate used as counter electrode andsaturated calomel electrode (SCE) used as
reference electrode. Then, ppy was preparedon electrode surface using constant current
method in the py/Tso- solution at 0.5 mA/cm2,
1mA/cm2 and1.5 mA/cm2, respectively.
The experiment process is shown in Fig.2.

3.

The morphology of the ppy/Tso- preparedon the pretreated SS surface by three current
density was compared with the control groupin Fig.3. The protrusions of two groups areboth grew up as the current density increases.
But the morphology of the experimentalgroup have more flat surface than the control
group. Additionally, the jagged protrusions
were changed to granular shape.
Fig.4
shows
the
electrochemical
impedance spectra of two groups in Artificial
Cerebrospinal Fluid (ACSF). It can be seenthat the impedance (in 1KHZ) of theexperimental group prepared in threedifferent current density are all lower than thecontrol group.

Results and Discussion

Fig 3 SEM images of the ppy/Tso-film (a), (b) , (c) and the ppy/Tso-film after silane pretreatment (d), (e), (f)The CV curve was obtained by cyclic
voltammetryusingelectrochemical
workstation in ACSF. The integration of CV
curves is CSCH. Tab.1 shows the CSCH of

two groups in different current density.
Obviously, compared to the control group,
experimental group which pretreated bysilane coupling agent have higher CSCH.

Bareelectrodeppycoatedelectrode
SCA/ppycoatedelectrode2
j=0.5mA/cm

5b

ppycoatedelectrode
SCA/ppycoatedelectrode2
j=1mA/cm

5

10

4

10

1492Ω
3

10

598Ω
542Ω
2

10

1

10

4

10

3

434Ω

10

2

10

358Ω

2

3

10

4

10

10

4

10

3

10

237Ω

2

10

207Ω

1

10

5

10

ppycoatedelectrode
SCA/ppycoatedelectrode2
j=1.5mA/cm

5

10

10

1

1

10

c
Impedance magnitude Z/ohm

Impedance magnitude Z/ohm

10

Impedance magnitude Z/ohm

a

1

2

10

3

10

10

Frequency/HZ

4

10

5

10

1

10

2

10

3

10

4

10

5

10

Frequency/HZ

Frequency/HZ

Fig. 4 The electrochemical impedance of different samples (a, 0.5 mA/cm2, b, 1.0 mA/cm2, c, 1.5
mA/cm2).
J

0.5

1

1.5

19.6

40.2

54.4

36.6

67.6

105.2

CH
Type
Without silane
pretreatment
After silane
pretreatment

Tab.1 The CSCH of different samples.

4.

Conclusion

The morphology of the coating waschanged after pretreatment in different
current density, the morphology changed
from jagged to granular shape. The
electrochemical impedance of ppy/Tso- film
after silane pretreatment decreased (1 KHZ)
15% in average compared with the control
group. The cathode charge of PPY/TSO-film
coated electrode after silane pretreatment had
relatively substantially risen.

Reference[1] C.X. The Development of conductive
plastics [J]. New Material Industry, 2002
(1):28-34.
[2] L.R., L.X, Wang, et al. Conductive
polymers andconductive polypyrrole Research[J].
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[3] Y.Y. Zhou ， M.Y ， S.Y. Conductive

polymer polypyrrole Research. Materials
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Properties of Porous Polypyrrole Thick
Films Prepared Electrochemically byMulti-step
Polymerization.
Actaphysico-chimica
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[6] Y.Q. Chen, D.Y. Pan Cheng Xian-su XU
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Electronic
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[7] F. Ran, G.J. Dang, L.B. Kong, et al.
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Sequence feeding preparation andsupercpacitive
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ypyrrole coated Fly Ash C enospheres.
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genolelectronics:

electrochemical
DNA
Biosensors.
Science&Technology
review.
2009;27(11):93-101.
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Biosensors Based on nanostructured
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filmsBiocorrosion of coated Mg-Zn-Ca alloy under constant compressive stressclose to that of human tibia
Bin Wang, Junheng Gao, Liguo Wang, Shijie Zhu, Shaokang Guan*
School of Materials Science and Engineering, Zhengzhou University,
97Wenhua Road, Zhenghzhou 450002, P.R. China*E-mail: skguan@zzu.edu.cn

Abstract
A self-designed immersion method simulating in vivo conditions more accurately thanconventional ones was used to investigate biocorrosion of coated Mg-Zn alloys for implants.
Results showed that the applied stress can change the degradation mode of coated samples andpromote corrosion process, but didn’t affect biocompatibility of the coated samples apparently.

1. Introduction
Magnesium (Mg) alloys are promising
candidates for degradable implants [1-3] dueto their favorable biocompatibility and wellmatched mechanical properties with humanbone. To solve their too rapid degradation
rate in physiological environment, various
techniques including surface modifications[2-3] are developed. However, because
implants degrade under stress when they
serve inside human body, the most common
immersion measuring biocorrosion without
considering the influence of stress [1-3] may
not be so precise in predicting thedegradation behavior of Mg alloy implants.
Therefore, it’s essential to study the
corrosion behavior of implants undercompressive
stressin
physiologicalenvironment, and biocorrosion of Mg alloys
under constant compressive stress close tothat of human tibia in SBF was firstly
investigated in this paper.

2. Experimental
All samples were from the substrate (castMg-2Zn-0.24Ca alloy (wt%)) and weredivided into three groups for studies: thesubstrate alloys without compressive stress instatic immersion; coated [2] alloys withoutcompressive stress in static immersion;coated alloys under compressive stress using
self-designed method, illustrated in Fig.1, theforce was adjusted to make sure that thecompressive stress applied to samples was

4.0MPa, close to the average stressshouldered by human tibia, 4.047±0.217MPa[4-5]. The coating consists of an inner
microarc oxidation layer and an outer nanohydrooxyapatite layer[2]. Kokubo’s SBF waschosen, and the immersion period was 5 days.
Corrosion morphologies of the samples werestudied by Quanta-200 scanning electron
microscope (SEM) equipped with energy
dispersion spectroscopy (EDS) facility, andpH values of the solutions were recorded
during the immersion at different intervals.

3. Results and Discussion
Fig.2a-b, c-d, e-f show the different
morphologies and elemental analysis of thethree groups of samples after 5 days’
immersion respectively. In Fig.2a-b, apparent
cracks, peeling-off parts and corrosion
products containing Ca, P and Mg, wereobserved on the substrate. Whereas in
Fig.2c-d, the coating was intact, andrandomly scattered osteoconductivitive Ca-P
particles were observed, demonstrating thatthe coating has improved corrosion resistanceand biocompatibility of the substrate. In
Fig.2e-f, microcracks, a few peeling-off
locations, globular Ca-P salt particles anddeposition products around some pores wereobserved. The main compositions of
peeling-off locations were major elements ofMAO
layer,
indicating
that
under
compressive stress, certain parts of outer HAlayer may fall off and the inner MAO layer
was exposed to SBF. Therefore, compressive

stress promoted the degradation of coating,but the coating still has effects improving thecorrosion resistance and osteoconductive
ability of the substrate.
Fig.3 shows pH variations of the samples
over five days immersion. It’s clear that pH
value of the substrate rose fastest and highest,
indicating fastest corrosion, followed bycoated samples under stress and then coatedsamples. For both the substrate and coatedsamples in static immersion, pH value
increased during the first stages, and thendecreased, exhibiting similar corrosion
process; however, pH value of the coatedsamples under stress rose all through thethree stages, indicating a different
degradation mode. By referring to corrosion
morphologies in Fig.2 and relevant research[6], it can be deduced that the appliedcompressive stress may promote the initiationof microcracks in the coating and assist thecrack propagation together with the corrosive

Fig.1. Schematic diagram of immersion
method for coated samples bearing compressive
stress.

effects of SBF, thus leading to the peeling-off
of outer HA layer around certain pores andmaking the inner part of the coating or eventhe substrate exposed to the solution.
Therefore, the coating degrades and thesubstrate corrodes gradually, generating more(OH)- and increasing the pH value.

4. Conclusion
The self-designed immersion method
including compressive stress close to humantibia provides essential information forpredicting in vivo degradation behavior ofcoated Mg alloy implants. The stress caused
certain peeling-off locations on the coating
and continuous increase of pH value.
Meanwhile, Ca-P salts precipitated inside themicropores and on the surface still showed
protective and osteoconductive ability of thecoating.

Fig.3. pH variations of the samples as afunction of immersion time.

Fig.2. SEM & EDS results of the samples after immersion for 5 days: (a) and (b) for the substrate; (c)and (d) for coated samples without compressive stress; (e) and (f) for coated samples with compressive
stressReferences[1] Gao JH, Guan SK, Ren ZW, Sun YF, Zhu
SJ, Wang B. Mater Lett 2011;65:691-3[2] Gao JH, Guan SK, Chen J, Wang LG, Zhu
SJ, Hu JH, et al. Appl Surf Sci
2011;257:2231-7[3] Li HF, Wang YB, Cheng Y, Zheng YF.

Mater Lett 2010; 64:1462-4[4] Fukubayashi T, Kurosawa H. Acta Orthop
Scand 1980;51:871-9[5] Wang F, Zhang Q, Hebei Medical Journal
2006; 28:83-5[6] Kannan MB, Dietzel W, Blawert C,
Atrens A, Lyon P. Mater Sci Eng A
2008;480: 529-3