Functional properties of citric acid capped TiO2 nanoparticles by hydrothermal
growth and dye-sensitized solar cell performance
J. Archana1, 2, M. Navaneethan2, and Y. Hayakawa1, 2*
1 Graduate School of Science and Technology, Shizuoka University
2 Research Institute of Electronics, Shizuoka University
3-5-1 Johoku, Naka-Ku, Hamamatsu, Shizuoka 432-8011, Japan.

Abstract
Anatase TiO2 nanoparticles are successfully synthesized via a simple hydrothermal
method with citric acid as a capping agent. The effects of systematic growth periods onthe morphological, structural and optical properties of TiO2 nanoparticles areinvestigated. XRD results confirm the formation of anatase phase while citric acid wasused. TEM measurement reveals that the particle size increased by increasing the
growth period. TiO2 particles synthesized under different growth periods such as 5, 15,
25 and 45 h are used to prepare a photoanode layer by spray deposition technique fordye-sensitized solar cell fabrication using N719 ruthenium as a sensitizer. It is foundthat the maximum efficiency () of 7.66 % is achieved for 15 h growth period and is
attributed to enhanced light harvesting caused by absorption of greater numbers of dye
molecules.
1. Introduction
Dye-sensitized solar cells (DSSCs)
have been considered as an alternative tosemiconductor solar cells due to their
good potential and cost effectiveness [1].
It is known that the DSSCs consist of
FTO substrate, photoanode material fordye absorption, platinum counterelectrodeandtheelectrolyte(iodide/tri-iodide). There are numerous
factors which determine the efficiency.
Among those, the photoanode isconsidered to be an important factor forthe light harvesting and charge transfer
properties. In DSSCs, the semiconductoroxide material with the wide band gap isused as the photoanode material [6].
Titanium-di-oxide (TiO2) gained goodattention due to their unique properties
such as well matched band alignment
with dyes [7]. It is regarded as apromising material preferred as aheterogeneous photo catalyst in solar
cells. TiO2 exists in three crystalline
polymorphs such as rutile, anatase andbrookite. Among those, rutile is the moststable phase, where as anatase andbrookite are in metastable phases.
However, the anatase phase has beenhighly employed in wide applications
such as DSSCs, photo catalysts, sensors
etc. In order to synthesis TiO2
nanoparticles several methods such as solgel, hydrothermal, laser ablation,
solvothermal
wereadopted.
Incomparison with the other methods,
hydrothermal method is a simple andinexpensive method to prepare well
crystalline materials. However, the fasthydrolysis process leads to the formationof irregular phase and morphology. In
order to synthesize the nanoparticles
without agglomeration, it is necessary to
use the capping agent. It is reported thatthe carboxylic acids have strong affinity
towards TiO2 material. The carboxylic
acid with a long hydrocarbon chain isconsidered as an important surfactant forthe synthesis of titania nanoparticles. It isvery important to identify the capping
ligand which offers unique size reduction
and better morphology.
In the present work, we reportsynthesize of anatase TiO2 nanoparticles
by facile hydrothermal method, using
citric acid as a capping agent. The
systematic investigations are carried out
for the growth period and the functional
properties.
The
photoanodes
arefabricated using spray technique and
DSSC performances are studied.
2. Results and Discussion
Figure 1(a) depicts the XRD patternof the samples at different growth periodsof 5, 15, 25 and 45 h. The phasecompositions of all the samples wereidentified from the XRD pattern. All thediffraction peaks were indexed to (101),(004), (200), (105), (204), (116), (220)and (215) planes of the crystal structureof anatase TiO2 phase and they matchedwith card (JCPDS: 21-1272). Whereas
the uncapped material shows other planes
such as (110) (101) and (211) whichcorresponded to the rutile phase andmixed with anatase phase. XRD patterns
evident that the role of citric acid as aphase directing ligand to obtain only theanatase phase.
Fig.1(b) Raman Spectrum of TiO2
nanoparticles with growth periods of 5, 15,
25 and 45 h.

Further confirmation for the electroniclevels of the samples was analysed by
X-ray photoelectron spectroscopy (XPS).
The binding energies obtained in the XPS
analysis were corrected by reference to
C1s at 284.60 eV. Figure. 2 represents the
XPS spectra obtained from Ti and O
regions of TiO2 nanoparticles. Figure 2(a) shows two strong peaks at 459.5 and464.9 eV which correspond to thebinding energies of Ti 2p3/2 and Ti 2p1/2.
Fig.1(a) XRD pattern of TiO2 nanoparticles
with growth periods of 5, 15, 25 and 45 h.

Figure 1(b) illustrates the Raman
spectra of the prepared samples.
Generally, the anatase phase has six
fundamental vibrational modes such as[A1g + 2 B1g + 3 Eg] and the rutile phasehas four fundamental vibrational modes
such as [A1g + B1g + B2g + Eg]. The citricacid capped nanoparticles grown at
different growth period have the four
Raman peaks at 148, 400, 158 and 649
cm-1 which can be assigned to Eg, B1g,
A1g and Eg modes of the anatase phase[2]. The uncapped material has theRaman peaks at 237, 400, 450, 518 and640 cm-1. Where the peaks 237 and 450
cm-1 can be assigned to Eg modes of therutile phase and the remaining lines
belong to the anatase phase as mentioned
above. Thus it confirms that the uncappedmaterial has the mixture phase of anataseand rutile. It has good agreement with the
XRD data.

Fig.2(a) XPS spectra of TiO2 nanoparticles
with growth periods of 5, 15, 25 and 45 h.

In Figure 2 (b), there was a strong
peak at 530.8 eV, which is attributed tosignature of the lattice oxygen O1s in the
Ti-O-Ti bonds. All the samples exhibitedthe similar peak values in the Ti and O
core level spectra. No obvious peaks forother elements of impurities wereobserved.

Fig.2 (b) XPS spectra of TiO2 nanoparticles
with growth periods of 5, 15, 25 and 45 h.

is clear that the Voc and FF show
somewhat constant whereas the Isc shows
an increasing behavior from 12.02 to
16.59 mA cm-2 when the particle size
increased from 8 - 11 nm (5 h) to 12 - 18
nm (15 h). When the particle size
increased as 24 - 29 nm (25 h) and 29 –37 nm (45 h), the Isc started to drop from16.59 to 13.44 and 12.16 mA cm-2. The
overall conversion efficiency () shows
the similar behavior of Isc thus theefficiency increased from 5.44 to 7.66%(8 – 11 nm (5h) to 12 - 18 nm (15 h))and it decreased as 6.45 and 5.61 %. It isclear that the DSSC performance ishighly dependent on the Isc factor. When
compared with the above data, the growth
period 15 h is optimized to yield the
maximum efficiency 7.66 % with theparticle size of about 12 - 18 nm. The
decrease in the efficiency as the particle
size increased may be due to the minimalof surface area for the greater absorbanceof dye molecules.
3. Conclusions

Fig.3 Illustration of I-V characteristics

and photocurrent spectra of TiO2
nanoparticles at growth periods of 5, 15,
25 and 45 h.
Table.1 Device parameters of DSSCThe DSSCs were fabricated using thenanoparticles synthesized at various
growth periods. Figure 3 shows thecurrent density versus voltage (I-V)
characteristics measured for 5, 15, 25 and45 h. The values of Voc, Isc, FF andconversion efficiencies () of the DSSCs
are listed in the Table.1. From the table it

TiO2 nanoparticles have beensuccessfully synthesized using facilehydrothermal method. The effect of citricacid on TiO2 nanoparticles has beenstudied. The functional properties of the
TiO2 nanoparticles were investigated by
TEM, XRD, Raman spectroscopy, UV-vis
spectrophotometery, FTIR spectroscopy
and XPS analysis. Citric acid promotes
the formation of anatase phase throughthe coordination of carboxylic groups
with the titanium complexes. The effectsof various growth periods (5, 15, 25 and45 h) have been investigated. The effectof the photoanode (with various growth
periods) on the DSSC conversion
efficiency was investigated. It was foundthat the maximum efficiency () of
7.66% was obtained for 15h growth
period.
Acknowledgments
The authors would like to thankProf. K. Murakami of the Nano-Device
Process Laboratory, for extending his
support to the I-V measurements. Authors
also
thank
Mr.T.Koyama
andMr.W.Tomada for their assistance during
the measurements. J. Archana would liketo thank MEXT, Japan, for their award ofa research fellowship.
References[1] O’Regan. B, Gratzel.M. Nature, 353(1991) 737.
[2] T.Ohsaka,, J. Phys. Soc. Jpn. 48(1980) 1661.
[3] P.Thistlethwaite, MSHook, Langmuir
16 (2000) 4993.

Synthesis of MoSi2 powder by Molten salt Technique using by MoS2
Daisuke Ishikawa
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu 432-8561, Japan

Abstract
The hexagonal phase MoSi2 was grown using MoS2 powder as the source material
by the molten salt technique. The structural and morphological properties of theresultant silicides are characterized. The layered structure of MoS2 with a hexagonal
crystalline structure would affect the growth morphology and enhance the formation of
the hexagonal MoSi2. Moreover, Synthesis of the h-MoSi2 was promoted by adding S
atom. A simple growth procedure to fix the crystalline structure using an appropriate
starting material under a preferable growth environment is proposed.
1. Introduction
Recently,
semiconducting
silicides,
which consist of non-toxic and abundantmaterials, have attracted much attention for
their potential to create new classes ofenvironmentally conscious electronics [1].
The group VI metal silicides are classified inthis material group [2]. MoSi2 with stable
tetragonal (t-) phases is metallic and has
been used in electrode wire materials. On theother hand, the hexagonal (h-) phase MoSi2
has lattice constants of a=0.4622 nm and
c=0.6646 nm, and is known to be a narrowgap semiconductor. It is reported that the
low temperature phase, h-MoSi2, is expected
to be one of the novel silicide
semiconductors, which has improved
thermoelectric properties [2]. However, theelectrical and thermoelectric property of thebulk crystal h-MoSi2 have not been
experimentally clarified. Since t-MoSi2 istransformed into h-MoSi2 with the Ti
addition by a Mechanical Alloying (MA)technique in the Mo-Ti-Si system [3], it has
been reported that Mo-silicides were grownusing Mo substrates and Mo-Ti alloy(Mo(Ti)) compacts by the molten salt
method. However, the tetragonal phase wasformed as well as the hexagonal phase in thesilicides. A nonhomogeneous compositional
distribution of Mo and Si atoms wasobserved in the silicides when using the Mo
substrate and Mo(Ti) compact [4]. The
single phase growth of h-MoSi2 andimproved compositional homogeneity of thesilicide are required.
On the other hand, MoS2 has a hexagonal
structure with lattice constants of a=0.31612
and c=1.22985 nm. The detailed structure of
MoS2 and extraordinary chemical propertiesof the two-dimensional MoS2 nanoclusters
are described in Refs.5 and 6, respectively.
MoS2 has a characteristic layered structure,i.e., metal chalcogenide slabs are formed by
two layers of close-packed chalcogenide
atoms sandwiching one metal layer between
them. These slabs are stacked with only van
der Waals forces between the slabs [7]. It isconsidered that the use of the thin separated
slabs would improve the homogeneity of theatomic distribution in the resultant alloys,
because the diffusing species easily can passthrough the materials during the treatment.
Moreover, it has also been reported that the
crystalline structure of the resultant
compounds is strongly affected by that of thestarting materials. For example, theformation of the zincblende-MnTe phase
was demonstrated using ZnTe [8]. Single
phase Ca2Si was obtained using Mg2Si as thestarting material [9]. In the same manner, it
is expected that the h-MoSi2 would be easily
obtained using MoS2 as the source material.
In this study, the synthesis of h-MoSi2
was examined using MoS2 powder toovercomethepreviously
mentioned
difficulties.
The
structuralandmorphological properties of the resultant
silicides were also characterized.

Fig.1 XRD spectra of the resultant powder byusing MoS2 as the source material at 600 ºC.

Fig.2 XRD spectra of the resultant powder byusing MoS2 as the source material at 700 ºC.

2. Experiments
The Mo-silicide powder was synthesizedusing MoS2 powder, S powder and a molten
salt, comprised of 36.58 mol%NaCl - 36.58
mol%KCl - 21.95 mol%NaF - 4.89
mol%Na2SiF6. The salt mixture, silicon
powder (99.999%, 21.85 mol% for the salt
mixture), S powder (99.99 %, 10 mol% forsalt mixture) and MoS2 powder (99.9%up,
2.73 mol% for salt mixture) were placed in
an SiO2–Al2O3 crucible with the powder.
They were thermally treated at 873 and973K. After the heat treatment, the salt wasremoved from the sample using deionized
H2O. The structural properties of the powder
were characterized by scanning electron
microscopy (SEM) with energy dispersion
spectroscopy (EDS), transmission electron
microscopy
(TEM)
and
scanning
transmission electron microscopy (STEM)
with EDS.

3. Results and Discussion
Figure.1 shows XRD spectra of theresultant powder by using MoS2 as thesource material at 600 ºC. The h-MoSi2
was formed with progress of growth time.
Figure.2 shows XRD spectra of theresultant powder by using MoS2 as thesource material at 700 ºC. The h-MoSi2
was formed with progress of growth time,
however, the t-MoSi2 formed as well as

h-MoSi2 with progress of growth time. The
stable t-MoSi2 was formed at higher
temperature.

SourcePotision
Material
MoS2

1
2
3

Composition[at.%]
Mo
Si
25.16 68.73
17.66 74.92
~100

S
6.11
7.43
-

Fig.3 SEM image and corresponding EDS
mappings of resultant powders grown using
MoS2 as the source material at 600 ºC / 4h.

Figure.3 shows an SEM image andcorresponding EDS mappings of theresultant powders grown using MoS2 as thesource material. The average composition of
the local area on the resultant silicide isapproximately Mo : Si ≈ 1 : 2~4. The region
with pure Si is also observed. According to
the Mo-Si phase diagram, no stable phase
exists between MoSi2 and Si [10]. The
excess Si is due to the presence of Si powder
as the source material.

respectively. Synthesis of the h-MoSi2
promoted by adding S powder.

Elements
Mo
Si
S

was

Composition [at.%]
12.82
21.46
0.36

Fig.5 STEM image and corresponding EDS
mappings of one of the powders grown using
MoS2 as the source material.

Fig. 4(a) TEM image and (b) an enlarged
lattice image of the resultant powder using
MoS2 as the source material at 600 ºC / 4h. The
FFT pattern of the images and corresponding
crystalline orientations of the domains areshown in the inset.

Figure.4 shows a TEM image and an
enlarged lattice image of the resultant
powder using MoS2 as the source material.
The Fast Fourier Transformation (FFT)pattern of the images and corresponding
crystalline orientations of the domains areshown in the figure. As shown in Fig.4(a),thin silicide films were synthesized, which isdue to the layered structure of the MoS2
starting material. The enlarged lattice fringes
show a hexagonal symmetry and the latticespacing of the fringe is 0.23 nm whichcorresponds to that of h-MoSi2(110) as
shown in Figure.4(b). The six-fold FFTpattern is consistent with the electron
diffraction patterns of the h-MoSi2 observed
along the [0001] zone axis.
Figure.5 shows an STEM image andcorresponding EDS mappings of one of thepowders grown using MoS2 as the source
material. The average chemical compositionof the powder is about Mo:Si=1:2, though asmall amount of S (~0.4%) remained in the
powder.
Figure.6 and 7 shows XRD spectra of theresultant powder by using S and MoS2 as thesource material at 600 and 700 ºC,

Fig.6 XRD spectra of the resultant
powder by using S and MoS2 as thesource material at 600 ºC.

Fig.7 XRD spectra of the resultant
powder by using S and MoS2 as thesource material at 700 ºC.
.

4. Conclusion
The h-MoSi2 was formed using MoS2
powder as the source material by the molten
salt method. Moreover, thin silicide firms
were synthesized, and the composition was
uniform. The layered structure of MoS2 withhexagonal crystals would affect the growthmorphology and enhance the formation of
h-MoSi2. Synthesis of the h-MoSi2 waspromoted by adding S atom. It is expectedthat this simple growth procedure to fix the
crystalline structure using the appropriate
starting material under the preferable growthenvironment would open a new field ofsilicide semiconducting technology.

Acknowledgements
I would like to thank W. Li and S. Oda
and T. Matsushita and K. Nakane and H.
Tatsuoka for their fruitful discussions.

Reference[1] APAC-SILICIDE, July 29-31, 2006,
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Yamaguchi and K. Akiyama.
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Y. Hayakawa and H. Tatsuoka, IOP Conf.
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Clausen, and H. Topsoe, J. Catal. 187,
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M.V. Bollinger, B.S. Clausen, H. Topsoe,
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620 (2003).
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Phase Equilibria 12, 493(1991).

Synthesis of calcium – Tin system composites by Thermal Annealing
Keisuke Kanda
Faculty of Engineering, Shizuoka University
TEL&FAX : 053-478-1099
E-mail : f0911042@ipc.shizuoka.ac.jp

Abstract
Nanostructured calcium – tin system composites were synthesized by thermal annealing.
These shapes of nanostructures grown on surface of composites were different each other.
Additionally, formed composites had reactivity against moisture. The structural properties of
the resultant composites were characterized by scanning electron microscopy (SEM)

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
The composites were synthesized fromSn lumps (Alfa Aesar, -100 mesh, 99.999%)
and Ca 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 details of experimental
conditions are described in Table 1. Eachcrucible was sealed in a stainless-steel tube
(inside diameter, 13mm; length, 80 mm) toprevent the vaporization of Ca. The sealed
tubes were heated at 500ºC for 4 hours in anelectric furnace.
The structural properties of the resultant
composites were characterized by scanningelectron microscopy (SEM).
Table 1 Preparation conditions

Composite#1

Molar ratio(Na : Sn)
1.0 : 1.0

3. Results and Discussion
Figure 1 shows the SEM image ofComposite #1.
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
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straight, and have sharp top end, were grownin composite #1. On the other hand, the rods,
which have an unhomogeneous shape and arugged surface, were grown in composite #2.
The synthesis condition of Na-Sn systemcomposites is clarified by characterization of
the cross-section of composite #1 and #2.
Fig.1 SEM images of Composite #1 shown in
Table1.

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).
While both composites were stored, bothcomposites reacted with moisture. Therefore,
it seems that peaks of compound having ahydroxy group and peaks of monohydratecompound appeared. In other words, it isconsidered that the nanostructures on thesurface of composites were formed by areaction with the moisture.
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.

4. Conclusion
Nanostructured sodium – tin systemcomposites were obtained by simple thermal
annealing method. Both composites had
reactively against moisture. Additionally,
different-shaped
nanostructures
weresynthesized on the surface of bothcomposites. The wires, which are thin,

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(1928) 947[7] J. T. Vaughey, J. D. Corbett, Inorg.
Chem. 36 (1997) 4316
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Chem. 110 (1998) 1641; Angew. Cem.
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42 (2003) 5474
[10] F. Dubois, M. Schreyer, T. F. Fassler,
Inorg. Chem. 44 (2005) 477

Synthesis of carbon based materials from organic wastes fordye - sensitized solar cells
R. Karthikeyan1, M. Navaneethan2, J. Archana1, M. Arivanandhan2 and Y. Hayakawa2*
1
Graduate School of Science and Technology, Shizuoka University
2
Research Institute of Electronics, Shizuoka University
3-5-1 Johoku, Naka-Ku, Hamamatsu, Shizuoka 432-8011, Japan.
*Email: royhaya@ipc.shizuoka.ac.jp
Abstract
Carbon materials were synthesized from the wastes of organic materials such as corn cob andlemon peel by a simple chemical route. XRD and Raman data indicated that thesynthesized materials were related to graphitic based structure with disordered grains.
Compared to lemon peel derived carbon, corn cob based carbon material had moreordered structure. FTIR spectra and elemental analysis concluded that the material had
oxygen and hydrogen.
1. Introduction
Dye-sensitized solar cells are the third
generation electrochemical solar cells, which
consist of four different major parts i.e., thephoto anode, dye, electrolyte and photo
cathode. The typical cell arrangement isshown in Fig. 1. The principle of DSSC’s
was discussed in 1991 by Gratzel et. al [1].
Photoelectrodes consist of nanostructured
TiO2 particles which are mainly used forphotovoltaic conversion. The mechanism of
DSSC’s is that the dye molecules adsorbed
on TiO2 are exposed to sunlight, as a result
photo electrons are generated and are injected
into the conduction band of photo anode. The
oxidized dye is then reduced by the Iodide(I-) / tri-Iodide (I3-) electrolyte. Counterelectrodes or photo cathodes are utilized forthe supply of an electron to the tri-Iodide [I3-]in the reduction process. This cycle should
proceed continuously. The overall efficiency
depends on every section of DSSC’s [2].
High catalytic activity and high cyclic
performance depend on the material at thecounter electrode. Commercially platinum isused as a best photo cathode. But, in terms of
cost and stability, variety of research focuseson the replacement of platinum material
without affecting its performance.
Carbon is one of the possible material forovercoming these issues. Because carbon
materials are generally considered as a better
catalytic material in the field of Fuel cells,
supercapactors and so on.

Fig.1 Schematic arrangement of dyesensitized solar cells
Many research groups have beensynthesizing the nanostructured carbon
materials as an efficient counter electrode [3].
However, in terms of cost and complicatedsynthesis methods it finds less attractive
towards the application of DSSC’s. To
overcome the above problems by employing
simple and easy synthesis procedures,organic sources are being attracted. In thiswork synthesis of carbon based materials
from the wastes of organics such as corn cob
and lemon peel are discussed.

2. Experimental Procedure
Fig. 2 shows the photograph of corn cob

were related to the (002) and (101) miller
indices of graphitic carbon material.

waste and lemon peel waste.
2.1. Corn cob carbon synthesis
Phosphoric acid was taken for
carbonization as well an activating agent in
one step process of chemical synthesis. In
our experiment 50 wt.% of phosphoric acidwas mixed with the carbon source of 10 g
and stirred for 3 h at 120 C. Residues ofthe filtrate were sintered at 500 C in argon
atmosphere. The final powder was washedwith distilled water several times tomaintain the pH in the range of 6 to 7.

Fig. 3 XRD spectra of corn cob carbon.
Raman spectra of corn cob carbon andlemon peel carbon are shown in Fig. 4 (a)and (b), respectively. The peaks at 1350 and
1604 cm-1 were the well known graphite Gband and defect D band which related to thegraphitic SP2 -bonding. Lemon peel carbon
as well corn cob carbon containd both peaks.
But the intensity raio of G and D bands weredifferent. In the case of corn cob carbon, IG/ID
ratio was 1.27 and for lemon peel carbon theratio was 1.

Fig. 2 Organic wastes Corn cob and lemon
peel.
2.2. Lemon peel carbon synthesis
Potassium hydroxide [KOH] was taken asa corrosive salt for carbonizing as wellactivating the wastes of lemon peel. The ratio
1:1 of lemon peel and KOH was taken with
100 ml of distilled water. The above mixer
was stirred for 3 h at 100 C. Then theresidue was subjected to sintering at 700 C
for 1 h in inert atmosphere. Further washingand filtration were carried out. A solution of 1
M Hcl is added to achieve the pH in therange of 6 to 7.

.
(a)3. Results and DiscussionCorncob
contains
cellulose,hemi-cellulose, and ligno cellulosicorganic
compounds.
Duringcarbonization the unsaturated valences ofedge carbon atoms were created becauseof the removal of other elements from theorganic sources. X-ray diffraction image ofcorn cob derived carbon is shown in Fig. 3. It
indicates that graphite based carbon wasformed. The two broad peaks at 26 and 44

(b)
Fig 4. Raman spectra of a) corn cob carbonb) lemon peel carbon

The FTIR spectra of the synthesized carbnfrom corn cob and lemon peel were shown in
Fig. 5 (a) (b), respectively. The peakcorresponded to 3400 cm-1 was related to theO-H vibrations of the hydroxyl groups. The
C=C aromatic rings vibration was located at
1564 cm-1. The bands located at 1330 and
1392 cm-1 were related to the C-O
carboxylate group. Other two peaks at 1170
and 1042 cm-1 were due to the C-O stretchingvibrations. Peaks at 667 and 578 cm-1 wereassigned to the out of plane of C-H bending
mode. The C-H out of the planecausedanother vibration at 870 and 811 cm-1.
In Fig 5 (b), the strong peak at 1550
cm-1 corresponded to the C=C carbon
material. The other peaks at 1213,1100 cm-1
and 1170, 1042 cm-1 were related to C-O
bending and stretching modes, respectively.
The peak at 872 cm-1 was due to the C-H out
of plane bending. From the Fig 5, it was
found that the presence of carbon and surface
oxygen were observed. From both the FTIR
spectra peak at 2400 cm-1 corresponded to theatmospheric CO2.

CO

respectively. Fig. 7(a) and (b) indicate theelemental analysis of corn cob carbon andlemon peel carbon, respectively.
From, Fig. 6(a) and Fig. 7(a), it was foundthat corn cob carbon based carbon contained
porous structure with different sizes andapproximately 90 at.% of the synthesizedmaterial contained carbon and 7 to 8 at.%oxygen. The remaining trace elements ofhydrogen, sulfur were present less than 0.5
at.%. Fig. 6 (b) and Fig.7 (b) indicated thatthree dimensional porous network structure
was synthesized from lemon peel carbon. The
porous arrangement was ordered throughout
the material. It contained 83 at.% of carbon ,14 to 15 at.% of oxygen and remaining oftrace elements.

(a)

(b)

2

Fig. 6 FESEM images of (a) corn cob carbon(b) lemon peel carbon

Carbon

(a)

CO

2

(b)
Fig. 5 FTIR spectra of (a) corn cob carbon
and (b) Lemon peel carbon(a)
FESEM images of corn cob and lemon
peel carbon are shown in Fig. 6(a) and (b),Carbon6,162-169 (2012).
3. Won Jae Lee, Eswaramoorthy Ramasamy,
Dong Yoon Lee and Jae Sung Song
Applied materials and interfaces 61145-1149 (2009).

(b)
Fig 7 Elemental analysis of corn cob carbon(a) lemon peel carbon (b)4. Conclusion
Synthesis of carbon based material
from the organic wastes such as corn cob
and lemon peel were performed. From
the XRD it is clear that material related tographitic based structure with disorderedgrains. Raman spectra explains that theamount of graphitic based carbon and its
disorderness. Compared to lemon peel
derived carbon, corn cob based carbon
material has a more ordered structure,
which can be understood from the ratio of
IG/ID. FTIR spectra and elemental
analysis conclude that the material has
oxygen and hydrogen, adsorbed from theatmosphere. The synthesized material ofhighly activated carbon can be utilized
further as a counter electrodes for dye
sensitized solar cells.
Acknowledgement
The authors would like to thank Mr.
T. Koyama and Mr. Tomada for theirsupport during characterization analysis.
One of the author R. Karthikeyan thanks
MEXT-Japan for the award of a research
fellowship.
References
1. O’regan, B. & Gratzel, M. Nature
353,737-740 (1991).
2. Brian E. Hardin, Henry J. Snaith andMichael D. McGehee Nature Photonics

Growth of CaSi2 layers on Si (111) substrates
Natsuki Kurebayashi

Faculty of Engineering, Shizuoka University
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
E-mail: f0711059@ipc.shizuoika.ac.jp

Abstract
CaSi2 to generate a thin film onto the Si(111) substrate in an atmosphere of Ca by vapor
deposition. Growth time becomes longer, began to increase in the proportion of Ca componentof the substrate surface. The structural properties of the resultant composites were characterized
by scanning electron microscopy (SEM), energy dispersive x-ray spectroscopy (EDS), and
X-ray diffraction (XRD).

1. Introduction
First, I will introduce silicon-based
nanosheet structure. Nanosheet is a
nanomaterials
withnanometer-scale
thickness obtained apart by a layer
constituting the layered compound, thehorizontal size of the micrometer scale.[1]There are the typical ones that are madegraphene in the current characterization
studies, and have been functionalized. Suchas high thermal conductivity and high carrier
mobility, as has been reported characteristics
are those of good quality[2]Those are expected to apply wiring material,
such as to chemical and bio-sensors. As withgraphene with silicon nanosheet effective
properties can be expected. Considered trying
to produce silicon-based nanosheet structure,and can be used as a starting material, which
is a layered crystal CaSi2 as a preliminary
step.[3]CaSi2 layered crystal structure is a Ca
atomic layer is inserted between the plane
Si(111) diamond-shaped. High vapor
pressure Ca, re-evaporation of calcium
occurs in the silicide formation temperature,
it is difficult silicide growth. Toxicity rather
than Ca, both Si, it is a material with low
environmental impact because it is anabundant element in the Earth's crust.[4]As mentioned earlier in a high vapor
pressure Ca, in the silicide formation

temperature growth is difficult because theevaporation will occur again.[5]Thus, the purpose of this study is on thegrowth of CaSi2 layers on Si(111) substrates
in an atmosphere of Ca using a quasi-closed
tube container and, the observation andcomposition analysis of the surface structureof those compounds.

2. Experimental
It was grown by vapor phase growth of
CaSi2. The composites were synthesized
fromCa
powder
(0.05g,particle
size :1-3mm) and Si substrate (0.03g). The
inside of the chamber is evacuated to about
10-3Pa degree of vacuum, heat-treated at
600 ℃ Growth time I was done in 2.5 hours
and 2.0 hours and 1.5 hours.
The structural properties of the resultant
composites were characterized by scanningelectronmicroscopy(SEM),
energydispersive x-ray spectroscopy (EDS), and
X-ray diffraction (XRD).

3. Results and Discussion
Fig.1 shows SEM images of each samplesof the surface of the substrate by the growthtime.

Fig.1 SEM images of each sample.
(a) 1.5h (b) 2.0h (c) 2.5h
Growth time of all, I look in particular at 1.5hours, as have grown big piled plate-likecrystals. The size of the plate is seem growthtime becomes longer and smaller. Crystals in
2.5 hours is quite small, it is seem to be themembrane.

Position at(%)

Ca

Si

C

1

30.67

48.80

20.53

2

28.72

53.27

13.60

3

34.28

52.95

12.77

Table.1 the results of EDS substrate surface
growth time 1.5 hours.
Table 1 shows the results of EDS substratesurface growth time 1.5 hours. The
composition ratio of Ca and Si wasapproximately 3:5 in all points.
Table 2 shows the results of EDS substratesurface growth time 2.0 hours. The
composition ratio of Ca and Si wasapproximately 4:5 in all points. It has
increased the proportion of Ca than 1.5 hours
was observed.

Position at(%)

Ca

Si

C

1

40.34

49.40

10.25

2

42.17

49.23

8.59

3

38.43

51.89

9.57

Table.2 the results of EDS substrate surface
growth time 2.0 hours.

Position at(%)

Ca

Si

C

1

63.70

28.69

7.61

2

65.28

28.58

5.88

3

64.53

27.28

8.19

Table.3 the results of EDS substrate surface
growth time 2.5 hours.
Table 3 shows the results of EDS substratesurface growth time 2.5 hours. The
composition ratio of Ca and Si wasapproximately 2:1 in all points. It has further
increased the proportion of Ca from theresults of 2.0 hours.
Based on the results of EDS, I haveobserved the XRD spectra of the sample
substrate surface obtained by heat treatmentof the composition ratio was close to 1.5
times the most CaSi2. These are shown in
Fig.2.

[4] S.Y. Jeong, J.Y. Kim, H.D. Yang, B.N.
Yoon, S.-H. Choi, H.K. Kang, C.W.
Yang, Y.H. Lee, Adv. Master, 15 (2003)
1172.
[5] J.-C. Gabriel, F. Camerel, B. Lamairem,
H. Desvaux, P. Davidson and P. Batail,Nature 413 (2001) 504

Fig.2 the XRD spectra of the substratesurface of the sample obtained by heat
treatment for 1.5 hours.
The diffraction peaks of CaSi2 wereobserved. And also, this most biggest perk issilicon. This result is believed, by performinga heat treatment for 1.5 hours in vacuo, andCaSi2 is generated on the substrate surface.

4. Conclusion
In this study we have conducted
observations of the surface structure of thecompounds obtained, the composition
analysis, and attempts film growth CaSi2
onto Si(111) substrate in an atmosphere of Cawith the container closed tube level. Growthtime becomes longer, began to increase in the
proportion of Ca component of the substratesurface. The results of the EDS, the substratesurface of the sample generated by thegrowth time 1.5 hours, it was possible formaking the composition ratio of Si and Ca is
the most close to 1:2. From XRD spectrum of
the sample is considered generated by thegrowth time of 1.5 hours in vacuum, to havebeen generated on the substrate surface
CaSi2.

References[1] T. Yui, Y. Mori, T. Tsuchino, T. Itoh, T.
Hattori, Y. Fukushima and K. Takagi,
Chem. Mater. 17 (2005) 206.
[2] M. Muramatsu, K. Akatsuka, Y. Ebina, K.
Wang, T. Sasaki, T. Ishida, K. Miyake
and M. Haga, Langmuir 21 (2005) 6590.
[3] A.S. Heintz, M.J. Fink, B.S. Mitchell,
Adv. Master, 19 (2007) 3984.

Study of the substrate treatment effect on initial growth of Ti-Oxide Nanorods by
Thermal Oxidation
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 pretreatment dependence of nanorods was investigated. The lengthand the diameter of nanorods grown on the substrate chemically-etched in hydrochloric acid by
the thermal oxidation at 800 ºC for 4 hours were approximately 1300 nm and 190 nm,
respectively. Layer and nanorods were composed of Ti-oxide such as TiO2 rutile. It is clarified
that the length and diameter of nanorods depend on pretreatment.

1. Introduction
TiO2 is an n-type semiconductor withwide band gap energy of 3.0-3.2 eV and has
been widely used for various applications in
dye-sensitized solar cells [1], gas sensors [2],photocatalyst [3] and negative electrode forlithium rechargeable batteries [4].
Nanostructured materials have attracted
much attention for its promising potential forelectronics, optoelectronics applications
because they have many merits such as ahigh surface to volume ratio and short andcontrolled electron pathway [5]. They havebeen reported in various forms such as
nanotubes [6], nanowires [7], nanoparticles[8] and nanoplates [9].
TiO2 nanostructures are synthesized bymany techniques such as sol-gel method [10],hydrothermal process [11], electrochemical
oxidation method [6] and thermal oxidation
technique [12]. Among of them, thermaloxidation technique is simple and cheap.
TiO2 nanostructures grown by this technique
was carried out under carrier gas such asacetone [7] and ethanol [12]. However,almost none of this technique is able to grow
the TiO2 nanostructures in the air.
In this study, we reported the growth of
Ti-oxide nanorods by thermal oxidation
technique of titanium substrates in the static
air. In addition, the pretreatment dependenceof nanorods was investigated.

2. Experiments
The TiO2 nanorods were grown by
thermal oxidation technique of titanium
substrates in the air. Titanium wires withdiameter of 1 mm were used as thesubstrates. The substrates were polished bythree
pretreatment
methods.
The
pretreatment methods were mechanicalpolish with #400 and #1500 sandpapers,
chemical etching in hydrochloric acid at 80ºC for 1 min and chemical etching insulphuric acid at 80 ºC for 1 min. And then,the substrates were ultrasonic cleaned in
acetone and ethanol for 10 min. The
substrates were placed at the center ofelectric furnace. The temperatures wereraised to 800 ºC for 1 hour. The substrates
were annealed in the air for 4 hour. The
temperatures were reduced to room
temperature in atmosphere.
The
Ti-oxidenanorods
werecharacterized by field emission scanningelectron microscopy (FE-SEM, JSM-7001F)
and X-ray diffraction methods (XRD, RINT
Ultima III).

3. Results and Discussion
Figure 1 shows SEM images of thesubstrate polished by three pretreatment
methods. Figure 1(a), 1(b) and 1(c) show

Fig. 1. SEM image of the substrate (a)polished with #400 and #1500 sandpaper, (b)chemically etched by hydrochloric acid, (c)chemically etched by sulphuric acid.

SEM images of the substrate polished with#400 and #1500 sandpapers, chemicallyetched in hydrochloric acid and chemicallyetched in sulphuric acid, respectively. In Fig.
1(a), 1(b) and 1(c), the surface of substrate
were different. In Fig. 1(a), the surface of thesubstrate polished with #400 and #1500
sandpapers is rough. In Fig. 1(b), on thesubstrate chemically etched in hydrochloric
acid, many ledges were formed. The
diameter of ledges was approximately 200
nm. In Fig. 1(c), on the substrate chemicallyetched in sulphuric acid, many voids wereformed. The diameter of voids was
submicron order.

Fig. 2. SEM image of Ti-oxide layer andnanorods on the substrate (a) polished with #400
and #1500 sandpaper, (b) chemically etched byhydrochloric acid and (c) chemically etched bysulphuric acid.

Figure 2 shows SEM images of TiO2
nanorods grown at 800 ºC for 4 hours. In Fig.
2(a) and 2(c), nanorods were formed and laid
randomly on the layer. In Fig. 2(a), thelength and the diameter of nanorods wereapproximately 910 nm and 130 nm,
respectively. In Fig. 2(c), the length and thediameter of nanorods were approximately
1400 nm and 190 nm, respectively. In Fig.
2(b), aligned nanorods were formed and laidon the layer. In Fig. 2(b), the length and thediameter of nanorods were approximately
1300 nm and 190 nm, respectively. The
length and the diameter of nanorods in Fig.
2(b) was formed as same as that in Fig. 2(c).

Fig. 3. XRD spectra of Ti-oxide layer andnanorods on the substrate (a) polished with #400
and #1500 sandpaper, (b) chemically etched byhydrochloric acid and (c) chemically etched bysulphuric acid. R represents TiO2 rutile structure.

However, the density of nanorods in Fig.
2(b) was more than that in Fig. 2(c).
Figure 3 shows XRD spectra of TiO2
nanostructures grown at the temperature of
800 ºC for 4 hours. TiO2 exist in three formsof tetragonal rutile, tetragonal anatase andorthorhombic brookite. In Fig. 3, somediffraction peaks of TiO2 rutile wereobserved. Overlapped diffraction peaks due
to other TiO2 structures and Ti-oxidematerials such as TiO, Ti2O3 and Ti3O5 wereobserved in Fig. 3(a). Hence, layer andnanorods in Fig. 2(a) were composed ofTi-oxide such as TiO2 rutile.

4. Conclusions
Ti-oxide nanorods were grown by
thermal oxidation technique in the air at 800ºC for 4 hour. Aligned nanorods wereformed and laid on the layer on the surface
chemically-etched by hydrochloric acid.
Layer and nanorods were composed ofTi-oxide such as TiO2, TiO, Ti2O3 and Ti3O5.

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] H. G. Yang et al., J. AM. CHEM. SOC.
131 (2003) 4078.
[10] Y. Lei et al., Appl. Phys. Lett. 78(2001) 1125.
[11] R. Yoshida et al., J. Solid State Chem.
178 (2005) 2179.
[12] S. Daothong et al., Scripta Mater. 57(2007) 567.

Structural Characterizations of Ag Nanostructures on Various Metallic Substrates
Meng Erchao
Graduate School of Science and Technology, Shizuoka University,
3-5-1, Johoku, Naka-ku, Hamamatsu, Japan
TEL&FAX : 053-478-1099
E-mail : f5145034@ipc.shizuoka.ac.jp

Abstract	  	
 Various Ag nanostructures were grown on metallic substrates using hexakisphosphate (IP6) as asilver carrier agent. The structural and morphological properties of the Ag nanostructures werecharacterized by SEM, EDS and XRD. SEM images shown that roughly three types of Ag
nanostructures were grown on different metallic substrates. EDS and XRD measurements shownthat the resultant nanostructures were silver. The growth mechanism of various Ag nanostructures
on metallic substrates was discussed.

1. Introduction
In recent years, silver has been extensively
investigated for its catalytic activities [1],anti-microbial activities [2], and used to avoid
infections and prevent spoilage. Many
researchers have focused on the anti-bacterial,
abilitytokillmicroorganisms
andmulti-functional
properties
ofsilvernano-particles [3–5]. Various methods havebeen employed for nano-silver production,and the syntheses of various shapes of Ag
nanostructures have been reported, such as
dendrite [6,7], wire [8], irregular-shaped wire[9], chain-like [10] and particle [11]structures.
For practical uses, it is important tofabricate the nanostructures using a simple
process with low costs. The systematicsynthesis study is also important to clarify thegrowth mechanism for various shaped Ag
nanostructures. In addition, the structural andmorphologicalcontrollingof
thenanostructures is also required to enhance thefunctional properties for specific applications.
In this study, a variety of Ag nanostructures
were synthesized on various metallic
substrates, and hexakisphosphate (IP6)[12–14] was used as a silver carrier agent.
Moreover, the growth mechanism of Ag
nanostructures on various metallic substrates
was discussed.

2. Experiments

Ag nanostructures were grown on various
metallic substrates using the LS21 solution
consisting of Ag-IP6. The experimental
process is as follows, as shown in Fig 1. Adrop of the solution was dripped onto metallic
substrates, then kept in air for several minutesat room temperature. Then the solution was
immediately dried using a gas burner or hot
plate. When the sample was dried, Ag
nanostructures appeared on the surface of
metallic substrate.
The structural and morphological propertiesof the Ag nanostructures were characterized
by scanning electron microscopy (SEM) alongwith energy dispersive X-ray spectroscopy(EDS), X-Ray Diffraction (XRD).

Fig 1 the experimental process ofnanostructures grown on metallic substrates

Ag

In addition, the surface condition of the Ag
nanostructure and dried IP6 complexes wascharacterizedby
X-ray
photoelectron
spectroscopy (XPS) using a VG, ESCA-LAB
Mk II with a non-monochromatized Al Kαsource (hν = 1486.6 eV). The energy
calibration for a charge correction in the
spectra was made using the C1s peak.

3. Results and Discussion
Figure 2 shows the SEM images of avariety of nano-, and micro-structures grownon various kinds of metallic substrates. As
shown in Fig 2, dense Ag nanostructures wereformed on (a) Cu, (b) Mn and (c) CaSi2
substrate. However, low density distributionof Ag structures are observed on (d) Fe, (e) Ti,(f) Al, (g) SrSi2, (h) Mo and (i) W substrates.

Figure 2. A variety of Ag nanostructures grown on
various kinds of metallic substrates.

As shown in Figure 2, roughly speaking,three types of Ag nanostructures are grown.
One is the dendrite structure, as shown infigures (a) – (c). It is also observed that thethree-dimensional Ag dendrite structures aregrown on (a) Cu and (b) Mn substrates.
However, the two-dimensional Ag dendrite
structures are grown on (c) CaSi2 substrate.
For cases (d) – (f), the fractal-shaped
structures are dominantly observed. In
addition, the particles are formed on the (g)SrSi2, (h) Mo and (i) W substrates.
A variety of Ag nanostructures were simply
fabricated using Ag-IP6 on the various typesof metallic substrates. It is interesting that thegrowth morphology of the nanostructures
depends on the metallic substrate materials.
The discussion on the substrate metal
dependence of the morphology and the growthmechanism will be discussed later.

Fig 3 the XRD patterns of Ag nanostructures on
Mn substrate

Fig 3 is the XRD patterns of Ag
nanostructures on Mn substrate. And the XRD
pattern of the Mn substrate is shown in Fig 3
for comparison. The figure shows that thepeak of Ag is detected, which indicates thatthe nanostructures grown on Mn substrate isAg.
So far, many mechanisms have been
discussed on the growth of Ag nanostructures.
In this study, Ag is from Ag-IP6. When the
LS21 solution is dripped on the metallic
substrate, the metallic atoms react with
Ag-IP6, and Ag ions are released into solution,
which provides Ag ions to form the Ag
nanostructures.
Some references reported that the transitionof the dendrite to the fractal structures iscaused by the density of the deposited atoms
in the solution. The existence of the dense
deposited atoms and anisotropic growth leads
to the formation of the dendritic structures,
and a low density of deposited atoms leads tothe formation of the fractal structures [15 –18]. Therefore, the deposition rate and silver
ions concentration are very important to thestructural control of Ag structures.

Fig 4 SEM images of the Ag nanostructure grownon (a) Al substrate and (b) Cu substrateFig 5 High magnification SEM images of the Ag
nanostructure grown on (a) Al substrate and (b)Cu substrate

Fig 4 and Fig 5 are the low-magnificationand high-magnification SEM images of theAg nanostructures grown on Al substrate andCu substrate, respectively. These two figures
shown a low density of fractal Ag
nanostructures are obtained on Al substrate.
For the Cu substrate, the dense dendrite Ag
nanostructures are grown on the surface.
Based on the previous discussion, thesubstrate metal dependence of the grown Ag
structures is caused by the reaction between
Ag-IP6 and the substrate metals. For the Cu
substrate, Ag-IP6 significantly reacted with
the Cu atoms of the substrate, and the Ag
atoms are supplied from the Ag-IP6 solution
to form the dense Ag nanostructures. Whilefor the Al substrate, the reaction of Ag-IP6
with the Al substrate is not very active. Thus,
the low density of Ag-IP6 reacted with the Al
substrate to form the low density Ag
nanostructures.
In order to demonstrate this mechanism, the
XPS spectra are carried out. Figures 6 (a) and(b) show the Ag(3d) and Cu(2p) XPS spectraof the dried metal-IP6 complexes after thereaction with the Cu substrate, respectively.
Figures 6 (c) and (d) show the Ag(3d) andAl(2p) XPS spectra of the dried metal-IP6
complexes after the reaction with the Al
substrates. As references, the Ag(3d) XPS
spectrum of the originally dried IP6, andCu(2p) and Al(2p) XPS spectra of the Cu andAl substrates, respectively, are also shown.
As shown in Fig. 6 (a), the intensity of the Ag
XPS spectra of the treated dried metal-IP6 issignificantly reduced compared to that of theuntreated Ag-IP6, which shows that the Ag
atoms are removed from the Ag-IP6 by thetreatment. On the other hand, the shape of theCu(2p) spectrum is significantly changed
compared with to that of the untreated Cu
substrate, which means that the Cu atoms

have combined with IP6 to form the metal-IP6
complexes, as shown in Fig. 6 (b). The
intensity of the Ag XPS spectra of the treated
and dried metal-IP6, as shown in Fig.6 (c), isslightly reduced compared with that of theuntreated Ag-IP6, which shows that some of
the Ag atoms have been removed from the
Ag-IP6 by the treatment. On the other hand,
the intensity of the Al(2p) spectrum is quite
low compared with that of the untreated Al
substrate, as shown in Fig. 6 (d), which meansthat Al atoms are only slightly replaced by Ag
atoms to combine with IP6 in the solution.
The XPS results agree with the previous
discussion.

Fig 6 . (a) Ag(3d) and (b) Cu(2p) XPS spectra ofdried metal-IP6 complexes after the reaction with
the Cu substrate. (c) Ag(3d) and (d) Al(2p) XPS
spectra of dried metal-IP6 complexes after thereaction with the Al substrates. As references,Ag(3d) XPS spectrum of originally dried IP6, andCu(2p) and Al(2p) XPS spectra of the Cu and Al
substrates are also shown.

Therefore, the metallic substrates
dependence of Ag nanostructrues issummarized. When Cu, Mn, CaSi2 react with
Ag-IP6, higher deposition rate and silver ion
concentration are provided, which introduces
Ag dendrite structures growth. Fe, Ti, Al
cause the low deposition of Ag, thus thefractal shaped Ag structures are formed. For
W, Mo and SrSi2, the reaction of IP6 withthese substrates rarely take place, thus ,particles are synthesized.

4. Conclusions
It was demonstrated that various kinds of
Ag nanostructures were prepared using IP6 onvarious metallic substrates. The dendritestructures on Cu, Mn and CaSi2, the fractal
structures on the Fe, Ti and Al substrates,
the particles on Mo, SrSi2 and W were
controllably obtained. It is noted that the IP6
plays an important role as a silver carrier
agent for the structural and morphologicalcontrol of the Ag nanostructures. It is
expected that this simple and easy
nanostructure fabrication technique using
metal complexes will encourage us to further
develop sophisticated materials.

Acknowledgments
I would like to thank K. Nakane, T.
Matsushita, W. Li, S. Oda and H. Tatsuoka
professor for their fruitful discussions

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Felicissimo, M. Scarpellini, D. R. Cardoso, R.
C. J. Vinhas, R. Landrs, J. F. Schneider, B. R.
McBarvey, M. L. Andersen and L. H.
Skibsred, J. Inorg. Biochem., 99, 1973-1982(2005)[13] Z. He, W. Honeycutt, T. Zhang and P. M.
Bertsch, J. Environ. Qual., 35, 1319-1328(2006)[14] Z. He, C. W. Honeycutt, T. Zhang and P.
M. Bertsch, Soil Sci. Soc. Am., 35, 1319-1328(2006)[15] M. Uwaha and Y. Saito, J. Cryst, Growth,99, 175-178 (1990)[16] H. Brune, C. Romainczyk, H. Röder and
K. Kern, Nature, 369, 469-471 (1994)[17] V. V. Agrawal, G. U. Kulkarni and C. N.
R. Rao, J. Coll. Int. Sci., 318, 501-506 (2008)[18] Y. Sawada, A. Dougherty and J. P.
Gollub, Phys. Rev. Lett., 56, 1260-1263(1986)

Characterizations of Sodium Carbonate and Stannate Nano-/Micro-rods
Kaito Nakane
Graduate School of Engineering , Shizuoka University
TEL&FAX : 053-478-1099
E-mail : f0230119@ipc.shizuoka.ac.jp

Abstract
Carbonate, stannate and silicate materials with various structures were fabricated by asimple experimental process. The molar ratio of Na : Sn affects the composition andstructural property of the sodium compound. Moreover, the structural property of
sodium silicate was affected by the growth temperature. The structural properties of
the resultant composites were characterized by scanning electron microscopy (SEM), and
X-ray diffraction (XRD).

1. Introduction
Recently, semiconducting carbonates,
silicates, germinates and stannates haveattracted much attention, as new wide gap
semiconductors. The band gap of
Sr-carbonate is estimated as 3.17 eV [1].
Zinc stannate is a wide-gap semiconductor
with a band gap of around 3.6 eV [2]. It has
been also reported that barium stannate has aband gap of 3.4 eV [3]. One-dimensional
materials have attracted much interest due totheir enhanced or modified optical,
electronic and mechanical properties
compared to those of bulk materials.
This paper shows a preliminary growthof carbonate, stannate and silicate
nano-/micro-rods. As a metallic element,
sodium is chosen here due to its low melting
point and high vapor pressure to make thereaction easy. The morphological andstructural properties of the sodium
compoundnano-/micro-rods
were
investigated.

2. Experimental
Nano-/micro- rods were prepared withstarting materials of Sn powder (99.999 %,-100 mesh), Si grain (99.999 %, 2 ~ 3 mm)
and Na grain (99 %), which were weighedin an Argon gas-filled glove box and charged
in boron nitride (BN) crucibles (DENKA;
inside diameter, 6 mm; depth, 17 mm). The
crucible was sealed in a stainless-steel tube.
The sealed tube was heated in an electric

Table 1 Preparation conditions

Na - Sn
Composite

Molar ratio
Na : SnTime Tempreture(hour)(ºC)#1

1.1 : 1.0

24

600

#2

1.0 : 2.0

24

600

Na – Si
Composite

Molar ratio
Na : SiTime Tempreture(hour)(ºC)#3

1.2 : 1.0

24

700

#4

1.2 : 1.0

24

700

#5

1.2 : 1.0

4

750

#6

1.2 : 1.0

4

650

furnace.
The details of experimental conditions
are described in Table 1.
The structural properties of the resultant
composites were characterized by scanningelectronmicroscopy(SEM), Energy
Dispersive X-ray Spectroscopy (EDS) and
X-ray diffraction (XRD).

3. Results and Discussion
Figure 1 shows the SEM images ofcomposite #1. From figures 1(a) and (b), a
large quantity of wire-like nanostructures
0.2-1µm in diameter and 6.7-11µm in length
can be observed. Moreover, it is found thatseveral wires have needle-like shape.
In addition, Figure 2 shows results of

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

Position(at.%)
(a)

Na

Sn

O

24.59

15.90

59.44

Fig.4 XRD spectra of (a) the composite #1 and(b) the composite #2.

Fig.2 SEM image and EDS analysis ofcomposite #1 shown in Table1.

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

SEM image and EDS analysis of composite#1. This results shows that the compositionof the wires in the images of Na : Sn : O≈
2:1:6. Therefore, this result suggests that theneedle-likenano-rods
consistofNa2(Sn(OH)6).
Figure 3(a) and (b) show SEM images ofcomposite #2. As shown in Fig.2 (a), a large
number of rods were formed on thecomposite surface. The inukshuk-likenanostructures are formed, as shown in theFig.2 (b). The average diameter of the rods is
1.8 µm and their lengths are in the range ofabout 7-10µm.
Figure 4 shows the XRD spectra of (a)composite #1 and (b) composite #2 shown in
Table1. From Fig.1 (a) and (b), it is foundthat sodium stannate and sodium carbonate
were synthesized respectively. In addition, it

Fig.5 SEM images of (a) the Composite #5, (b)the Composite #6, (c) the Composite #7 and(d) the Composite #5.

is considered that the needle-like nano-wires
synthesized with molar ratio of Na:Sn=1.1 :1.0 consist of Na2(Sn(OH)6) and Sn. The
inukshuk-like micro-rods synthesized withmolar ratio of Na : Sn=1.0 : 2.0 consist of
Na2CO3-H2O and Sn.
Figure 5 shows SEM images of various
sodium silicate structures formed at different
experimental
conditions.
Coral-likemicro-rods and nano-wires of sodium silicate
are formed at 700 ºC for 24 h, as shown in
the Fig.5 (a) and (b). Moreover, theneedle-like and blade-like sodium silicate
nanostructures are formed at 650 ºC and 750ºC for 4h, respectively, as shown in the Fig.5(c) and (d).

4. Conclusion
Carbonate, stannate and silicate materials
with various nano-structures were fabricatedby simple experimental methods.
The molar ratio of Na : Sn affect thecomposition and structural property of thesodium compound. Moreover, the structuralproperty of sodium silicate was affected by
the growth temperature.
In future, new semiconducting carbonate,silicate and stannate nano-/micro-rods will
be grown and their electrical and opticalcharacterizations will be expected.

Acknowledgments
I would like to thank S. Oda, T.
Matsushita, D. Ishikawa, W. Li and H.
Tatsuoka for their fruitful discussions.

Reference[1] S. Ni, X. Yang, T. Li, Mater. Lett. 65,
766–768 (2011).
[2] S. Baruah and J. Dutta, Sci. Technol.
Adv. Mater. 12, 013004 (2011).
[3] G. Larramona, C. Gutierrez, I. Pereira,
M.R. Nunes, F.M.A. Da Costa, J. Chem.
Soc.Faraday Trans. 85, 907-916 (1989).

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 wasinvestigated. 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, and theelectronic structures of Ca2Si have been
calculated with the expectation of a directtransition [5, 6].
It was also reported that the orthorhombic
Ca2Si is a stable phase with a direct 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. Warashina et al. synthesized
Ca-silicide by MA. The resulting powders
were contained Ca2Si and/or Ca5Si3 phaseswith additional cubic phase crystals [10]. It
is noted that the cubic phase of Ca2Si weresynthesized in the condition of the molar
ratio of Ca:Si=7:3.
In this study, it was attempted thatCa-silicides were synthesized different
growth condition by MA. The molar ratio of
Ca:Si=7:3 was fixed. In addition, resultant
powders were characterized.

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 70 hours in a SPEX 8000
vibratory mill.
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 patternsof the powders mechanically alloyed of Ca
and Si powders under milling conditions of(a), (b) and (c), respectively defined in Fig. 1.
The diffraction peaks of Ca-silicides and Si
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.

Fig. 1 XRD spectra of resultant powder
synthesized using the Ca granule and fine Si
powder with (a) four 1/4 inch steel millingballs for 70h, (b) sixteen 1/8 inch steel millingballs for 70h, (c)sixteen 5/32 inch steel millingballs for 70h. Ca and Si powders were
weighed in the molar ratio of Ca:Si = 7:3 inthese cases.

Fig. 2 SEM micrograph and the corresponding
EDS analysis for the powders synthesized from
Ca and Si powders with (a) four 1/4 inch steel
milling balls for 70h, (b) sixteen 1/8 inch steel
milling balls for 70h, (c)sixteen 5/32 inch steel
milling balls for 70h. Ca and Si powders were
weighed in the molar ratio of Ca:Si = 7:3 inthese cases.

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.

Reference[1] Y. Makita, in: Mc Connel (Ed.), The
First NREL Conference on Future
Generation Photovoltaic Technologies,
AIP, New York, 1997, p. 3.
[2] Y. Maeda, K.P. Homewood, T. Sadoh, Y.
Terai, K. Yamaguchi and K. Akiyama(Eds.), Thin Solid Films, 515 (2007) 22.
[3] M. Akasaka. T. Iida, K. Nishio, Y.
Takanashi, Thin Solid Films 515 (2007)
8237.
[4] J. Tani, H. Kido, J. Alloys Compd. 466(2008) 335.
[5] Y. Imai, A. Watanabe, M. Mukaida, J.
Alloys Compd. 358 (2003) 257.
[6] D. B. Migas, L. Miglio, V. L.
Shaposhnikov, V. E. Borisenko, Phys.
Rev. B 67 (2003) 2052031.
[7] C. Suryanarayana, Mechanical Alloying
and Milling, (Marcel Dekker, New York,
1997).
[8] Y. Imamura, H. Muta, K. Kurosaki, S.
Yamanaka, in: C. Uher, Y. Grin,
Ballreich, Sugiwara, G. Pastorino, M.
Cauchy, M. Udagawa, (Eds.), 25th
International
Conferenceon
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.

Investigation on dissolution process of Si into Ge meltby X-ray Penetration Method
M.Omprakash1,2, M.Arivanandhan1, R. Arun Kumar 1, H.Morii1, T.Aoki1, T.Koyama1,
Y.Momose1, H.Ikeda1, H.Tatsuoka3,Y.Okano4, T.Ozawa5, Y.Inatomi6, S.Moorthy Babu7 andY.Hayakawa1
1
Research Institute of Electronics, Shizuoka University, Hamamatsu-432 8011, Japan
2
Graduate School of Science and Technology, Shizuoka University, Hamamatsu-432 8011,
Japan
3
Faculty of Engineering, Shizuoka University, Hamamatsu-432 8011, Japan
4
Osaka University, Osaka, Japan
5
Shizuoka Institute of Science and Technology, Fukuroi, Shizuoka 437-8555, Japan
6
Japan Aerospace Exploration Agency, 3-1-1 Yoshinodai, Kanagawa 229-8510, Japan
7
Crystal Growth Centre, Anna University, Chennai, India
E-mail-royhaya@ipc.shizuoka.ac.jp

Abstract
Si dissolution into Ge melt and crystal growth of SiGe was in-situ observed by
X-ray penetration method. The rectangular shaped sandwich sample of Si (seed)/Ge/Si (feed) was used for the experiment. The penetrated X-ray intensities through thesample were recorded by rectangular shaped CdTe detector as a function of time andtemperature. The experimental results demonstrate that the dissolution of Si seed waslarger compared to Si feed crystal. Moreover, the crystal growth of SiGe was clearlyobserved from the drastic variation of penetrated X-ray intensity near the growthinterface. The growth mechanism for the observed SiGe growth process was
discussed based on the penetrated X-ray intensity profile and composition profile
measured by EPMA.
1. Introduction
Si1-xGex is a promising material forthermoelectric applications especially atelevated temperatures. However, it ishighly challenging to optimize the figureof merit by a suitable compromise
between the thermal conductivity, k,Seebeck coefficient, α, electricalconductivity, σ of the material. Moreover,
the thermoelectric properties of theSi1-xGex alloy are strongly dependent on
the composition of the material [1].
Therefore to enhance the thermoelectricperformance of the material, it isnecessary to grow the alloy crystal withhomogeneous composition. However,bulk growth of Si1-xGex withhomogeneouscomposition
fromsolution is very difficult because of thelarge miscibility gap. Any smallchanges in the solidification rate lead tosignificant composition variations [2].
Online monitoring of the solution,
composition profile and solid-liquidinterface is highly useful to understandthe solute transport and growthmechanism from the high temperature
solution.
The enormous efforts have been
made by the researcher using different
methods. Nakajima et.al [3] and his
group [4] developed an automatedfeedback control (AFC) system to keep
the temperature at the interface constantduring growth. Sazaki et al. observed
the solution temperature distribution
during the solution growth of SiGe bulk
crystal using an infrared camera[5].They investigated the position of thegrowth interface and measured thetemperaturearoundthegrowthinterface.
In the present study, we haveinvestigated the dissolution of Si into
Ge melt, and crystal growth of SiGe byusing X-ray penetration method. The
composition of the grown crystal wasmeasured by EPMA, and discussed thedissolution process.
2. Experimental Method
Fig. 1 shows the sample configurationand the temperature profile of thefurnace. Rectangular shaped Si (6×6×3mm3) (seed) /Ge (6×3×3 mm3)/Si(6×6×3 mm3) (feed) sandwich sample
was used for the experiment. The
samples were mirror-polished usingalumina abrasive powder and etched inan acid mixture of HF: HNO3 (1:1) (forSi) and HF:H2O2 (1:1) (for Ge) toremove the oxide layer. The sandwich
sample filled BN crucible was inserted
into the quartz ampoule and evacuatedabout 10-4 Pa before sealing. The sealedampoule was fixed vertically inside afurnace whose temperature profile wascarefully
measured
beforetheexperiment. The vertical temperature
gradient at the sample position was
fixed at 0.6 °C/mm. The reference
temperature of the furnace wasmonitoredusing
theR-typethermocouple connected at middle ofthe furnace.
Φ12

Φ10

QuartzCarbon sheet
Feed interface

765.5

6 mm

Si

6
mm
3
mm
6
mm

3 mm

Temperature ( °C )Si
15 mm Ge

765.0

0.6 °C / mm

764.5

764.0

763.5

Seed interface

BN
5

6

7

8

9

10

11

12

slowly decreased from 1200 °C to1170 °C at a cooling rate of-0.42 °C/min to observe the growthfrom seed-solution interface.
The penetrated X-ray intensities
were continuously recorded up to
800 °C during post-growth cooling.
Finally, the sample was cooled down toroom temperature (RT) and thepenetrated X-ray intensity through thesolidified sample was recorded for 1 s.
The grown sample was cut parallel tothe growth direction and polished withan alumina abrasive powder. The Sicomposition profile of the grownsample was measured along the growth
direction by Electron probe micro
analyzer (EPMA).

Distance (mm)Carbon sheet

2. Results and Discussion

Fig.1 Sample configuration and the insidetemperature profile of the furnace.
SampleX-Ray source
Voltage 150 kV
Current 0.1 mA

Heater

CdTe detector

Fig. 2 X-ray penetration measurement
system configuration.

X-ray source is the tungsten target withan acceleration voltage of 150 kV and acurrent of 0.1 mA. The detector is arectangular shaped CdTe line sensor,
which has 64 cells in the x-direction and1510 cells in the y-direction (Fig.2).
The temperature of the furnace wasraised up to 1200 °C at a
pre-determined heating rate andmaintained as constant for 5 h toobserve the dissolution and crystal
growth of SiGe. During heating after
660 °C, the X-rays were allowed topenetrate through the sample and thepenetrated X-ray intensities were
continuously recorded for every one
second. Ge had completely melted when
the reference temperature reached
970 °C. The composition profile of thesolution was continuously observed
when temperature maintained constantfor 5 h. After 5 h, the temperature wasFig. 3 X-ray images of the sample as afunction of time and temperature.

The image at RT shows the initialdimensions of the solid Si and Ge
sandwich sample. In the RT image, theSi seed and feed regions looked light
blue whereas the Ge region looked dark
blue due to density variation between Si
and Ge. The image corresponded to Ge
melting at 970 °C shows the volume
change of Ge during melting. When thetemperature increased, the dark blue
region increased which shows thedissolution process of Si seed into Ge
melt. Moreover, the darkness of thesolution decreased as the time andtemperature increased due to large
amount of Si dissolution.
Fig. 4 illustrates the variations of
the penetrated X-ray intensities during
heating the sandwich sample as afunction of time and temperature. The
penetrated X-ray intensity of Ge wassmaller than that of Si due to large
density difference of Ge (5.323 g/cm3)
and Si (2.390 g/cm3).

Seed

Feed

70000

80000

60000

70000

50000

RT
970 ℃
1098 ℃
1190 ℃

40000

9 min
63 min67 min 21 sec

30000
20000

Seed

Feed

60000

Intensity (cps)Intensity (cps)

80000

RT
1200 ℃
1200 ℃
1200 ℃
1200 ℃

50000
40000

68 min128 min
184 min
255 min

30000

10000
0

20000
0

2

4

6

8

10

12

14

16

Length (mm)Fig. 4 Variations of penetrated X-ray
intensities during dissolution with respect totime and temperature.

When the reference temperaturereached 970 °C in 9 min, Ge wascompletely melted. The penetrated
X-ray intensity of molten Ge was low
compared with RT intensity profile andthe total length of the Ge decreased
from 3 to 2.7 mm due to density
variation of Ge solid and liquid.
The temperature of the furnace wasincreased up to the desired growthtemperature (1200 °C). The Si seed had
started to dissolve into Ge melt as ashoulder region in the intensity profile
at 1098 °C was observed in Fig. 4. The
shoulder region near the seed interface
and the intensity of SiGe solution region
increased as the temperature increased
due to continuous dissolution of Si seed
into Ge melt at 1190 °C. Due to thecontinuous solute transport from seedinterface to the feed interface, thepenetrated X-ray intensity near the feedinterface increased due to Si
accumulation. These changes are
possibly due to solute transport fromseed interface by solutal convection andthus the solution near the feed interface
becomes Si richer.
Once the temperature reached
1200 °C, it was kept constant for about
5 h. The intensity profile changes areshown in Fig. 5. During the dwell-time,
the Si seed dissolved continuously.
Moreover, the overall intensity of SiGe
solution increased as the timeincreased
and
a
very
steep
compositional gradient was establishedin the solution.

10000
0

2

4

6

8

10

12

14

16

Length (mm)Fig. 5 Intensity profile at constanttemperature 1200 °C with various time.

As a consequence, the solution near the
feed interface got supersaturated since
the inside temperature gradient between
seed interface and feed interface of thesample was very low (0.6 °C/mm) as
measured by dummy sample prior to theexperiment (Fig. 1). Subsequently, thespontaneous growth of SiGe was started
near the feed interface and it was clearlyobserved from the intensity profile of
128 min data. Moreover, the growth wasstarted at the reference temperature ofabout 1200 °C after 128 min, thus the
crystal was Si richer and the solution
near the interface was Ge richer sincemost of the Ge was rejected from thegrowth interface due to segregation
phenomena. The heavier Ge moved toseed interface and the dissolution ofseed was proceeded further shown in
Fig. 5.
Fig. 6 shows the variations ofpenetrated X-ray intensity profile during
the cooling process of the sample. The
profile at 1200 °C after 375 min showsthat the dissolution stopped at the seed
and feed interfaces and the growth
proceeded towards seed interface. The
intensity profiles at 1000 °C after 516
min and at 800 °C after 521 min
indicate the overall intensity of the SiGe
regionincreased.
Itshows
crystallization of the residual solution
during cooling process.

8000070000

RT1200 ° C
1197 ° C
1000 ° C
800 ° C

50000
40000
30000

375 min
465 min
516 min
521 min20000

4. Conclusion

10000
0

2

4

6

8

10

12

14

16

Length (mm)Fig. 6

Intensity profile during cooling.

1 00

Si composition (% )

80

60

40

20

0
0

2

4

6

8

10

12

14

16

Length (mm)Fig. 7 Si composition distribution by EPMA
measurement system.

Dissolution of Si into Ge melt andgrowth process of SiGe was in-situ
observed using Si(seed)/Ge/Si(feed)
sandwich sample by X-ray penetration
method. The experimental results revealthat the dissolution of Si seed was
dominant when compared to Si feed.
The growth process was clearlyobserved from the time dependent
X-ray intensity profile. Due to the lowtemperature gradient, and continuous
solute transport from seed interface tofeed interface, spontaneous growth had
occurred near the feed interface. The
composition of the grown sample wasanalyzed by EPMA and it well agreed
with standard sample calibrated dataintensity.
References

70000

Solid
LiquidCrystal

60000

[1]

Z. Tan, W. A. Jesser and F.D. Rosi,
Mat. Sci. Eng. B 33 (1995) 195.

[2]

N.V. Abrasimov, S.N. Rossolenko,
W. Thieme, A.Gerhardt, W.
Schroder, J. Crystal Growth 174(1997) 182.

50000

Intensity (cps)Intensity (cps)60000

during the cooling process. The
composition of the grown sample wasanalyzed by comparing the penetrated
X-ray intensity data with the standard
sample and the composition of thegrown sample was determined nearly
Si0.7Ge0.3 as shown in Fig. 8.

40000

30000

20000
10000
0.0

0.2

0.4

0.6

0.8

1.0

Si composition (%)

Fig. 8. Grown sample intensity comparedwith standard sample intensity.

After the growth process, the grownsample was removed from the ampouleand polished to study the composition
distribution. Fig. 7 shows the Sicompositional distribution of the grownsample by EPMA measurement. Fromthe compositional profile it is found thatthe crystal was grown from the feedinterface to the seed interface with the
Si composition of 0.7. Near the seedinterface,thecompositionwasfluctuated and was Ge richer due to the
crystallization of residual solution[3] K. Nakajima, S Kodama, S.
Miyashita, G. Sazaki, S. Hiyamizu,
J. Cryst. Growth205 (1999) 270.
[4]

Y. Azuma, N. Usami, T. Ujihara, G.
Sazaki, Y. Murakami, S. Miyashita,
K.Fujiwara, K. Nakajima, J. Cryst.
Growth 224 (2001) 204.

[5]

N. Armour and S. Dost, J. Crystal
Growth 311(2009)780.

PbSrS based quantum well for optically pumped surface emitting laser
Yutaro Sugiyama
Department of Electrical and Electronic Engineering, Shizuoka University,
Johoku 3-5-1, Naka-ku Hamamatsu 432-8561, Japan

Abstract
This paper introduces theoretical background and design of quantum well for optically
pumped mid-infrared surface emitting laser. To obtain high optical gain we use multiple
quantum well (MQW) structure. We calculated the theoretical value of quantum well’s gain.

1. Introduction
Tunable mid-infrared lasers are usefulfor spectroscopic applications such as trace
gas analysis, high speed gas detection. A
high non-radiative Auger-recombination
probability and a large free carrier
absorption which is proportional to thewavelength squared are making difficult toprepare mid-infrared semiconductor lasers
operating at room temperature. Lead salt
IV-VI materials such as PbTe, PbSe, andPbS have direct gaps at L-points ofBrillouine zone and have almost symmetricband structure between conduction andvalencebands.
Thus
the
Auger
recombination probability is much smaller
than that of III-V and II-VI materials, andrelatively higher temperature laser operation
have been obtained in the IV-VI materials[1,2]. Optically pumped lasers are useful todecrease the optical losses by free carrier
absorption [3]. We recently prepared
SrS/PbS multiple quantum well (MQW)VECSELs with optically excitation (OE)
layers in both sides of the active layers [4,5].
The MQW laser operated in high efficiency
and high power at -70˚C with external
quantum efficiency of 16% and maximum
pulsed output power of 2W. We describe
here a theory and design of quantum well for
optically pumped PbS-based active layer.

E-k relationship at the bottom of conduction
band or the top of valence 2 k 2 band is given by
E=
2m
,
(2)
where  is plank constant and m iseffective mass. Therefore, density of states
for bulk material is given by

D3D ( E ) =2m 3 Eπ 2 3

.

(3)
Also

m
D 2D ( E ) = 2 ∑ Θ(E − En ) two-di
mensi
π

onal density of states for quantum well isgiven by,(4)

where En is the band edge energy of the n-th
subband, and Θ(E-En)=1 for E-En≥0 andΘ(E-En)=0 for E-En<0.
IV-VI materials has the property that

2. Density of States(DOS)Density of states for bulk material isgiven by⎛ 4 3 ⎞πk ⎟d ⎜ 3⎜⎟ .
D 3D ( E ) = 2
dE ⎜ (2π )3 ⎟
⎜
⎟
⎝
⎠

Fig.1 Schematic band diagram of a superlattice.

(1)

shapes of constant energy surfaces is
ellipsoids of revolution along the <111> axes.
Therefore, electron and hole have two kindsof effective masses; ml and along thesymmetry axes and mt transverse to thesymmetry axes. So the E-k relationshipbecomes as follows.
]2⎛E ⎞  2 ⎛ k 2 k y k 2 ⎞E⎜1 + ⎟ = ⎜ x + + z ⎟⎜ E ⎟ 2 ⎜ m m m ⎟
g ⎠
t
l ⎠
⎝ t
⎝

, (5)

where Eg is the band gap and the term
(1+E/Eg) represents nonparabolicity near
the band edge, based on two-band model.
Thus, three dimensional density of states isgiven by2mt2 ml E

D3D ( E ) =π 2 3E ⎛⎜ 2 E ⎞⎟1+
1+
E g ⎜⎝E g ⎟⎠,

mt2 ml / mzz ⎛ 2 E ⎞
⎜1 +⎟∑ Θ(E − En )⎜π 2
E g ⎟⎠⎝

(7)where mzz is the effective mass along
quantum well direction given by

(

) (

)

1 / m zz = cos 2 θ / ml + sin 2 θ / mt . (8)As usual, boundary conditions of the
wave function at the materials boundary I / IIare
Φ I = Φ II
18
[×10 ] 1

Density of states [meV-1cm-3]

In superlattice shown in Fig.1, the Blochcondition Φ(z + L) = Φ(z ) exp(ikL) must be
fulfilled, resulting in the following equation
for the subband structure of the superlatticewith layer thicknesses LI and LII:
ξ +1/ ξsin (k zI L I ) sin (k zII L II )2
+ cos(k zI L I ) cos(k zII L II ) = cos(k (L I + L II ))−, whereξ = k zI (E + E gII − U )/ k zII (E + E gI ) .

3. SrS/PbS SPSL and optical gain
We
used
SrS/PbS
short-periodsuperlattices (SPSLs) with atomic layer SrS
for the barrier layers in the MQWs. The
SPSL has higher carrier mobility comparedto PbSrS alloy owing to smaller alloy

DOS(2d) SL

500
energy[meV]600

Fig.2 Density of states for PbSrS/PbS (111)
MQW with 12nm PbSrS barrier and 8nm PbS
well. Small and large steps in the density ofstates correspond to [111] and other <111>
oblique valleys.

(9)
and

Band gap (meV)

1000

0.2
0
400

(12)

By solving the above equation, the E-krelationship of superlattice consisting ofseveral subbands and density of states are
obtained as shown in Fig.2. The density ofstates for each subband is represented by a
two-dimensional density of states.
In PbS, there is one valley along [111]superlattice direction and 3 oblique valleys
along other <111>.

DOS(bulk)
PbS
dΦ0.8
dΦ II / dz
I / dz=E +0.6E gI
E + E gII − U
0.4

(11)

(6)

and two-dimensional density of states isgiven by
D 2D ( E ) =

(10)

25 20

900

15

at 300K

SrS content (%)
10
PbSrS alloy

800
700

SrS/PbS SPSL

600
500
400
1

2

3
4
PbS well width (nm)5

6Fig.3 Band gaps of SrS/PbS SPSL withmonolayer SrS and PbSrS alloy with the same
SrS content.

Transmittance (%)

100

1

80

0.8

QW

60

γ = α { f c (1 − f v ) − f v (1 − f v )} = α (PbSrS
f c − f v ) 0.6
SPSL
buffer40
20

PbSrS(8nm)/PbS(12nm)
40QWs on PbSrS(1.05 µm)
0
1000
2000
3000
4000
Wavenumber (cm-1)

0.4
0.2

5000

0
6000

Absorption coefficient (µm-1)

scattering. And the higher mobility is useful
to decrease free carrier absorption because
free carrier absorption is inversely
proportional to the mobility [4,5]. Figure3
shows the band gap of the PbSrS SPSL, andthat of PbSrS alloy with same SrS content.
The band gap of the SPSL increases with thedecrease of the PbS layer thickness.
Absorption coefficient is proportional todensity of states. The absorption coefficientis obtained simulating of the theoreticalFig.4 Optical transmission spectra of
PbSrS/PbS 40MQW, grown on PbSrS buffer
layer. BaF2(111) was used for the substrate.
PbSrS layer consists of SrS/PbS(12nm) SPSL.
Black dashed line shows theoreticaltransmission spectrum calculated from theindicated absorption coefficients (solid line in
red color), and blue line shows experimental
spectrum.

Optical gain for one QW (10-3)

10
PbSrS/PbS(12/8nm) [111] 5QWs
8
6

80K

4
2
300K
0
0

1
2
3
4
5
6
18
-3
Injected carrier concentration (10 cm )Fig.5 Dependence of optical gain on injected
carrier concentrations for PbSrS/PbS MQWfrom 80 to 300K.

transmission spectrum to the experimentalone.
The optical gain γ is obtained byfollowing equation.
(13)

Fig.5 shows the dependence of optical
gain for one PbS quantum well grown in[111] direction on excitation carrier densitywith the parameter of temperature from 80K
to 300K. Maximum gain per one quantum
well is about 0.002 at 300K with excited
carrier concentration n=p=5×1018cm-3. Thus,
if we prepare the active layer with 5QWs,
mirror reflectivity above 99% is necessary toobtain laser operation.
4. Summary
Theoretical background and designof quantum wells for optically pumped
mid-infrared surface emitting lasers were
described. To obtain enough optical gain amultiple quantum well (MQW) structure isnecessary. With the mirror reflectivity about
99%, number of quantum wells around fiveis necessary for room temperature laser
operation.
Reference[1] D. L. Partin, Superlattices and Microstructures 1, 131 (1985).
[2] B. Spanger, U. Schiessl, A. Lambrecht,
H. Böttner, and M. Tacke, Appl. Phys.
Lett. 53, 2852 (1988).
[3] A. Khiar, M. Rahim, M. Fill, F. Felder, F.
Hobrecker, and H. Zogg, Appl. Phys.
Lett. 97, 151104 (2010)[4] A. Khiar, M. Rahim, M. Fill, F. Felder,
H. Zogg, D. Cao, S. Kobayashi, T.
Yokoyama, and A. Ishida, J. Appl. Phys.
110, 023101 (2011)[5] A. Ishida, Y. Sugiyama, Y. Takano, H.
Sakata, M. Rahim, A. Khiar, M. Fill, F.
Felder, and H. Zogg, submitted to Appl.
Phys. Lett. 99, 121109 (2011)

The effect of Ca substitution for Ba on the superconductor YBa2Cu3Oy
Toru Tanikawa* and Yasuhiro Hayakawa
Research Institute of Electronics, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8011, Japan
TEL&FAX: +81-53-478-1338*
e-mail: rt2senor@hotmail.co.jp

Abstract
Characteristic changes by the substitution of Ca for Ba in YBa2Cu3Oy superconductors havebeen investigated. The pellets were characterized by scanning electron microscope (SEM) together with X-ray energy dispersive spectroscopy (EDS), X-ray diffraction (XRD) and magnetic susceptibility measurement. It was found that the sample didn’t show superconductivity
between 56 K and room temperature. CaO rich precipitates were formed.

1. Introduction
The high-temperature superconductivity
at 30 K in the oxide compound (LaBa)2CuO4
was discovered by Bednorz and Muller [1] in1986. Soon after that Cho et al. and Wu et al.
[2-3] found the superconducting transition
temperature around 90 K in Y-Ba-Cu oxidesystem. These discovery of high-Tc superconductors simulated many researchers toinvestigate new oxide superconductors withhigher Tc and extensive studies have been
made on various properties of these oxides.
Up to now, however, the highest Tc is 164 K
which is recorded on Hg2Ba2Ca2Cu3Ow athigh pressure [4].
In order to find high-Tc superconductors,
we have investigated about the effect of substitution of Ca for Ba on the YBa2Cu3Oy superconductor.

2. Experimental
At first, the standard YBa2Cu3Ox superconductor was synthesized for a reference.
Y2O3, BaCO3 and CuO of reagent-grade purity were used as source materials. It was
sintered at 900 ˚C for 12 h in air and then it
was reground and cold-pressed into
disk-shape pellets. The pellets were sinteredat 960 ˚C for 12 h in air, after that the furnace was cooled to 500 ˚C and annealed in
O2 atmosphere at 24 h, then, finally the furnace was cooled to room temperature at aslow speed.
Three kinds of samples (samples A, B,

C) were synthesized from powders of Y2O3,
CaCO3 and CuO of reagent-grade purity. For
the synthesis of all samples, the appropriate
amounts of powers were mixed. Figs. 1(a)and (b) show the pre-sintering, sintering andannealing processes of the sample A, respectively. It was sintered at 960 ˚C for 12 h inair, thoroughly reground and cold-pressed
into disk-shape pellets. The pellets were sintered at 960 ˚C for 12 h in air, after that the
furnace was cooled to 500 ˚C and annealedin O2 atmosphere at 24 h, then, finally the
furnace was cooled to room temperature at aslow speed.
Sample B was sintered under O2 atmosphere in all the above processes. The temperature and time conditions of sample B
was the same as the conditions of sample A.
For the sample C, the temperature-condition
was changed during the first sintering process. The temperature was decreased to 1000
˚C for the sintering process of samples C.
The atmosphere and time conditions of sample C was the same as the conditions ofsample A.
The surface morphological, compositional and structural properties of all thesamples were analyzed using Scanning Electron Microscopy (SEM), Energy Dispersive
Spectroscopy (EDS) and X-ray powder diffraction (XRD) apparatus. Magnetic susceptibility measurements were performed by an
AC susceptibility- temperature measurement,using a CHINO SK03-01A.

Fig. 1(a) Pre-sintering and sintering processes.

Fig. 1(b) Annealing process.

3. Results and Discussion
3-1 YBa2Cu3Ox superconductor
Fig. 2 shows XRD pattern of the standard YBa2Cu3Ox superconductor. It indicatedan orthorhombic structure and the lattice
constant was a = 3.7845, b = 3.9225 and c =11.6367 Å
The temperature dependence of the magnetic susceptibility is shown in Fig. 3. The
critical temperatures which are 10 % and
90 % of the change of susceptibility from thenormal state were 89.4 and 90.2 K, respectively.

Fig. 2 XRD pattern of the standard YBa2Cu3Ox
perconductor.

Susceptibility [a.u.]
80

90
Temperature [K]100Fig. 4 XRD pattern of the sample A.

Fig. 3 The temperature dependence of the magtic susceptibility of the standard YBa2Cu3Ox supernductor.

3-2 Substitution of Ca for Ba on the
YBa2Cu3Oy
The XRD pattern of sample A is shownin Fig. 4. The pattern was different from the
YBa2Cu3Oy (Fig. 2). It suggested that the
crystal structure of sample A was not sameas the YBa2Cu3Oy structure and Ca was not
simply substituted of the Ba site. The samples B and C also indicated the same structures.
Figure 5 shows the surface morphologyof sample A recorded by SEM. The composition of the P and Q area were measured by
EDS. The ratio of Y : Ca : Cu : O was 13.3 :16.7 : 34.3 : 35.7 at the area P and 1.2 :51.8 : 3.7 : 43.3 at the area Q. It indicatedthat CaO-rich precipitates were formed at thearea Q. To make clear the distribution of theelements, the mapping of the Y, Ca. Cu andO was measured as shown in Fig. 6. It was
clearly seen that the Ca rich area were included in the sample.

Fig. 5 Surface morphology of sample A
easured by SEM.

The temperature dependence of themagnetic susceptibility is shown in Fig. 7.
The sample A did not indicate the superconductivity between 56 K and room temperatureFig. 6 the mapping of the Y, Ca. Cu and O.

4. Conclusion

Susceptibility [a.u.]

YBa2Cu3Ox standard sample was synthesized and measured the superconducting
properties. Samples which substitute Ca for
Ba on the YBa2Cu3Oy did not show superconductivity between 56 K and room temperature. The crystal structure was different
from YBa2Cu3Oy. CaO rich precipitates wereformed.

Reference
0

100
200
Temperature [K]300

Fig. 7 The temperature dependence of theagnetic susceptibility of the sample A.

[1] J.G.Bednorz, K.A.Mueller, Z.Phys.B64(1986) 189.
[2] C.W.Chu, P.H.Hor, R.L.Meng, L.Gao,
Z.J.Haung,
Y.Q.Wang,
J.Bechold,
D.Cambel, M.K.Wu, J.Ashburn and
C.H.Haung Phy.Rev.Lett.58 (1987) 405.
[3] M.K.Wu, J.R.Ashburn, C.J.Torng,
P.H.Hor, R.L.Meng, L.Cao, Z.L.Hung,
Q.Wang and C.W.Cho, Phys. Rev. Lett.
58 (1987) 908.
[4] A.Schilling, M.Cantoni, D.J.Guo and
H.R.Ott, Nature 363 (1993) 56.