Developing controllable anisotropic wet etching to achieve silicon nanorods,
nanopencils and nanocones for efficient photon trapping
Ho-Yuen Cheung,‡ Hao Lin,† Fei Xiu,† Fengyun Wang,† TakFu Hung,† Johnny C. Ho†,*,Chun-Yuen Wong‡,*†
Department of Physics and Materials Science, City University of Hong Kong, Tat Chee
Avenue, Kowloon Tong, Kowloon, Hong Kong‡
Department of Biology and Chemistry, City University of Hong Kong, Tat Chee Avenue,
Kowloon Tong, Kowloon, Hong Kong
TEL&FAX: 3442 4897 / 3442 0538
*E
-mail: johnnyho@cityu.edu.hk

Abstract
This paper introduces a simple wet-chemistry based anisotropic etching for fabricatinghigh aspect ratio nano-arrays with different morphologies. This method has a great potential
for application of producing efficient, low-cost and large-scale photon trapping devices.

1. Introduction
Since silicon (Si) nanostructures, including
nanowires, nanopillars and nanocones,
contain unique optical and electricalproperties
andextraordinary
surface-area-to-volume ratio, they have been
widely used in device applications such as
solar cell sensors, vertical transistors, etc. inrecent years. One of the famous methods forfabricating those nanostructures with highaspect ratio is reactive ion etching (RIE). Yet,
the drawbacks include high cost, induced
damage of the crystal orientation on thesurface and chemical contamination.
On the other hand, using wet anisotropic
etching approach to produce sharp andwell-controlled features has not been wellachieved because of the difficulties inchoosing appropriate chemicals as etchants.
Most of the etchants preferentially erode theunderlying substrate in an isotropic manner.
Here,
we
introduceasimplewet-chemistry-only fabrication scheme forthe controllable hierarchy of highly regular,
single-crystalline and high aspect ratio Si
nanostructures with different geometricalmorphologies, ranging from nanorod,
nanopencil and nanocone arrays. Moreover,these nano-arrays can be manipulated withtunable dimensions and spacing over a wide
range in large areas, which is particularly
attractive for applications requiring efficientphoton trapping (e.g. photovoltaics).

2. General Experimental Details
In this work, the Si nanopilliar array
acted as fabrication templates made bynano-sphere lithography using polystyrene(PS) nanospheres followed by themetal-assisted chemical etching. Briefly, theTi/Au mesh film is deposited on the Si
surface, then is performed etching in
HF/H2O2 solution. The fabricated templates
were next treated with a mixture of AgNO3,
HF, and HNO3 or H2O2 for the selective
deposition of Ag clusters on the nanopillar
tips for the further Si etching. Figure 1
shows the SEM images of different
morphologiesofnanopillar
arraysmanipulated by a number of repeating timesof Ag cluster deposition and removal
processes.

3. Results and Discussion

Fig. 1. Representative SEM images of nano-arrays with different geometrical morphologies (ranging
from nanopillars, nanorods, nanopencils to nanocones; panel a to d with the pitch of 0.7 μm; panel e to hwith the pitch of 1.3 μm). © Royal Society of Chemistry

Figure 2 shows the proposed formation
mechanism of different morphologies ofsuch arrays. The fabrication templates werefirst etched by [AgNO3+ HF + HNO3 / H2O2]
system. The anisotropic behavior was due tothe presence of Ag clusters which acted as
catalysts formed by AgNO3. Because of thedifference of crystallinity, the Ag clusters

Fig. 2 Schematic illustrations of the formation
mechanism of Si nanostructures with different
morphologies observed in Fig. 1. © Royal
Society of Chemistrywere mainly formed at the apical position of
each
nanopillar.
Themorphologicaldifference could be controlled by theconcentration of AgNO3 (nanocone formed
from higher concentration) and changing theoxidant (nanorod formed by using H2O2
instead of HNO3)
Figure 3 shows the UV-vis reflectancemeasurement in the spectral range of 400 –800 nm of Si nanopillars with different
morphological geometries. According to
the reflective spectra, nanorods andnanopencils exhibit very low reflectancedown to <5% compared to the nanopillar

Fig. 3 Optical reflectance measurement ofdifferent fabricated morphological nanopillar
arrays. Nanopillars, nanorods and nanopencils
with the pitch of 0.7 μm and 2 μm depth. ©
Royal Society of Chemistry

and bare flat substrates, which is mainly
due to the apical region with a smaller
pillar diameter for more efficient photon
capturing, trapping and transmission
down to the base region with a higher
material filling ratio and enhanced
absorption.

4. Conclusion
Si nanorod, nanopencil and nanocone
arrays with large scale have beensuccessfully fabricated by the simple
wet-chemistry. Those fabricated nano-arrays
were shown with high performance forphoton trapping ability.

Reference[1] H. Lin, H. Cheung, F. Xiu, F. Wang, S.
Yip, N. Han, T. Hung, J. Zhou, J. C. Hoand C. Wong, J. Mater. Chem. A,
2013, 1, 9942

Iron Phthalocyanine Coated Carbon Nanohorn for Oxygen Reduction Reactionin Proton Exchange Membrane Fuel Cells (PEMFCs)
Guofa DONG, Fengyun WANG, Johnny C. HO*
Department of Physics and Materials Science, City University of Hong Kong, 83 Tat Chee
Avenue, Kowloon, Hong Kong
TEL&FAX: +852-3442 4897，+852-3442 0538*
E-mail: johnnyho@cityu.edu.hk

Abstract
A composite of iron phthalocyanine (FePc) coated carbon nanohorn (CNH) was
fabricated by a simple method and its catalysis on oxygen reduction reaction (ORR) was
studied. Comparing with FePc, FePc/CNH composite shows much better catalytic property
for ORR. The composite can give an ORR peak potential of -0.13 V, higher than the potential
on commercial Pt/C. The composite also has a good selectivity for ORR when methanol
coexists.

1. Introduction
Oxygen reduction reaction (ORR) is afundamental process for the energy transfer
in PEMFCs. The reactive efficiency of ORR
will determine the entire performance of a
PEMFCs1. Up to now, the most efficient
ORR catalysts in commercial application are
still platinum (Pt) or Pt-based materials, but
Pt is a noble metal and its reservation islimited2. So developing cheap andhigh-performance ORR catalysts is a maintask in the research of PEMFCs.
Metallic phthalocyanines (MPc, shown in
Fig. 1) are macrocycle molecules and used asdyes for many years. MPcs are also perfectcatalysts for many electrochemical reactions3.
In 1964，Jasinski first reported that CoPc heldcarbon nanomaterials such as carbon
nanotubes, carbon fibers and graphene, havebeen used for loading MPcs to improve thedispersion of MPc and accelerate the electrontransfer10, 11. The ORR catalytic performanceof the composite based on MPc wasenhanced a lot, but it still needs a further
improvement to meet the common
applications.
Carbon nanohorn (CNH) is an individual
particle composed of an aggregate of manyhorn-shaped sheaths of graphene sheets12. Its
shape looks like dahlia or bud, and it holds
very large surface area, excellent electricalconductivity and outstanding thermal andchemical stability. It is an ideal supporter13, 14.
Here we select carbon nanohorn as the
supporter and fabricate a composite of FePc
and CNH by a simple method in order toimprove the catalytic performance of FePc on
ORR and achieve a close to perfect ORR
catalyst.

2. Experimental section
2.1 Preparation of the Composite

Fig. 1 Scheme of metallic phthalocyanine

good catalytic property on ORR in alkaline
media4. From then on, many different
catalysts based on MPcs for ORR aredeveloped5, 6, but being prone to aggregation7,8
and its poor electric conductivity9 limit theelectrocatalytic performance of MPc. To
improve the performance of MPc, many

Iron
Phthalocyanine(FePc)
wascommercially available. Carbon nanohorn
was produced by a direct current arc
discharge15. The composite of FePc and CNH
was fabricated by mixing of FePc and CNH
in CHCl3 under ultrasonication for 8 hours,
then the black suspension was stirred andkept still for 10 hours. The final composite
was separated from the solution by vacuumrotary evaporation and then dried undervacuum and 80 oC overnight. The

commercial Pt/C (20 wt % Pt on Vulcan
XC-72) was obtained from Alfa Aesar.

but there are just a few weak characteristic
peaks for CNH. In the FePc/CNH composite
the characteristic peaks from FePc are very

2.2 Characterization
The final composite was denoted as
FePc/CNH and characterized by Fourier
Transform infrared (FTIR) spectroscopy, andtransmission electron microscopy (TEM,
Philips CM-20).
2.3 Electrochemical Measurements
Electrochemical measurements were
performed on an electrochemical workstation(CHI 660C). A standard three-electrodeelectrochemical system connected with a gas
flow systems was employed. A glassy carbon
disk electrode (GCE, diameter of 5 mm), an
Ag/AgCl (saturated KCl-filled) electrode anda platinum wire were used as the working
electrode, reference electrode and counterelectrode,
respectively.
The
ORR
performance of the catalysts was determinedby cyclic voltammogram (CV) measurements
in an aqueous solution of 0.1 M NaOH. CV
curves were recorded at a scan rate of 100
mV s-1.
The electrode surface was polished withthe slurry of α-Al2O3 powder with different
particle sizes (500nm, 100nm, and 50 nm),then washed with DI water, ethanol, and DI
water accordingly and separately for 5
minutes and finally dried in hot air. For
modifying the GCE, 2.0 mg FePc/SWCNT,
1.0 mL isopropanol and 50.0 μL Nafion (DuPont Corp., 5 wt %) were mixed into a ink.
30 μL of the ink was loaded on the GCE
surface. For comparison, the GC electrodes
coated with FePc, SWCNT or Pt/C (20 wt %Pt on Vulcan XC-72) were also fabricatedwith the same procedure.

3. Results and Discussion
3.1 Characterizations of FePc/CNH
Fourier Transform infrared (FTIR)
spectroscopy is a common and powerful tool
to characterize materials, and it can provide
the characteristic peaks for one certain
compound. The FTIR spectra of CNH, FePc
and FePc/CNH are shown in Fig. 2. FePc
shows rich and strong characteristic peaks

Fig. 2 Infrared spectra of CNH, FePc and
FePc/CNH

obvious, but the intensities of these peaks
become weak, which may be attributed to the
π-π electronic interaction between FePc andthe graphene layer of CNH. From the FTIR
spectra, it can be concluded that FePc wasloaded on the CNH successfully. The TEM
images of CNH and FePc/CNH are shown in
Fig. 3. The CNH is very neat and its layers
can be observed clearly in Fig. 3(a), but afterFig. 3 TEM images of CNH (a) and FePc/CNH
composite (b)mixing with FePc the structure of CNH
becomes vague due to the black material
covered on the surface. Obviously it is thelayer of FePc fixed on CNH.
3.2 Electrochemical Performance of the
FePc/CNH for ORR Catalysis
ORR activity was determined based on
steady-state polarization curves, which wererecorded after 20 pre-sweeps. Fig. 4 shows
the CV curves of FePc/CNH in NaOH
solution saturated by N2 and O2, respectively.
In N2 saturated solution, only some weak

peaks can be found and these peaks may
result from the intrinsic properties of
FePc/CNH. In O2 saturated solution, a strong
peak is achieved at a potential of -0.13 V,
which is obviously caused by the ORR.

highest of the three and it is 40 mV and 240
mV higher than FePc and CNH, respectively.
The peak current intensity on FePc/CNH
composite is also the highest of the threevalues, and it increases to -2.01 mA cm-2
which is about twice of the peak current
catalysis, comparing with FePc. This can be
attributed to the high dispersion of FePc andfast electron transfer by CNH. The catalytic
activity of the FePc/CNH composite for ORR
is also compared with the commercial Pt/C,
which is shown in Fig. 6. There is an increaseabout 40 mV in potential for FePc/CNH
comparing with Pt/C.

Fig. 4 CVs of FePc/CNH in 0.1 M NaOH solution
saturated with N2 (black) or O2 (red)The CV curves for FePc and CNH in O2
saturated NaOH solution are shown in Fig. 5
for comparing the ORR activity of different
materials. It can be found that the peakpotential on FePc/CNH composite is theFig. 6 CVs for Pt/C and FePc/CNH in 0.1 M
NaOH solution

Fig. 5 CVs of FePc, CNH and FePc/CNH in 0.1
M NaOH solution saturated with O2

The selectivity for ORR is also an
important property for an efficient ORR
catalyst. Though Pt-based catalysts have highactivity for ORR, they often suffer from acrossover effect caused by the interference ofcarbon oxide, methanol and so on, and thenlose their ORR catalytic activity completely.
Here, the ORR selectivity of FePc/CNH
composites was also assessed and shown in
Fig. 7. For FePc/CNH there is almost noFig. 7 CVs of FePc/CNH (a) and Pt/C (b) in 0.1 M NaOH solution saturated with O2
with and without 1 M CH3OH.

apparent change before and after 1 M
CH3OH was added into the NaOH solution(Fig. 7 (a)), which implies that FePc/CNH
composite has a high selectivity for ORR.
But there is an obvious change in the CV
curves of Pt/C and the ORR peak at -0.17V
disappeared completely and two new upside
peaks at -0.06 V and 0.07 V appear after the
addition of CH3OH.

4. Conclusion
In conclusion, we have prepared acomposite of FePc and CNH, and alsostudied its catalytic performance for ORR.
The FTIR spectra and TEM images provedthat FePc was fixed on CNH successfully.
The electrochemical results implied that
FePc/CNH shows good activity andselectivity for ORR and it is an excellent
catalyst for the reduction of oxygen.

6.

7.

8.
9.
10.

11.

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Monolayer molecular doping for cost-effective Si photovoltaics
Fei Xiu, Johnny Ho*
Department of Physics and Materials Science, City University of Hong Kong, 83 Tat Chee
Ave., Kowloon Tong, Kowloon, Hong Kong*
E-mail: johnnyho@cityu.edu.hk

Abstract
This paper introduces our recently developed low-cost, scalable, monolayer surface dopingof semiconductor materials for the Si photovoltaic device fabrication without using
conventional costly doping schemes such as the ion-implantation. This novel technique isbased on the formation of self-assembled monolayers of dopant-containing molecules on thesurface of Si followed by the subsequent thermal diffusion of dopant atoms via rapid thermal
annealing. To maximize the absorbed light into the device, surface of crystalline silicon wafer
is randomly texturized by the chemical anisotropic etching. In our preliminary study, energy
conversion efficiency of ~6% under the illumination of AM1.5G is achieved based on thisprocess scheme.

1. Introduction
In the past decades, diminishingsupplies in natural fuel resourcesand environmental issues from combustion
gases highlight the need for new energy
sources [1–3]. Being clean, renewable anduniversally abundant, solar energy is a viablechoice to meet the increasing energy demand.
However, it is also the least harvested formof renewable energy.[2] The major obstacles
come from the substantial fabrication andmaterials costs of existing high performancesilicon solar cells which are not economical
for terrestrial daily applications, due to thefact that solar cell modules already take up
approximately 60 % of the total cost for thesolar power installation.[3] In this regard,
tremendous effort has been made to develop
highly efficient and low-cost Si photovoltaic
devices.
Typical solar cells rely on p/n junctions
to collect the photo-generated electron-holepairs and then convert them into
electricity.[4] During fabrication processes,
impurities such as dopants are thermally
introduced to the Si crystal to produce thosejunctions; however, this is not possible tocontrol the junction profile near the surface
region to attain so-called shallow junctions(USJs) to push the ultimate energy

conversion to the thermodynamic efficiency
limit.[4] Device physics modeling has shownthat the formation of ultra-shallow junction
is preferred to achieve high PV efficiency
since the majority part of electron-hole pair
optical generation occurs just below thesurface of the devices. Thus the shallow
junction can effectively separate charge
carriers.
Recently, we have developed an
innovative approach to achieve the low-cost,large-scale, high throughput and controllable
monolayer doping (MLD) for shallow
junctions in Si, which is based on theformation of self-assembled monolayers of
dopant-containing molecules on the Si
surface followed by the subsequent thermal
diffusion of dopant atoms via rapid thermal
annealing.[5,6] Notably, in contrast to thecompeting technologies, monolayer doping
does not require expensive vacuum
equipment and long vacuum process time
which is attractive to be used as amass-production platform. By texturizing thesubstrate into randomly distributed pyramid
structure, light absorption properties of thedevice is improved which contributes to thehigher conversion efficiency. In our initialstudy, energy conversion efficiency of 6.2 %under the illumination of AM1.5G wasachieved based on this process scheme.

Figure 1 Schematic of the shallow junction formation process.

Figure 2. SEM images of the Si pyramid structure.

2. n+/p junction formation by MLD
and photovoltaic device fabrication
For the phosphorus-MLD (P-MLD)
process, p-type Si wafers were first treatedwith dilute hydrofluoric acid to remove thenative SiO2 in the glove box. The Si surface
wasthenreactedwith
diethyl
1-propylphosphonate (DPP, Alfa Aesar) andmesitylene as a solvent (25:1, v/v) for 2.5 h
at 120 °C to assemble a P-containing
monolayer [6, 7]. Following, a layer of 50
nm thick SiO2 is electron-beam evaporated
as a cap, and the substrate is spike annealed
between 900-1050 °C in Ar ambient to drivein the P atoms and achieve n+/p USJs. The
spike annealing is performed in a rapid
thermal processing tool with a fast ramping
rate of 100°C/s to the target temperature.
Schematic of shallow junction formation

process is demonstrated in Figure 1.
After the formation of n+/p junction, ITO
and Ni were deposited as the top and bottom
electrodes, respectively. Ag finger grids werethen thermal evaporated on top of the ITO
layer to help collecting charge carriers more
effectively.

3. Results and Discussion
Here the random Si pyramid structure
was achieved by the chemical anisotropic
etching using a potassium hydroxide andisopropyl alcohol solution under 80 oC. It
could be observed from Figure 2 that the Si
substrate has been well texturized withrandom pyramids with an average height of~ 8 um.
Light reflection of the planar siliconwafer and silicon pyramid structures etchedfor different etching time is compared in
Figure 3. It is found that light reflection of30 min65 min90 minplanar wafer

Reflection (%)

50
40
30
20

400

500

600

700

800

Wavelength (nm)Figure 3. Light reflection property of Si wafer and Si pyramid structures.

pyramid structures is much lower than theuntreated Si wafer, indicating a better lightabsorption property of the pyramid structure.
Notably, as the etching time increases, lightreflection of the obtained pyramid structure
decreases. And the pyramid structure under
90 min etching shows the lowest lightreflection property.
On the other hand, to decrease the backcontact side carrier recombination, a rear p+
back surface doping is achieved by ion
implantation with boron. Sheet resistance ofthe back side of the wafer is measured by thefour point probing and a low sheet resistanceof 24 ohm/sq is achieved. Solar cell devices
are then fabricated after the deposition ofcontacts, as shown in Figure 4, andcorresponding photovoltaic performance isinvestigated by solar simulation under AM
1.5G.
Figure 5 demonstrates the I-V characteristics
of pyramid structured solar cell device withp+ back side doping. This pyramid solar cellhas the Voc of 0.45 V, Jsc of 28.6 mA/cm2
and fill factor (FF) of 48.34%, with theenergy conversion efficiency of 6.2 % underAM1.5G illumination.
For the qualitative comparison, Table 1
summarized the Voc, Jsc, FF and efficiency
values of the solar cell devices with different
surface structures and the effect of p+ backAg contact grids

Figure 4. Schematic of the pyramid solar cell
device structure with the p+ back side doping.

side doping. Contrasted with the planar
device, there is a great improvement in Jsc
and conversion efficiency for the pyramidsolar cells due to the efficient light trapping
properties of the pyramid structure. A little
decrease in FF after the pyramid texturizing
is observed probably because of the increasein the series resistance. After the formationof p+ layer in the back side by the boron
doping, Jsc is improved significantly forboth planar and pyramid devices. The p+
doping formed at the back side of Si
substrate could generate an inner electricfield to repel photo-induced minority carriers
at the backside and decrease theback-surface recombination, which leads tothe improvement of Jsc.

30

Jsc /mA

20

10

0
0.0

0.2

0.4

Voc / V
Figure 5. I-V characteristics of the pyramid solar cell device with p+ back side doping.

Table 1. Performance parameter of solar cell devices with different structures.
DeviceVoc

Jsc
2

Fill Factor

Efficiency

structure(V)(mA/cm )(%)(%)

n+/p/p+ (planar)

0.44

21.49

53.97

4.99

n+/p (planar)

0.46

15.31

64.76

4.52

n+/p/p+ (pyramids)

0.45

28.61

48.34

6.2

n+/p (pyramids)

0.41

23.8

52

5.03

4. Conclusion
In this work, studies of monolayer
doping approach for the shallow p-n junction
formation and application in single
crystalline silicon solar cells were presented.
This cell fabrication process is a simple andlow-cost approach based on a phosphorus
doping step achieved by the monolayer
doping technique at the front to create then+/p junction. In our study, 6.2% efficiency
was accomplished based on the pyramidsurface texturization. To further explore thehigher photovoltaic performance, additional
studies on the fabrication of three-dimension(3D) nanostructured solar cell devices withsuperior light harvesting properties will be
carried out.

Reference[1] M. S. Dresselhaus, I. L. Thomas, Nature.
414(2001) 332–337.
[2] M. Grätzel, Nature. 414(2001) 338–344.
[3] G. Goncher, R. Solanki, Proc. SPIE.

(2008) 7047.
A.L. Fahrenbruch, R.H. Bube,
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[4][6] J.C. Ho, R. Yerushalmi, G. Smith, P.
Majhi, et al., Nano Lett. 9(2009)725-730.

Surface roughness induced electron mobility degradation in InAs Nanowire
SenPo Yip, Fengyun Wang, Ning Han, KitWa Fok, Guofa Dong, TakFu Hung, K S Chan,
Johnny C Ho*
Department of Physics and Materials Science, City University of Hong Kong, 83 Tat Chee
Avenue, Kowloon Tong, Kowloon, Hong Kong
TEL: +852-3442-4897*
E-mail: johnnyho@cityu.edu.hk

Abstract
In this work, the surface roughness dependent electron mobility in InAs nanowires isstudied. The electron mobility of InAs nanowires degrades monotonically as the nanowire
surface roughness increases after detailed electrical characterizations. Under low temperature
measurements, the surface roughness shows a dominant impact on the electron mobility after
the surface/interface traps and phonon scattering have been decoupled. All these highlight thatthe nanowire geometries and surface conditions must be carefully considered for the device
design of optimal performances.

1. Introduction
III-V semiconductor nanowires (NWs),
especially InAs NWs with the high electron
mobility, have attracted large researchinterests
aschannelmaterials
forhigh-performance transistors. The electricalproperties of these NWs are dependent on
their dimension and surface condition due toits high surface-to-volume ratio. This work
presents a semi-quantitative study of thedependency of electron mobility on thesurface roughness of InAs NWs, in which
the intrinsic electron mobility decreases with
the increase of surface roughness. The
dominant role of the surface roughness
scattering on the electron mobility
degradation is highlighted under the
low-temperatureelectricaldevicecharacterization. The experimental results
are important for the design and optimizationof device performance constructed by NWs.

respectively, under H2 environment (200sccm) at ~1.0 Torr for 60 min. The grown
substrate was cooled down under H2environment. The grown InAs NWs were
harvested by sonication in anhydrous ethanol
and then were drop-casted onto a heavily
doped Si wafer (with 50 nm thermal grownoxide on top) for fabricating back-gatedfield-effect transistors (FETs). The electrodepositions were defined by standard
lithography process. Followed by the 5 s of
HF (1%) treatment (to remove the nativeoxide on the NW), the substrate was put into
vacuum chamber for Ni (50 nm) deposition
and then lift-off. Electrical characterization
was performed in a cryogenic probe station(Janis ST-500) with a semiconductor
parameter analyzer (Agilent 4155C).
Transmission electron microscope (TEM)
analysis was performed in JEOL 2100F by
drop-casting the harvested NWs onto a
copper grid.

2. Experimental

3. Results and Discussion

InAs NWs investigated here wereobtained by a solid-source catalytic chemical
vapor deposition (SSCVD). Briefly, 1.2 g of
InAs powder (99.9999% purity) and Ni
catalyst (0.2 nm thick on the SiO2/Si
substrate and pre-annealed for 10min at 800oC) were heated at the upstream (690 oC) anddownstream (470 oC) zone of a tube furnace,

The obtained NWs from SSCVD are
long, dense, without tapering and free fromsurface coating on the substrate (Fig. 1a).
Such NWs are ideal for studying thecorresponding electrical properties because
any tapering and surface coating would
degrade the mobility of NWs. The average
diameter of these NWs is 31 ± 6 nm (Fig.

Figure 1. (a) SEM image of grown InAs
NWs (scale bar = 2 μm). (b) TEM image of
InAs NWs (scale bar = 200 nm). Insert isthe diameter distribution of 100 NWs. (c)XRD spectrum of the grown InAs NWs. (d)HRTEM image of an InAs NW (scale bar =5 nm). Insert is the relative FFT reciprocallattice spots. [1] © 2013 IOP Publishing

Figure 2. (a) Peak field-effect mobility as afunction of surface roughness for NWs. (b)Peak field-effect mobility as a function ofdiameter for two different surface roughness
ranges. [1] © 2013 IOP Publishing1b), based on the TEM investigation of 100

NWs. All NWs obtained are zincblend in the
crystal structure (Fig. 1c) and the dominant
growth direction is <110>. A conformal
native oxide layer (~3 nm) is observed which
will be subtracted from the diameter formobility calculation (Fig. 1d). Afterinvestigation of different NW devices withwide range of diameter and roughness, themobility is observed to drop monotonicallywith increasing roughness with similar
diameters (Fig. 2a); on the other hand, thick
NWs have higher mobility than those of thin
NWs when they have similar roughness.
Furthermore, thinner NWs are moresensitive to the impact of surface roughness(Fig. 2b).
Low-temperature electrical measurements
further emphasis the effect of surface
roughness on mobility degradation bydecoupling the effect of phonon scattering

Figure 3. (a) Surface-dependent peakfield-effect mobility of NWs with thediameter of 26.9 ± 0.7 nm measured at 77 K
and 298 K. (b) NW diameter-dependent peakfield-effect mobility of NWs with the surface
roughness of 2.0 ± 0.2 nm measured at 77 K
and 298 K. [1] © 2013 IOP Publishingand thermal activated surface/interface traps.
At 77 K, the mobility of all devices is
improved. The improvement is more
profound for smoother NWs. On the other
hand, the field effect mobility has lessimproved at 77 K for thinner whichhighlights the dominant role of surface
roughness scattering for thinner NWs
because of the higher surface-to-volume
ratio. All the experimental results highlightthe impact of enhanced surface roughness tothe field-effect mobility in miniaturized andsurface disordered NWs.

Conclusion
In conclusion, this study presents a detailelectrical characterization to investigate theeffect of surface roughness on the mobilityof InAs NW FETs. With the surface
roughness increases, the mobility of the FET
decreases monotonically. This effect is more
profound with smaller diameter. Low
temperature electrical characterizations
further decouple the phonon scattering andsurface/interface traps degradation andfurther highlight the dominant role of surface
roughness induced degradation on mobility
in both miniaturized and surface disordered
NWs.
Reference[1] Wang F.Y., Yip S. P. Han N. Fok K. W.,
Lin H., Hou J. J., Dong G., Hung T. F., Chan
K. S., Ho J.C. "Surface roughness induced
electronmobilitydegradation

nanowire", Nanotechnology,

24,

in

InAs

375202,

2013. doi:10.1088/0957-4484/24/37/375202

pH-responsive Graphene in Aqueous Suspensions
Ming Fang, Johnny C. Ho*
Department of Physics and Materials Science, City University of Hong Kong,
83 Tat Chee Avenue, Kowloon Tong, Kowloon, Hong Kong
TEL&FAX: +852-3442 4897/0538*
E-mail: johnnyho@cityu.edu.hk

Abstract
Homogeneous aqueous suspensions of graphene were prepared by chemical reduction ofgraphene oxide in the presence of chitosan. Strong noncovalent interactions between thegraphene sheets and the chitosan molecular chains allow reversible manipulation of the sheets
between well dispersed and more aggregated states by using pH-change as a stimulus. The
strategy of using stimuli-responsive polymers to functionalize graphene opens up newopportunities for rational design of smart graphene based materials.

1. Introduction
Graphene has attracted huge attention
recently due to its exceptional electronic,
mechanical and thermal properties as well as
great application prospects.[1–4] So far,graphenehas
been
preparedbymicromechanical cleavage,[5] chemical
vapor deposition[6] and chemical reductionof graphene oxide (GO).[7] Unfortunately,
as-produced graphene generally shows
limited solubility in solvents, which hinders
its practical applications in certain areas.
This way, chemical modifications could
serve as a powerful tool to solve thismentioned
problem.
Particularly,
noncovalentmodifications
using
dispersants/surfactants are often as usedbecause they are not only simple for
operation, but also cause no damage to thebonding structure of graphene. Various
polymers and surfactants such as PSS,[8]PmPV,[9] and TWEEN[10] have been used
to facilitate the dispersion of graphene insolutions. However, to our knowledge, little
attention has been paid to the utilization of
stimuli-responsive polymers as modifiers
which could possibly endow graphene withmore interesting properties.
Here, we demonstrate the use of chitosan,
a pH-responsive biopolymer, as anoncovalentmodifier
toobtainhomogeneous suspensions of graphene andshow that the chitosan modified graphenesheets can be reversible manipulated

between well-dispersed and aggregatedstates by applying pH-change as a stimulus.

2. Experimental
Graphite oxide was synthesized fromgraphite powder (<30μm) by a modified
Hummers method.[11] A stock suspension of
GO was obtained by sonication of thegraphite oxide in water.[7] The GO
suspension (100mL, 0.1mg/mL) was thengradually added to an acidic aqueous
solution of chitosan (100mL, 0.5mg/mL, in
0.5 M HCl) with stirring to yield a uniform
brown-colored dispersion. L-ascorbic acidpowder (100 mg) was subsequently added tothe above dispersion and the mixture washeld at 60 °C for 6 h to reduce the GO into
graphene.

3. Results and Discussion
(a)

(b)

Fig. 1 (a) Aqueous suspension of exfoliatedgraphite oxide. (b) Typical AFM image of GO
sheets deposited on a mica surface

Graphite oxide can be easily exfoliated andbecome monolayers in water undersonication. Obtained GO suspensions arebrown colored and highly stable and AFM
images show that the sheets are about 1.2 nm
in thickness (Fig.1).

(a)

(b)

highly dispersed

agglomerated

Fig. 2 Schematic illustration of the influence ofthe mixing manner on the properties of the
mixtures. (a) Adding GO into a solution of
chitosan produces a homogeneous dispersion.
(b) Adding chitosan into a GO suspension leads
to agglomeration of the GO sheets.

The GO suspensions were then used for
preparations of GO/Chitosan solutions.
Interestingly, we found that the mixing
manner have huge influence on properties ofthe mixtures. Homogeneous GO/chitosan
solution was prepared by simply adding asuspension of GO into a solution of chitosan
under stirring. However, when we mix thetwo component solution in the opposite
order, that is, adding a solution of chitosan
into a suspension of GO, we get irreversible
agglomerations. It is known that GO bears
lots of carboxylic and hydroxyl groups andchitosan has abundant amino groups along
its molecular chains. Ionizations of thesefunctional groups in water make GO andchitosan charged, negatively for the former
and positively for the latter, which was
confirmed by ζ potential. While the GO
suspension had a negative ζ potential of
about -46 mV, the solution of chitosan held apositive ζ potential of about +100 mV. Once
GO and chitosan meet with each other insolution, strong ionic interactions take place
and yielded a new balanced system. The
distinct dispersion states of GO sheets indifferent mixtures can be explained throughthe following model. As illustrated in Fig. 2,
when GO was added to the solution of
chitosan, chitosan is in mass relative to GO,
and every GO sheet was immediately
wrapped by large amounts of chitosan
molecular chains through electrostatic
attractions (ionic interactions), by which the
GO sheets were isolated as individuals and a
uniform dispersion formed. For mixing thetwo solutions in opposite order, there was alarge excess of GO relative to the chitosan,
and one chitosan polymer chain can possiblyattached two or more GO sheets, leading to“bridging” flocculation.
Generally, GO can be chemically reduced
into graphene. However, direct reduction of
GO causes irreversible aggregations of the
product, which is not favorable for its
processing and applications. In this work, weused chitosan as the dispersant, and reduced

Absorbance (a.u.)rGO

Intensity (a.u.)

(b)

(a)

GO
rGO

GO
200

300
400
500
Wavelength (nm)

600

1320

1650

1980

-1

Raman Shift (cm )Fig. 3 (a) UV-vis absorption spectra of aqueous dispersions of bare GO and the reduced GO withchitosan. The inset shows the photographs of the dispersions. (b) Raman spectra of the GO and rGO.

the GO in situ by adding L-ascorbic acid(vitamin C) into a well-dispersed
GO/chitosan suspension. After reduction, thesuspension changed its color from lightbrown to dark, and no visible aggregation
observed (Fig.3a, inset). The reduction was
confirmedbyUV-visabsorption
spectroscopy and Raman spectroscopy. As
shown in the UV-vis absorption spectra (Fig.
3a), after the reaction, the maximum
absorption of the GO dispersion red-shiftedfrom 230 to 270 nm, and meanwhile theabsorption intensity in the entire spectrum(especially the region above 300 nm)
dramatically increased, suggesting therestoration of the electronic conjugations inthe sheets.[12] In the Raman spectra (Figure3b), the reduced GO (rGO) gave a muchhigher intensity ratio of D band to G band
(D/G, where the D band is induced by
defects and the G band is due to the in-planevibration of sp2 carbon atoms[13]). Since theRaman D/G intensity ratio is proportional tothe average size of the sp2 domains, the riseof the D/G intensity ratio was actually
attributed to the increase in number of small
in-plane sp2 domains, confirming theeffectiveness of the conventional reduction.

(a)

(b)

Fig. 4 SEM images of (a) rGO sheets fromreduced, chitosan-containing solution, and (b)GO reduced with no additional dispersants.

SEM was utilized to analyze thedispersing state and morphology of the rGO
in the suspension by dropping the suspension
on silicon wafers. As shown in Fig.4,
plentiful “transparent” separated rGO sheets
appeared on the substrate, suggesting therGO sheets in the suspension were welldispersed. In contrast, reduction of a GO
without chitosan under a same condition
produced black clusters which exhibited ahighly agglomerated morphology in SEM.
This indicates the chitosan played an
important role for the dispersion of the rGO
in water.

Fig. 5 AFM image of the rGO sheets from thechitosan-containing solution. The sample wasrinsed by 0.1% HCl after deposition.

It is well known that, during reductions,
GO gradually loses its oxygen-containing
groups (OCGs) and becomes hydrophobic,
so that direct reduction of GO in water
without proper mediating agents generally
causes irreversible aggregations. On theother hand, reduction of GO does not
generate perfect graphene[12], and in thepresence of chitosan, the residual OCGs of
the rGO had the opportunity to capture theamino and hydroxyl groups of chitosan via
zwitterionic interaction and hydrogenbonding and endowed with dispersibility inwater. AFM showed that the rGO were about
3 nm in thickness (Figure 5), which is muchthicker than GO (1.2nm) and the reported
value for clean rGO (1 nm), indicating that
there was a layer of chitosan anchored on thesides of the rGO sheets.
Chitosan is a weak polyelectrolyte andits molecular conformation and solubility
can be reversibly altered by pH changes;
therefore, noncovalent bonding of the rGO
with chitosan offers the possibility ofcontrolling the behaviors of the rGO in
suspensions by using pH as a stimulus. For
practical
experiments,
rGO/chitosan
suspensions were tuned to different pHvalues (from 2 to 12) and theirtransmittances (at 500 nm) were recorded bya
UV-visspectrophotometer.
The
transmittances were plotted against the pH
values (Figure 6a), which clearly shows apH-dependent dispersibility of the rGO in
the suspensions. Sharp transitions fromdispersed to agglomerated states wereobserved when the pH changed from 6 to 7,
which is consistent with ζ potential
measurements (Fig. 6b).

Transmittance (%)
Zeta Potential (mV)

(a)

80

[4] A.A. Balandin, S. Ghosh, W. Bao, I.
Calizo, D. Teweldebrhan, F. Miao, and40

C.N. Lau, Nano Lett. 8 (2008) 902.
[5] K.S. Novoselov, A.K. Geim, S. V

0

2

4

6 pH 81012Morozov, D. Jiang, Y. Zhang, S. V

(b)

30

0

Dubonos, I. V Grigorieva, and A.A.
Firsov, Science 306 (2004) 666.
[6] X. Li, W. Cai, J. An, S. Kim, J. Nah, D.

2

4

6

8

10

12

pH

Fig. 6 (a) Transmittance of the supernatants ofrGO/chitosan suspensions with different pHvalues after sedimentations. (b) Zeta potentialof the rGO/chitosan suspension at different pH.

We also found that the agglomerated rGO
can be redispersed when the suspension was
tuned to acid conditions again, indicatingthat the dispersion/aggregation of the rGO in
the suspension could be controlled by using
pH change as a stimulus.
4. Conclusion
We demonstrate the preparation ofhomogeneous aqueous suspensions ofgrapheme, which is prepared by chemical
reduction of GO in the presence of chitosan.
The graphene in the suspensions can bereversibly switched between well-dispersed
and aggregated states by changing the pH
value. The study may serve as a prototype
for future studies on using stimuli-responsive
moieties to functionalize graphene foradvanced applications.

Reference[1] A.K. Geim and K.S. Novoselov, Nat.
Mater. 6 (2007) 183.
[2] C. Lee, X. Wei, J.W. Kysar, and J.
Hone, Science 321 (2008) 385.
[3] K.I. Bolotin, K.J. Sikes, Z. Jiang, M.
Klima, G. Fudenberg, J. Hone, P. Kim,and H.L. Stormer, Solid State Commun.
146 (2008) 351.

Yang, R. Piner, A. Velamakanni, I.
Jung, E. Tutuc, S.K. Banerjee, L.
Colombo, and R.S. Ruoff, Science 324(2009) 1312.
[7] S. Stankovich, D.A. Dikin, R.D. Piner,
K.A. Kohlhaas, A. Kleinhammes, Y. Jia,
Y. Wu, S.T. Nguyen, and R.S. Ruoff,
Carbon 45 (2007) 1558.
[8] S. Stankovich, R.D. Piner, X. Chen, N.
Wu, S.T. Nguyen, and R.S. Ruoff, J.
Mater. Chem. 16 (2006) 155.
[9] X. Li, X. Wang, L. Zhang, S. Lee, and
H. Dai, Science 319 (2008) 1229.
[10] S. Park, N. Mohanty, J.W. Suk, A.
Nagaraja, J. An, R.D. Piner, W. Cai,
D.R. Dreyer, V. Berry, and R.S. Ruoff,
Adv. Mater. 22 (2010) 1736.
[11] W.S. Hummers and R.E. Offeman, J.
Am. Chem. Soc. 80 (1958) 1339.
[12] D. Li, M.B. Müller, S. Gilje, R.B.
Kaner, and G.G. Wallace, Nat.
Nanotechnol. 3 (2008) 101.
[13] M.S. Dresselhaus, A. Jorio, M.
Hofmann, G. Dresselhaus, and R. Saito,
Nano Lett. 10 (2010) 751.

Growth of Ca Silicide Layers by Thermal Treatment of Si Substrates in Ca Vapor
Yusuke Sato*1
1

Faculty of Engineering, Shizuoka University, Hamamatsu, Shizuoka, Japan*e-mail: f0911064@shizuoka.ac.jp

Abstract
Ca silicide layers were formed by thermal treatment of Si substrates in Ca vapor.
I examined the difference in structure of two samples, which were synthesised by different
growth time. The structural properties of the resultant composites were characterized by fieldemission scanning electron microscopy (FE-SEM) and X-ray energy dispersive spectroscopy(EDS).

1.

Introduction

Recently, semiconducting silicides haveattracted much attention for their potential tocreate new classes of environmentallyconscious
electronics.
Silicon-based
nanoscale materials, such as nanotubes,
nanowires,andnanoparticles,
haveremarkable electronic and optical properties
and are suitable for nanodevice applications
owing to their extraordinary structures.
Whilezeroorone-dimensional
nanomaterials may be suitable for nanoscale
fabrication, two-dimensional nanomaterials
can bridge the gap between the quantum
world and three-dimensional bulk because of
their nanoscale thickness and microscale area.
To synthesize two-dimensional nanomaterials,
chemical exfoliation is performed artificiallyfor several classes of layered materials bycertain soft-chemistry procedures. In adition,
many researchers have attempted to prepare
two-dimensional silicon sheets, there havebeen no reports of a successful fabrication ofsilicon monolayer sheets.
One of the method for silicon monolayer
sheets synthesis, it is chemical exfoliation of
calcium disilicide [1]. Therefore, as a firststep for synthesized silicon monolayer sheets,
we have a purpose of growth of Ca silicide
layers by thermal treatment of Si substrates in
Ca vapor.

2.

ExperimentFirst, we prepared Ca ball (99.5 %, 1-3mm) and Si (111) substrate. Ca ball was
adjusted to be 0.05 grams. The Si (111)substrate was etched in HF solution(HF:HNO3:CH3COOH = 1:2:1) and rinsed
with deionized H2O. Next, the substrate was
immersed in a diluted hydrofluoric acidsolution (HF:H2O = 1:9). Then Ca ball andSi substrates were located in a loosely sealed
quartz container, which was loaded into a
vacuum chamber. Then the container washeated up and the temperature wasmaintained at 600ºC.
During the heat treatment, Si substrates
were exposed to Ca vapor, after which
Ca-Silicide layers were formed on the Si
substrates. We synthesized two samples bychanging the time for the heat treatment. One
is the sample(a) of the growth time of 15
minutes. The other one is the sample(b) of
the growth time of 30 minutes.
The structural properties of the resulting
structures were characterized by X-ray
diffraction (XRD), field emission scanningelectron microscopy (FE-SEM) and X-ray
energy dispersive spectroscopy (EDS).

3.

Results and Discussion

Figure 1 shows the SEM images ofsample(a) and sample(b). And Table 1 shows
results of EDS analysis of sample(a) andTable 2 shows results of EDS analysis ofsample(b).

Fig.1 : SEM image of sample(a) (growth time:15min) and sample(b) (growth time:30 min).

4.

Conclusion

We tried synthesized Ca silicide on Si
substrates. As a result, Ca silicide had
synthesised by thermal treatment of Si
substrates in Ca vapor. In addition, we
considered another material has beensynthesized in the lower layer and the upper
layer of the layer. In future, to find the
optimum growth time and growth condition
for synthesis Ca silicide.

Table1 EDS analysis of sample(a)5.

Reference[1] Hideyuki Nakano, Takuya Mitsuoka,
Masashi Harada,Kayo Horibuchi, Hiroshi
Nozaki, Naoko Takahashi,Takamasa Nonaka,
Yoshiki Seno, and Hiroshi Nakamura Angew.
Chem. Int. Ed. 2006, 45, 6303 –6306

Table2 EDS analysis of sample(b)From figure1, We confirmed that the layer
was synthesized on a Si substrate. And we
can be confirmed that the sample(b) is thicker
than sample(a). In addition, from table1andtable2, the layer can be seen that it iscomposed of Ca and Si. Therefore, We
considered the calcium silicide has beensynthesized on a Si substrate.
From sample(a) and sample(b), we can be
confirmedthat the atomic ratio of Ca to Si isapproximately 1:2 near the boundary of thesubstrate and layer. We considered anothermaterial has been synthesized in the lower
layer and the upper layer of the layer.

Thermal Annealing of Si3N4 with Gallium
Naoki Yamamoto
Faculty of Engineering, Shizuoka University
3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka 432-8561, Japan
TEL&FAX: +81-53-478-1099
E-mail: f0711151@ipc.shizuoika.ac.jp

Abstract
Thermal annealing of Si3N4 with gallium in a nitrogen gas atmosphere. Wire is generated on theupper surface of the gallium. The structural properties of the resultant composites werecharacterized by field emission scanning electron microscopy (FE-SEM), and X-ray diffraction(XRD).

1.

Introduction

Recently, gallium nitride attracts muchattention in various fields. For example, it isused in car headlights and signal using a lightemitting diode. This is a material necessary to
the present life, which is to enrich life. Other
than a light-emitting diode, the application toa high power transistor solar battery is
expected.
In addition, Photon Enhanced Thermionic
Emission (PETE) was an opportunity to begin
the study.[1] Photon Enhanced Thermionic
Emission is a newly proposed form of solar
energy harvesting which relies upon a
combination of quantum and thermal
processes to generate electricity. Whileminimizing costs and integration with existing
infrastructure solar, PETE module canincrease the power generation efficiency in apractical scale. Thus, it is possible to reducethe cost of power generation.
The purpose of this study, gallium nitride
will be prepared by a simple way, and making
a cheap solar battery.

2.

Synthesized sample was heated at 900ºC for
5 hours in an electric furnace. Electric furnace
was filled with nitrogen when heated.
The structural properties of the resulting
structures were characterized by X-ray
diffraction (XRD), field emission scanningelectron microscopy (FE-SEM).

3.

Results and Discussion

Figure 1 shows the appearance of thesample after heat treatment. Wire had grownfrom the surface of the sample.

Experimental

The composites were synthesized fromliquid gallium (purity 99.99%) and silicon
nitride powder (Si3N4). Silicon nitride was so
as to be covered by the gallium at this time.

Figure 1, Appearance of the sample

4.

Conclusion

Gallium and silicon nitride were heatedat 900ºC for 5 hours in an electric furnace.
Wires are grown on the surface of thesample. Ga2O3 was generated from theresults of XRD.
GaN was not able to be formed on thisexperiment condition. Therefore, it isnecessary to examine the growthtemperature and the growth time GaN isgenerated in the future.
Figure 2, XRD spectrum of the sample

5.

Reference[1] Nicholas Melosh. Zhi-Xun Shen
Photon Enhanced Thermionic Emission
for Solar Energy Harvesting
April 20, 2012

Figure 3, SEM image and EDS analysisof the sample
Figure 2 shows XRD spectrum of part of thesample. Ga2O3 was included in the sample
from the result. However, GaN was notincluded in the sample from the result.
Figure 3 shows SEM image and EDS
analysis of the sample. It is possible that
nitrogen is lost from the silicon nitride fromthe results of EDS. As a result, it is consideredthat it has changed to the silicon.

Deposition of Silicide-based Microstructures on Ca-silicide/Si substrates
1

Kaito Nakane*1
Graduate School of Engineering, Shizuoka University, Hamamatsu, Shizuoka, Japan*e-mail: f0230119@shizuoka.ac.jp

Abstract
Various Silicide-based microstructures were formed by CrCl2 treatment of Ca-silicide/Si
substrates. The CrCl2 treatment time affects the structural property of micro/nano-structures. The structural properties of the resultant composites were characterized by fieldemission scanning electron microscopy (FE-SEM) and X-ray energy dispersive spectroscopy(EDS).

1. Introduction
Recently,
semiconducting
silicides,
which consist of non-toxic and abundantmaterials at low cost, have attracted muchattention for their potential to create newclasses of environmentally conscious
electronics. In addition, low-dimensional
materials have attracted much interest due totheir enhanced or modified optical,
electronic and mechanical properties
compared to those of bulk materials.
It has been reported that synthetic
method of the silicon nano sheet which usedpure Cl2 and Metal-chlorides, such as SbCl3,
FeCl3, S2Cl2, PCl5, SnCl2, however, thestructures of the sheets have not beenclarified yet [1-2]. Using metal chlorides, theformation of silicides and siliconnanostructures
and
their
structuralmodification
have
been
intensively
investigated. CrCl2 is one of appropriate
metal chlorides source materials tosynthesize silicide nanostructures [3].
This paper shows Synthesis of
micro-/nano-structures by CrCl2 treatmentof Ca - silicide/Si substrates. The structuralproperties of micro/nano -structures were
investigated.

First,

2. Experimental
we prepared Ca-silicide/Si

substrates with starting materials of Ca ball
(99.5 %, 1-3 mm) and Si (111) substrate. Caball and Si substrates were located in aloosely sealed quartz container, which was
loaded into a vacuum chamber. Then thecontainer was heated up and the temperature
was maintained at 600ºC. During the heat
treatment, Si substrates were exposed to Ca
vapor, after which Ca-Silicide layers wereformed on the Si substrates.
For the CrCl2 treatment, the reaction
apparatus was flushed with argon gas toremove any oxygen and moisture in the
furnace. After the furnace reactor was heated,then the Ca-silicide/Si substrates and 0.08 gof the CrCl2 source materials were placed in
the reactor. The CrCl2 source material was
evaporated at 872 ºC, then reacted with theCa-silicide/Si substrates in the range of 650ºC and maintained for 2 - 15 min, thencooled to room temperature without the flowof argon gas [3].
The morphological and structuralproperties of the resulting structures werecharacterized by X-ray diffraction (XRD),
field emission scanning electron microscopy(FE-SEM) and X-ray energy dispersive
spectroscopy (EDS).

3. Results and Discussion
Figure 1 shows the SEM images ofSample grown at 2 min CrCl2 treatment.
From figure 1(a), it is found that large
quantity of white nanostructures were
deposited on the sample surface. In addition,
white nanoparticles 0.5-0.8 μm in diametercan be observed, as show in the Fig.1 (b).
Moreover, it is found that microwires weregrown with nanoparticles.
Figure 2 and Table 1 shows results of
SEM image and EDS analysis of Sample

Fig. 1.

SEM image of nanoparticles (CrCl2

Fig. 3.

SEM image of microrods (CrCl2

treatment time:2min)treatment time:15 min)Fig. 2. SEM image and corresponding EDS
mappings of the nanoparticles grown at 650 ºC, 2
min CrCl2 treatment
Table1 EDS analysis

Fig. 4. SEM image and corresponding EDS
mappings of the microrods grown at 650 ºC, 15
min CrCl2 treatment
Table2 EDS analysis

Element

Ca

Si

Cr

Cl

O

Element

Ca

Si

Cr

Cl

O

[at%]

3.12

0.60

39.50

0.56

56.21

[at%]

6.01

36.47

7.13

6.91

43.47

grown at 2 min CrCl2 treatment. This resultshows that the composition of thenanoparticles were mainly consist of Cr andO.
Figure 3 shows the SEM images ofSample grown at 15 min CrCl2 treatment. As
shown in Fig.3 (a), a large number ofmicrorods were formed on the sample
surface. From figures 1(b), microrods 3-4
μm in length and nanoparticles wereobserved. In addition, it is found that wires
are bundle. Moreover, nanoparticles werelocated on the surface of the nanorods andon the top of the wires.
Figure 4 and Table 2 shows results of
SEM image and EDS analysis of Sample
grown at 15 min CrCl2 treatment. This resultshows that the composition of the
microwires is mainly consist of Si and O.

4. Conclusion
Compiteswithvariousmicro-/nano-structures were synthesized by
CrCl2 treatment of Ca-Silicides/Si substrate.
The CrCl2 treatment affect the structuralproperty of resulting comppsotes.
In future, Well-aligned silicide-based
micro- /nano-structures will be grown andtheir electrical and optical characterizations
will be expected.

Reference[1] H. Kautsky and L. Haase, Chem. Ber. 86(1953)1226

[2] E. Bonitz, Chem. Ber. 94(1961)220[3] W. Li, E. Meng, T. Matsushita,
S. Oda, D. Ishikawa, K. Nakane,
J. Hu, S. Guan, A. Ishida andH Tatsuoka, J. Cryst. Growth365(2013)11

Synthesis of Si nanowires using MnCl2 and Si powders source on Si substratewith Au catalyst
Erchao Meng*
Graduate School of Science and Technology, Shizuoka University, 3-5-1 Johoku,
Naka-ku, Hamamatsu 432-8011, Japan
TEL: +81-53-478-1099, Fax: +81-53-478-1099
E-mail: f5145034@ipc.shizuoka.ac.jp
Abstract

High-density and high-quality Si nanowires were synthesized using MnCl2 and Si
powder on the Si substrates with Au catalyst. As comparison, under the same growth
condition, when Si powder was used as source materials, there was no nanowires grown. The
Si nanowires grown are characterized by SEM, TEM and STEM with EDS.

1. Introduction

Recently, novel devices, based on
nano-materials, have attracted a great dealof attention because of their potentials formany industrial applications. Therefore, thehuge efforts have been devoted tonanomaterials researches. Si nanowries havegood potentials for application, thus, whichhave been synthesized by many methods
such as chemical vapor deposition [1] andhigh temperature thermal evaporation [ 2, 3].
On the other hand, many kinds of Si
precursors were used for Si nanowires
growth such as SiH4 and Si2H6, and so on[4]. However, for these growth methods andSi precursor, usually complex processes, thetoxic and dangerous precursor gases or thehigh reaction temperature are needed.
Therefore, it is necessary to develop a new
synthesis method of Si nanowires.
In this study, Si nanowires weresynthesized using MnCl2 and Si powder assource materials on the Si substrate, and Au
is used as catalyst for Si nanowries growth.
2. Experiments
The nanowires were synthesized on
Si(111) substrates. The Si(111) substrates
was degreased with solvent, etched in HF
solution (HF:HNO3:CH3COOH = 1:2:1) andrinsed with deionized H2O. Next, thesubstrates were immersed in a dilutedhydrofluoric acid solution (HF:H2O = 1:9),
and then dried prior to loading into a
vacuum for Au deposition. The Si substrates
were prepared with Au deposition with the
thickness of about 3 nm as a catalyst.
On the other hand, MnCl2 and Si
powders with a mole ration of 1:3 were usedas source materials. The Si substrates andthe source materials were located in aloosely sealed quartz container, which was
loaded into a vacuum chamber. Then thecamber was pumped down to a basepressure of 10-3 Pa. Then the container washeated up and the temperature wasmaintained at 700-850
for 60 min. A
schematic illustration of the growthequipment used in this study was shown in
Ref. [5].
The morphological and structuralproperties of the resultant nanostructures
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 SEM images of the resultant
products synthesized using (a) MnCl2 andSi powder source (b) Si powder source
Figure 1(a) shows SEM image of Si
nanowires, which were synthesized using
MnCl2 and Si powder source on the Si
substrates with Au catalyst. The source andsubstrate temperatures were maintained for
1 h at 850 and 700 . As comparison,
Figure 1(b) shows SEM image of theresultant product synthesized using Si
powder under the same growth condition.

From SEM images, it can be seen that whenusing MnCl2 and Si powder as source
materials, Si nanowires with high-density
and high-quality were grown on the Si
substrates. However, using Si powder
source, just some particles were grown on
the Si substrate, no nanowires wereobserved. The results indicated that MnCl2
plays an important role in fabricating Si
nanowires.

Figure 2 SEM, STEM images and EDS
mappings of the Si nanowires grown on thesubstrate with Au catalysts using MnCl2 andSi powder source.
Figure 2 shows the SEM, STEM
images and EDS mappings of the Si
nanowires grown on the substrate with Au
catalysts using MnCl2 and Si powder source.
Figure 2(a) shows the conventional
columnar Si nanowires with high-density
were grown, and the catalyst particles can
be clearly observed on the top of Si
nanowires. The STEM image in Figure 2(b)shows that the surface of Si nanowires issmooth, and the diameter of nanowires isalmost the same as that of catalyst particles
on the top of nanowires.
EDS mapping shows that Si wasdistributed on the entire nanowire, while Au
was mainly concentrated in the top of
nanowire, which suggested that Si
nanowires were synthesized by Au catalyst
via VLS growth mechanism. It is also
observed that oxygen and manganese atoms
were detected in the catalyst particle, andfew oxygen is detected in the nanowire.
Manganese atoms is adsorbed into Au-Si
solution, and is oxidized, which suppressed
the oxidation of Si nanowires.
Figure 3 shows TEM image, HRTEM
image and the corresponding FFT pattern of
Si nanowires, which were synthesized using
MnCl2 and Si powder as the source
materials. TEM image shows that thestraight and smooth surface Si nanowires
with hemispherical catalyst particles on the

top of nanowires were grown. HRTEM
image shows that the lattice spacings
between two neighboring fringes measured
are 0.331nm and 0.201nm, and the angle
between these two planes is 35º, whichcorresponds to those of the (111) and (220)planes of Si. It demonstrated that Si
nanowires were synthesized, which isconsistent with the mapping result shown in
Figure 2. The corresponding FFT patternindicated that Si nanowire was grown along
the [111] direction.

Figure 3 TEM image, HRTEM image andthe corresponding FFT pattern of Si
nanowires synthesized using MnCl2 and Si
powder source on the Si substrate with Au
catalyst.
It is reported that when Si powder wasused as raw materials for growth of silicon
nanowires, the evaporation temperature ofsource materials used in experiments ismore than 1000 in order to fabricate Si
nanowires [6]. However, here in our study,
the source temperature is 850 (actually,
according to our experimental results, Si
nanowires can be fabricated at lower source
temperature), which is lower than theevaporation temperature of Si powder. Thus,
Si powder can not be evaporated, and there
is no nanowires grown on the Si substrate,
as shown in Figure1(b).

4. Conclusions

High-density and high-quality Si
nanowires were synthesized using MnCl2and Si powder on the Si substrates with Au
catalyst at lower source temperaturecomparing with Si powder source. This
simple and low-cost growth technique is
expected to be applied to othersemiconductor nanostructures.

References[1] Nong M. Hwang, Woo S. Chenong, Duk
Y Yoon and Doh-Y. Kim, Journal of Crystal
Growth 218 (2000) 33.
[2] Y. Shi, Q. Hu, H. Araki and T. Noda etal., Applied Physics A, 80 (2005) 1733.

[3] A. Colli, A. Fasoli, P. Beecher and J.
Robertson et al., Journal of Applied Physics
102 (2007) 034302.
[4] W. Molnar, A. Lugstein, P. Pongratz, N.
Auner, C. Bauch and E. Bertagnolli, Nano
Lett. 10 (2010)3957-3961.
[5] Y. Mizuyoshi, R. Yamada, T. Ohishi, Y.
Saito, T. Koyama, Y. Hayakawa, T.
Matsuyama and H. Tatsuoka, Thin Solid
Films 508 (2006)70.
[6] Y. Shi, Q. Hu, H. Araki, H. Suzuki, H.
Gao, W. Yang and T. Noda, Applied
Physics A, 80(2005) 1733-1736.

Effect of Nickel Addition on Growth Evolution of Silicon Nanostructures using
MnCl2 and Si powders
Hiroaki Suzuki
Faculty of Engineering, Shizuoka University
TEL&FAX: 053-478-1099
E-mail: f0911042@ipc.shizuoka.ac.jp

Abstract
Faceted nanowires can be fabricated with Vapor-Liquid-Solid-Solid (VLSS) Growth
Mechanism and square shape nanowire using Si powders and MnCl2 as source is reported. In
this paper, I added Nickel as impurity in the source and investigated the effect of Nickel
addition.

1.

Introduction

Thermoelectric device is generic name ofthe element, using a phenomenon to connect
electricity with heat. For example power
generator, heating and cooling applications.
Recently, attention has focused on thethermoelectric device as a power generator
because of diverse reasons. For example, it
can be small, doesn’t need external source,
etc. Therefore, thermoelectric device has
been researched for a long time and improved
efficiency.
The efficiency of thermoelectric material
is represented by its dimensionless figure of
merit, ZT, which is calculated by thefollowing relational expression:=+
where S is the Seebeck coefficient, σ isthe electrical conductivity, T is thetemperature, and κlattice and κelectronic are thelattice and electronic contributions to thethermal conductivity, respectively. In thisexpression, Seebeck coefficient, electricalconductivity
andelectronic
thermalconductivity have influences each other, but
only lattice thermal conductivity doesn’t haverelation with other parameters because it
depends on phonon although others dependon electron. It means that, to lower the lattice
thermal conductivity makes high ZT and tochange the structure can make low lattice
thermal conductivity.
The fabrication of faceted squarenanowire with Vapor-Liquid-Solid-Solid(VLSS) Growth Mechanism using Si
powders and MnCl2 as source is reported [1].
Therefore in this paper I aim at theinvestigation of the effect of Nickel addition
to the source.

2.

Experiment

The Si nanostructures were synthesizedon Si (111) substrates. The Si (111)substrates was degreased with solvent, etchedin HF solution (HF:HNO3:CH3COOH =1:2:1) and rinsed with deionized H2O. Next,
the substrates were immersed in a dilutedhydrofluoric acid solution (HF:H2O = 1:9),
and then dried prior to loading into a vacuum
for Au deposition. The Si substrates wereprepared with Au deposition with the
thickness of about 3 nm as a catalyst.
On the other hand, Si powders, MnCl2
and Ni with a mole ration of 1:1:0.1 and
1:1:0.02 were used as source materials. The
Si substrates and the source materials werelocated in a loosely sealed quartz container,
which was loaded into a vacuum chamber.
Then the camber was pumped down to a basepressure of 10-3 Pa. Then the container washeated up and the temperature wasmaintained at 700-900 oC for 25min.
The structural properties of the resultant
nano/microstructures were characterized byscanning electron microscopy (SEM).

3.

Results and Discussion
4.

Fig. 1 SEM images of (a) plan-view and (b)angle images (source materials mole ration =1:1:0.1)

Figure 1 shows SEM images of plan-view
and angle images (source materials moleration = 1:1:0.1). In this figure, I observed
triangular structures on the surface. And theheight of structures was approximately 1 μm.
From this result, I thought Ni had a large
effect. So next experiment, I used less Ni.

Fig. 2 SEM images of plan-view images(source materials mole ration = 1:1:0.02).

Figure 2 shows SEM images of plan-view
images (source materials mole ration =1:1:0.02). I found triangle-like hexagonal
microrods this time.
From the last result of triangle structuresand microrods at this time, possibly we canget triangle nanowires.

Fig. 3 SEM images of angle images. Source
materials mole ration = 1:1:0.1 (a), 1:1:0.02 (b).

And in these experiments, I came acrossthe nanosheets. Figure 3 shows SEM imagesof nanosheets with source materials moleration of (a) 1:1:0.1, (b) 1:1:0.02. If we use Si
powders and MnCl2 as source, we can hardly
get the nanosheet.

Conclusions

I investigated effect of Nickel AdditiononGrowthEvolutionof
SiliconNanostructures using MnCl2 and Si powders.
Triangleandtriangle-likehexagonal
microstructures
werefabricated,but
nanowires can’t be found. And at the sametime, I came across some nanosheets.
In the future, I want to clarify themechanism of effect of Ni addition andcontrol nanostructures.

References[1] E. Meng, W. Li, K. Nakane, Y.
Shirahashi, Y.Hayakawa, and H.
Tatsuoka, AIP Advances 3, 092107(2013)

Nitric acid treatment of MoSi2 Nanosheets synthesized by Molten Salt Techniqueusing MoS2
Daisuke Ishikawa
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu 432-8561, Japan
E-mail tehtats@ipc.shizuoka.ac.jp

Abstract
The hexagonal phase MoSi2 was formed by the molten salt technique using MoS2
powder as the source material. Nitric acid treatment was performed to the resultant
silicides. The structural and morphological properties of the resultant silicides arecharacterized. 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.
Impurities decreased by performing Nitric acid treatment.
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. The
t-MoSi2 has lattice constants of a=0.3205 nm
and c=0.7845 nm. On the other hand, thehexagonal (h-) phase MoSi2 has lattice
constants of a=0.4622 nm and c=0.6646 nm,
and is known to be a narrow gap
semiconductor. It is reported that the lowtemperature phase, h-MoSi2, is expected to
be one of the novel silicide semiconductors,
which
has
improved
thermoelectric
properties [2]. However, the electrical andthermoelectric property of the bulk crystal
h-MoSi2 have not been experimentally
clarified. Since t-MoSi2 is transformed intoh-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 grown using Mo
substrates and Mo-Ti alloy (Mo(Ti))
compacts by the molten salt method.
However, the tetragonal phase was formed
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 and improved
compositional homogeneity of the silicide
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 wasdemonstrated 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.
Moreover, MoS2 has layered structures. It is
expected MoSi2 synthesized by Molten
salt technique by using MoS2 havenanostructures. Nanostructures haveunique properties and novel applications
compared to those of bulk materials. A
decrease in the thermal conductivity and anincrease in the power factor, σS 2 , of thenanostructures have been demonstrated by thetheoretical calculations and proof of principle
experiments, relative to the corresponding bulk
structures.[10-13]

In this study, the synthesis of h-MoSi2
was examined using MoS2 powder. In orderto remove impurities, Nitric acid treatment
was performed. The structural andmorphological properties of the resultant
silicides were also characterized.

Fig.2 Intensity ratio of MoS2(002) andMoSi2(111) by XRD spectra of the resultant
powder by using S and MoS2 as the source
material at 600 ºC and 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. Then, Nitric acid treatment was
performed to the sample. Figure.1 shows
Nitric acid treatment process. The structuralp roperties of the powder were characterizedby X-ray diffraction (XRD) and transmission
electron microscopy (TEM).

Fig.1 Nitric acid treatment process

Figure.3 shows a TEM image of theresultant powder using MoS2 as the source
material. As shown in Fig.3, plate likestructures were grown, which is due to thelayered structure of the MoS2 as source
material.

Fig.3 TEM image of the resultant powder using
MoS2 as the source material at 600 ºC / 4h.

3. Results and Discussion
Figure.2 shows the intensity ratio of
MoS2(002) and MoSi2(111) by XRD spectraof the resultant powder by using S and MoS2
as the source material at 600 ºC and 700 ºC.
Synthesis of the h-MoSi2 was promoted byadding S powder.

Figure.4 shows XRD spectra of Nitricacid treatment performed to the resultant
powder by using MoS2 as the source material
at 600 ºC. After Nitric acid treatment, thenumber of peaks of impurities decreased.

Fig.4 XRD spectra of the resultant
powder by using MoS2 as the source
material at 600 ºC after Nitric acidtreatment.

4. Conclusion
The h-MoSi2 was grown using MoS2
powder as the source material by the molten
salt method. Moreover, thin silicide firms
were synthesized. Synthesis of the h-MoSi2
was promoted by adding S atom. It is
expected that this simple growth procedure
to fix the crystalline structure using theappropriate starting material. Impurities
decreased by performing Nitric acidtreatment.

Acknowledgements
I would like to thank K. Nakane andH.Koyama and H. Tatsuoka for their fruitful
discussions.

Reference[1] APAC-SILICIDE, July 29-31, 2006,
Kyoto, Thin Solid Films 515(22), pp.
8101-8300, edited by Y. Maeda, K.P.
Homewood, T. Sadoh, Y. Terai, K.
Yamaguchi and K. Akiyama.
[2] T. Nonomura, C. Wen, A. Kato, K. Isobe,

Y. Kubota, T. Nakamura, M. Yamashita,
Y. Hayakawa and H. Tatsuoka, Physics
Procedia 11, 110 (2011).
[3] A. J. Heron and G. B. Schaffer, Mater.
Sci. Eng. A352, 105 (2003).
[4] T. Nonomura, C. Wen, K. Shirai, K.
Isobe, A. Kato, Y. Kubota, T. Nakamura,
Y. Hayakawa and H. Tatsuoka, IOP Conf.
Ser.: Mater. Sci. Eng. 18, 14 (2010).
[5] L. S. Byskov, J. K. Norskov, B.S.
Clausen, and H. Topsoe, J. Catal. 187,
109 (1999).
[6] J. V. Lauriitsen, M. Nyberg, R.T. Vang,
M.V. Bollinger, B.S. Clausen, H. Topsoe,
K. W. Jacobsen, E. Laegsgaard, J. K.
Norskov
andF.
Besenbacher,
Nanotechnology 14, 385 (2003).
[7] C. F. Castro-Guerrero, F. L. Deepak, A.
Ponce,J.
Cruz-Reyes,
M.
D.
Valle-Granados, S. Fuentes-Moyado, D.
H. Galván and M. Jose-Yacamán, Catal.
Sci. Technol. 1, 1024 (2011).
[8] T. Matsumoto, Y. Souno, H. Tatsuoka, Y.
Nakanishi, H. Kuwabara, Appl. Surf. Sci.
169/170, 325 (2001).
[9] T. Hosono, M. Kuramoto, Y. Matsuzawa,
Y. Momose, Y. Maeda, T. Matsuyama, H.
Tatsuoka, Y. Fukuda, S. Hashimoto and
H. Kuwabara, Appl. Surf. Sci. 216, 620(2003).
[10] L. D. Hicks et al., Phys. Rev. B 47,
12727 (1993).
[11] L. D. Hicks et al., Phys. Res. B 47,
16631 (1993).
[12] R. Venkatasubramanian et al., Nature
413, 597 (2001).
[13] A. I. Hochbaum et al., Nature 451, 163(2008).

Effect of Ca addition on microstructure, mechanical properties anddegradation behavior of Mg–Zn–Sr alloy
Xiang Meng a, Hirokazu Tatsuoka b, *, Shaokang Guan caGraduate School of Science and Technology, Shizuoka University, 3-5-1 Johoku, Naka-ku,
Hamamatsu 432-8011, Japanb
Faculty of Engineering, Shizuoka University, 3-5-1 Johoku, Naka-ku, Hamamatsu, Shizuoka
432-8561, Japanc
Materials Research Center, School of Materials Science and Engineering, Zhengzhou
University, Wenhua Road 97, Zhengzhou, Henan 450002, China* E-mail: tehtats@ipc.shizuoka.ac.jp

Abstract
The effect of trace Ca addition on the microstructure, mechanical properties anddegradation behavior of the as-cast Mg-2wt. % Zn-0.3wt. % Sr alloys was investigated in thisstudy. The results indicated that there always were a decrease in grain size and an increase instrength with the increase of Ca content which were under 0.3wt%. Degradation was studied
by using immersion tests in Hanks’ solution. With more Sr addition the corrosion resistance ofalloys declined.

1. Introduction
Magnesium (Mg) alloys have gained
significant interests for their potential
applicationasbiodegradable
andbio-absorbable implant materials [1, 2].
Unfortunately, the high degradation rate of
Mg in physiological conditions is always asignificant challenge to its application.
Element alloying is one of the effective
measures to improve the mechanicalproperties and corrosion resistance of Mg
alloys. For preventing or minimizing
cytotoxicity or adverse tissue reactions,using of the alloying elements that exist in
the human body or have been shown to havebeneficial effects on tissue regeneration andhealing can be beneficial. Such elements,
like zinc (Zn), Calcium (Ca) and strontium(Sr) are non-toxic or less-toxic to human,
which were chosen for investigation in thisstudy. The content of Ca was changed inorder to study its effect on microstructure,mechanical properties and in-vitro corrosion
properties of Mg-Zn-Ca-Sr alloys.

2. Experimental
In the present study, Mg-2wt. %Zn- 0.3

wt. % Sr –x wt. % Ca (x=0, 0.1, 0.25, 0.5)
alloys were prepared in a mild steel crucible
under CO2/SF6 (volume fraction rate,
3000:1) atmosphere in an electronic
resistance furnace. The microstructure of thespecimens was observed by optical
microscopy (Olympus, H2-UMA) andscanning electron microscopy (Philips,
Quanta-2000). The tensile tests (69 mm ×
25 mm × 5 mm) were carried out on an
SHIMADZU （ AG-1C ） material testingmachine. The potentiodynamic polarization
experiments were conducted on anelectrochemical workstation (RST5200).

3. Results and discussion
3.1. Microstructure
Fig. 1 shows the optical microstructuresof the quaternary Mg–2wt. %Zn–0.3 wt. %Sr-x wt. % Ca (x = 0, 0.1, 0.25, 0.5) alloys.
With the increase of Ca content, the grain
size of alloys decreased gradually. This
mostly can be attributed to the growth
restriction by solute effect or by the
nucleating particles as fine precipitates wereobserved in the grains and along grainboundaries [3].

(a)

(b)

(c)

(d)

Fig. 1. Optical microstructures of alloys. (a) Mg-2Zn-0.3Sr alloy; (b) Mg-2Zn-0.3Sr-0.1Ca alloy; (c)Mg-2Zn-0.3Sr-0.25Ca alloy; (d) Mg-2Zn-0.3Ca-0.5Ca alloy.

3.2. Mechanical properties
Fig. 2 shows the mechanical propertiesof Mg-2Zn-0.3SrSr-x Ca alloys. It can beseen clearly that with the increasing Ca
content, the tensile strength increased firstly

and then decreased; Mg-2Zn-0.25Ca-0.3Sr
alloy exhibited the highest tensile strength
which was higher (24%) than Mg-2Zn-0.3Sr
alloy. When the content of Ca was over 0.3
wt. %, trend of tensile strength obviously
decreased.

190

3.4. Immersion test

180
170

b

σ /MPa

160
150
140
130
120
0.0

0.1

0.2

0.3

0.4

0.5

Ca content(wt.%)

Fig.

2.

Mechanical

Mg-Zn-Sr-x Ca alloys.

properties

ofFig. 3(a) illustrates the variation of pH
value with immersion time in SBF for 72 h.
More Ca addition brings more second phase,
which lets alloys generate more Mg2+ ionswhich increased the pH values [4]. Fig. 3(b)shows the mass loss curves of different
samples after immersion time in SBF
solution at 37
for 120 h. It can be seenthat the corrosion rate of the as-cast alloy
was similar with the variation of pH value.
Ca addition decreases the corrosion
resistance of alloys, but not obviously.

0.050

(a)

8.6

(b)

0.045
0.040

Weight loss/ mg/cm

28.4

pH Values

8.2

8.0

7.8

Mg-Zn-0.3Sr
Mg-Zn-0.1Ca-0.3Sr
Mg-Zn-0.25Ca-0.3Sr
Mg-Zn-0.5Ca-0.3Sr

7.6

0.035
0.030
0.025
0.020
0.015

Mg-Zn-0.3Sr
Mg-Zn-0.1Ca-0.3Sr
Mg-Zn-0.25Ca-0.3Sr
Mg-Zn-0.5Ca-0.3Sr

0.010
0.005

7.4

0.000
-0.005

7.2
-10

0

10

20

30

40

50

60

70

80

time/h

0

20

40

60

80

100

120

Time/hFig. 3. (a) The pH values and (b) the mass loss curves of different samples with immersion time in
SBF solution at 37 ℃.

4. ConclusionThe
microstructural,mechanicalproperties and degradation behavior of
Mg-Zn-Ca-Sr alloys were investigated. The
Mg-2.0wt%Zn-0.25wt%Ca-0.3wt%Sr
showed to be a promising biodegradable
alloy candidate due to their low degradation
rate and moderate mechanical properties. It
is important to control the microstructure for
optimization of both mechanical anddegradation behavior of the alloys.

Reference[1] Witte, F., Kaese, V., Haferkamp, H.,
Switzer, E., Meyer-Lindenberg, A.,
Wirth, C.J., Windhagen, H., 2005. In
vivo corrosion of four magnesium alloys
and the associated bone response.
Biomaterials 26, 3557-3563.
[2] Brar,
H.S.,
Keselowsky,
B.G.,
Sarntinoranont, M., Manuel, M.V., 2011.
Design considerations for developing
biodegradable magnesium implants.
JOM 63, 100-104.
[3] HUANG Xiao-feng, ZHU Kai, CAO
Xi-juan. The Roles of Alloying
Elements in Magnesium Alloys.
Foundry Technology, 2008, 29(11),1574-1578.
[4] Harpreet S. Brar, Joey Wong, Michele V.
Investigation of the mechanical anddegradation properties of Mg-Sr and
Mg-Zn-Sr alloys for use as potential
biodegradable implant materials. Journalof the Mechanical Behavior of
Biomedical Materials, 7(2012): 87-95.

Sulfur Concentration Dependence on Structural Property of MoSi2 Synthesized by
Molten Salt Technique using MoS2
Hajime Koyama
Faculty of Engineering, Shizuoka University
TEL&FAX: 053-478-1099
E-mail: f0911061@ipc.shizuoka.ac.jp

Abstract
The hexagonal phase MoSi2 was formed by the molten salt technique using MoS2
powder as the source material. And it is reported synthesis of the Hexagonal MoSi2
was promoted by adding S atom. In this paper, I investigated sulfur concentration
dependence on structural property of MoSi2 synthesized.
1. Introduction
Recently, semiconducting silicides haveattracted much attention for theirpotential to create new classes ofenvironmentally conscious electronics.
The group VI metal silicides areclassified in this material group [1].
MoSi2 with stable tetragonal (t-) phases
is metallic and has been used in electrodewire materials. The
t-MoSi2 has lattice constants of a=0.3205
nm and c=0.7845 nm. On the other hand,
the hexagonal (h-) phase MoSi2 haslattice constants of a=0.4622 nm and
c=0.6646 nm, and is known to be anarrow gap 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,
the electrical and thermoelectric propertyof the bulk crystal h-MoSi2 have notbeen experimentally clarified. Sincet-MoSi2 is transformed 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 grown using 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
the silicides. A nonhomogeneouscompositional distribution of Mo and Si
atoms was observed in the silicides whenusing the Mo substrate and Mo(Ti)compact [4]. The single phase growth of
h-MoSi2 and improved compositional

homogeneity of the silicide 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 andextraordinary chemical properties of thetwo-dimensional MoS2 nanoclusters aredescribed 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 thehomogeneity of the atomic distribution in
the resultant alloys, because the diffusing
species easily can pass through thematerials during the treatment. Moreover,it has also been reported that the
crystalline structure of the resultant
compounds is strongly affected by that of
the starting materials. For example, theformation of the zincblende-MnTe phase
was demonstrated using ZnTe [8]. Single
phase Ca2Si was obtained using Mg2Si
as the starting material [9]. In the samemanner, it is expected that the h-MoSi2
would be easily obtained using MoS2 as
the source material. Moreover, MoS2 haslayered structures. It is expected MoSi2
synthesized by Molten salt technique by
using MoS2 have nanostructures.
Nanostructures have unique properties
and novel applications compared to thoseof bulk materials. A decrease in thethermal conductivity and an increase in
the power factor, , of the nanostructures
have been demonstrated by thetheoretical calculations and proof ofprinciple experiments, relative to thecorresponding bulk structures.[10-13]
In this study, the synthesis of h-MoSi2
was examined using MoS2 powder.

using the appropriate starting material.

2.

[1] APAC-SILICIDE, July 29-31, 2006,
Kyoto, Thin Solid Films 515(22), pp.
8101-8300, edited by Y. Maeda, K.P.
Homewood, T. Sadoh, Y. Terai, K.
Yamaguchi and K. Akiyama.
[2] T. Nonomura, C. Wen, A. Kato, K.
Isobe, Y. Kubota, T. Nakamura, M.
Yamashita, Y. Hayakawa and H. Tatsuoka,
Physics Procedia 11, 110 (2011).
[3] A. J. Heron and G. B. Schaffer, Mater.
Sci. Eng. A352, 105 (2003).
[4] T. Nonomura, C. Wen, K. Shirai, K.
Isobe, A. Kato, Y. Kubota, T. Nakamura,
Y. Hayakawa and H. Tatsuoka, IOP Conf.
Ser.: Mater. Sci. Eng. 18, 14 (2010).
[5] L. S. Byskov, J. K. Norskov, B.S.
Clausen, and H. Topsoe, J. Catal. 187,
109 (1999).
[6] J. V. Lauriitsen, M. Nyberg, R.T. Vang,
M.V. Bollinger, B.S. Clausen, H. Topsoe,
K. W. Jacobsen, E. Laegsgaard, J. K.
Norskov and F. Besenbacher,
Nanotechnology 14, 385 (2003).
[7] C. F. Castro-Guerrero, F. L. Deepak,
A. Ponce, J. Cruz-Reyes, M. D.
Valle-Granados, S. Fuentes-Moyado, D.
H. Galván and M. Jose-Yacamán, Catal.
Sci. Technol. 1, 1024 (2011).
[8] T. Matsumoto, Y. Souno, H. Tatsuoka,
Y. Nakanishi, H. Kuwabara, Appl. Surf.
Sci. 169/170, 325 (2001).
[9] T. Hosono, M. Kuramoto, Y.
Matsuzawa, Y. Momose, Y. Maeda, T.
Matsuyama, H. Tatsuoka, Y. Fukuda, S.
Hashimoto and H. Kuwabara, Appl. Surf.
Sci. 216, 620 (2003).
[10] L. D. Hicks et al., Phys. Rev. B 47,
12727 (1993).
[11] L. D. Hicks et al., Phys. Res. B 47,
16631 (1993).
[12] R. Venkatasubramanian et al., Nature
413, 597 (2001).
[13] A. I. Hochbaum et al., Nature 451,
163 (2008)

Experiments

The Mo-silicide powder was synthesizedusing MoS2 powder, S powder and amolten 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 %, (5 mol% 10 mol% and 20
mol% for salt 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 773 and 873K.
The structural properties of the powder
were characterized by X-ray diffraction(XRD) .
3.

Results and Discussion

Fig.2 Intensity ratio of MoS2(002) andMoSi2(111) by XRD spectra of the resultant
powder by using S and MoS2 as the source
material at 500 ºC and 600 ºC.

4.
Conclusion
Synthesis of the h-MoSi2 was promoted
by adding S atom. In case of density ofS(10mol%), the effect was enhanced . It
is expected that this simple growth
procedure to fix the crystalline structure

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

Review on the fabrication techniques of SiGe- and Ge-on-insulator wafer
Veerappan Manimuthu1,2, Shoma Yoshida1,3, Yuhei Suzuki1,4, Hiroya Ikeda1,2,4
1.Research Institute of Electronics, Shizuoka University
2.Graduate school of Science and technology, Shizuoka University
3.Faculty of Engineering, Shizuoka University
4.Graduate school of Engineering, Shizuoka University*
e-mail: manimuthu@rie.shizuoka.ac.jp

Abstract
The nanostructured thermoelectric materials have been expected to enhance theperformance of various applications such as conversion of heat to electricity (i.e.) in the fieldof green energy technology, thermocooling and sensors. In order to fabricate SiGe and Ge
nanowires, in this paper we survey various approaches for preparing SiGe- and Ge-on-insulator(SGOI and GOI) wafers such as SIMOX (separated by implanted oxygen) process, Ge
condensation technique, direct wafer bonding process and also discusses the merits anddemerits to choose better fabrication technique.

1. Introduction
Waste heat is produced on a daily basis
in our life, for example laptop liberates
waste heat and in the car engine, more than
70% of the energy is wasted as heat. By
capturing some of these waste heat,
thermoelectric materials has the ability tocapture the waste heat and to charge thebattery if it is placed on car. So thethermoelectric materials are the possibilities
and endless.

Fig. 1 Importancematerials.

ofstill as waste heat. This energy loss or waste
heat from different sources can betransferred to electrical energy bythermoelectric (materials) system. The
thermoelectric effect is a phenomenon, in
which the temperature difference between
the hot and cold side causes the diffusion of
charge carriers in the materials and thus
produces electricity. Usual thermoelectric
system has a pair of n- and p-typesemiconductor materials (thermocouple) as
shown in Fig. 2 for obtaining a large
thermoelectromotive force.

thermoelectric

From the Fig. 1, only 34% of total heatfrom different sources such as thermal power
plant, nuclear power plant, automobiles,
factories etc., is useful and the other 66% is

Fig.

2

Thermoelectric effect principle.

Nanostructured thermoelectric materials
are expected for the enhancement of power
conversion efficiency because of following
factors,

1. Decrease in the thermal conductance byan increase in boundary scattering ofphonons
2. Enhancement of Seebeck coefficient
3. Realization of a large number density ofthermocouples by miniaturization
Another way for the improvement isusage of silicon germanium (SiGe) which
has emerged as an important semiconductor
technology. SiGe generally has lower
thermal conductance because of an increasein defect scattering of phonons. In order tomake SiGe and Ge nanowires for different
thermoelectric applications this paper
discusses about the different fabrication
techniques for preparing SGOI and GOI
wafers.

2. Fabrication processes
2.1 SIMOX process
Figure 3 shows the SIMOX process [1].
It is commonly known as separated by
implanted oxygen process, which is a mostpopular method for Si-on insulator (SOI)
fabrication. In this process there is an
implantation of O ions into the Si substrate,
after which ex-situ annealing at elevatedtemperatures removes implant damage andforms a uniform buried insulating SiO2 layer.
By optimizing fabrication conditions such as
implantation substrate temperature andpost-implantation annealing conditions, thecrystal defects and micro-roughness of thematerials can be reduced.

higher-Ge-content SiGe, the process does
not generate good quality structures for Geconcentrations greater than 30%. Surface
oxidation of SiGe occurs during post-implantannealing, and thermal instability of SiGe isobserved at elevated temperatures. SIMOX
for SGOI and GOI fabrication has notachieved a prominent role, and othertechniques have been pursued anddeveloped.

2.2 Ge Condensation ProcessThis technique begins with the formationof a SiGe layer with a low Ge fraction (0.15)
which was grown epitaxially by ultrahighvacuum chemical vapor deposition over acommercially available bonded SOI wafer.
Next, dry oxidation is carried out attemperatures lower than the melting
temperature of SiGe in O2 atmosphere. In
this oxidation process, Si atoms in the SiGe
layer are oxidized selectively and Ge atoms
are rejected from the SiO2 layers, while
out-diffusion of Ge atoms is suppressed by
the buried and top oxide layers, resulting in
the increase of Ge fraction in the SiGe layer
when the SGOI thickness becomes thinnerthan the initial SiGe thickness. Finally, theoxidation process is stopped when almost
all Si atoms in the SGOI layer are oxidized.
The schematic view of the process is shownin Fig. 4 [1, 2].

Fig.
4
Schematic
diagramGe-condensation technique.
Fig. 3 Schematic diagram of SIMOX
process.
The merits and demerits of SIMOX
process are described below. SiGe substrates
lead to acceptable buried oxide formation for
low Ge concentrations, including 10% and
18% Ge. Unfortunately, when applied toof

The merits and demerits of SIMOX
process are described below. SiGe layer withvery high Ge concentration can be produced.
Very thin and pure layer of SGOI/GOI
substrates can be produced. Growth of thick
Ge layer (needed for optoelectronics) isdifficult because long oxidation time isneeded.

2.3 SmartCut approach

Figure 5 shows the SmartCut approachwith stop layer modification [2]. The process
begins with a planarized Si1-xGex virtual
substrate with the desired Ge content,
followed by the growth of the stop andtransfer layers at low temperatures to avoid
interdiffusion. After hydrogen implantation,
bonding and splitting, the created structure
consists of a layer of remaining Si1- xGex
with the stop and transfer layers bonded tooxidized Si wafer. The remaining Si1- xGex isthen removed using any selective processthat does not attack the stop layer, leaving
behind a smooth surface. Optionally, thestop layer can then be removed leaving only
the transfer layers.
The merits and demerits of SIMOX
process are described below. The stop layer
method has great flexibility, and can beapplied to any number of transfer layers andstop layers. In addition, all layers are defined
epitaxially. Thicknesses of the transferredlayers are precisely controlled, allowing forthe fabrication of ultrathin layers on oxide.
This is ideal for fully-depleted MOSFETs.
	
  
	
  
	
  
	
  
	
  

Implanted	
  H+	
  Relaxed	
  SiGe	
  

	
  
	
  
Si	
  

Relaxed	
  SiGe	  The merits and demerits of SIMOX
process are described below. Direct wafer
bonding method has the ability to create
ultra-thin layers of SiGe- and Geon-insulator wafers. By using this process,one can control the impurity concentrationand the layer thickness precisely.

On comparing the above discussed
fabrication techniques, the direct wafer
bonding process seems useful to study thethermoelectric characteristics of SiGe andGe nanostructures and also for future highperformance CMOS applications because of
its simple procedure. Now we are confirming
the bonding conditions for GOI fabrication.

SiGe	
  
Si	
  substrate	  Relaxed	
  SiGe	  Implanted	
  H+	
  

Fig. 6 Schematic diagram of direct wafer
bonding process.

3. Conclusion

Relaxed	
  SiGe	
  

SiGe	
  

into contact by van der Waals attraction in aclean room at room temperature underatmospheric pressure. Ge side of the wafer isthinned by mechanical polishing. Finally, thebonded wafers are subjected to annealing forstrong bonding at the interface. The
schematic view of the process is shown in
Fig. 6 [1, 3].

Si	
  
SiGe	
  
Relaxed	
  SiGe	  Implanted	
  H+	
  

Relaxed	
  SiGe	  Si	  Oxide	  Si	  SiGeOI	  Oxide	
  Fig. 5 Schematic diagram of SmartCut
approach.

2.4 Direct wafer bonding process
In this fabrication method, at first the
SiO2 films are formed on the silicon
substrates by dry oxidation. The surfaces ofGe and the oxidized Si wafer are chemicallytreated (ammonium hydroxide and water).
The Ge and SiO2 surfaces were then brought

References[1] Gianni Taraschi *, Arthur J. Pitera,
EugeneA.
Fitzgerald,
Solid-StateElectronics 48 (2004) 1297–1305.
[2] Shu Nakaharai, Tsutomu Tezuka,
Naoharu Sugiyama, Yoshihiko Moriyama,
and Shin-ichi Takagi, Applied Physics
Letters 83, 3516 (2003).
[3] Gianni Taraschi,z Arthur J. Pitera, Lisa
M. McGill, Zhi-Yuan Cheng, Minjoo L. Lee,
Thomas A. Langdo, and Eugene A.
Fitzgerald, Journal of The Electrochemical
Society, 151 (1) G47-G56 (2004).

Advantage of KFM for Seebeck Coefficient
Measurement of Nanostructures
1,2

Takuro Oda 1,3,4Faiz Salleh

1,3,5

Hiroya Ikeda

1

Research Institute of Electronics , Shizuoka University
2
Faculty of Engineering , Shizuoka University
3
Graduate School of Science and Technology, Shizuoka Uviversity
4
Research Fellow of the Japan Society for the Promotion of Science5
Graduade School of Engineering , Shizuoka University

Abstract
With the aim of enhancing thermoelectric property, one of the methods is introduction ofnanostructures. Therefore, a new measurement method is desired for characterizing thethermoelectric properties of nanostructured materials. In the present study, we propose a newtechnique of Seebeck coefficient measurement of nanostructures by KFM (Kelvin-probe force
microscopy). This report demonstrates a principle of KFM and its advantage.

1.

Termoelectricity andnanostructures

Thermoelectric
conversionisthephenomenon to pick up heat from electricity or
to pick up electricity from heat. The efficiencyof thermoelectric power generation can be
expressed as a function of the dimensionlessfigure-of-merit, ZT, which is shown in Fig 1.
Z is the figure-of-merit and T is thetemperature. It is found from Fig.1 that
efficiency increases monotonously withincreasing the ZT. Therefore, the researchers
in this field have investigated to make Z
large.

Fig. 1: Relationship between ZT andefficiency

Z is usually defind as the following
equation,,where S is the Seebeck coefficient, σ is theelectrical conductivity and κ is the thermal
conductivity. Hence, the efficiency isenhanced by increasing Seebeck coefficientand by decreasing κ.
We focus our attention on silicon
nanowire to increase S and to decrease κ. S
is theoretically expected to increase when
the diameter of silicon nanowire is smaller
than 8nm [1]. This is due to the carrierconfinement effect. It can be expected thatwhen the diameter of silicon nanowire issmaller than the mean-free-path of phonon,

Fig. 2: Schematic diagram of boundary
scattering of phonon in silicon nanowire

the value of κ decreases by increasingboundary scattering of phonon, as
schematically shown in Fig.2.

Fig. 3. The Coulomb force is expressed as
  follows:

2. Charactarization of nanomaterials
Unfortunately, it is difficult to measure
the thermoelectric characteristics of silicon
nanowire because of its small dimensions.
Therefore, it is necessary to develop a newmeasurement method that can characteristic
nanostructures. In practice, there are mainly
two ways to measure the thermoelectriccharacteristics.
One is the fabrication of complicatedsample structure by MEMS. It means thatelectrodes and heaters are built in the sample[3-5]. This measurement method has merits
that the usual measurement devices areavailable and that the analysis of obtained
date is simple. However, the preparation of
the sample is very difficult.
The other is the usage of scanning probe
microscopy [2]. A probe touches the sample
surface, and thermoelectric characterizationis performed in the vertical direction.
Therefore, the contact between the metallic
probe and the sample surface, which disturbs
the thermal environment of sample, is not
avoidable
andspecially-customized
equipment is required.
We have developed a new measurement
technique using KFM (Kelvin-probe force
microscopy). The KFM gives us theinformation about the local surface potentialof the sample corresponding to the
Fermi-energy difference between the sample
and cantilever. KFM has merits that themeasurement is non-contact and easy andthat special sample structure is not required.
3. Measurement principle by KFMFirst of all, the measurement principle of
KFM itself is explained. Fig. 3 describes themeasurement principle of the Fermi energyof a sample by KFM. When a cantilever is inclose proximity to the sample surface and iselectrically connected to the sample, theFermi energy of the cantilever coincides
with that of the sample. Then, a Coulomb
force will occur between the cantilever andthe sample due to the work-function
difference qΔΦ, as shown in the left side of,where V is the potential difference between
the cantilever and the sample ( or ΔΦ ), C isthe cantilever-sample capacitor, and z is thecantilever-sample distance.
Note that KFM nullifies this force byapplying an external bias between the
cantilever and the sample ( VDC in the rightside of Fig. 3 ). The magnitude of this bias
corresponds to the potential difference.
Using the above procedure, the Fermi energyof the sample with respect to that of thecantilever can be obtained from the applied
bias.

Fig. 3: The schematic view of a cantilever andthe sample at the time of the KFM measurement
and the corresponding band diagram

The Seebeck coefficient is defined asfollows:,

where ΔV is the thermoelectromotive force,and ΔT is the temperature difference. S canbeobtainedbymeasuringthethermoelectromotiveforce for a given
temperature difference across the sample.
Equation of S is converted to,where TH and TL are the temperatures in thehigh-temperatureandlow-temperature
regions, respectively. ΦTH and ΦTL are thesurface potentials in the high-temperatureand low-temperature regions, respectively.
Therefore, S can be estimated from themeasurement of the Fermi energy for eachregion. Figure 4 shows the band structure of
the cantilever and n-type Si sample under atemperature difference. When a temperature
difference is applied to the sample, the Fermi
energy of the sample will be spatiallydistributed. Generally, EFTH is lower than
EFTL in n-type Si. Since KFM can measureΦTH and ΦTL as shown in Fig. 4, we canevaluate the Seebeck coefficient.
4. Future planNow, we are preparing a Si nanowire
sample and are scheduled to evaluate its
Seebeck coefficient by KFM.

Fig. 4: Band diagram of n-type silicon andcantilever under a temperature difference

Reference[1] N. Neophytou et al., Phys. Rev. B,
Vol.83, 245305 (2011)[2] Y. Zhang, C. L. Hapenciuc, et al., Appl.
Phys. Lett., vol.96, pp.62107, (2010)[3] F. Völklein, M. Schmitt, et al., J. Electro.
Mater. vol.38, pp.1109, (2009)[4] C.-H. Lee. G. -C. Yi, et al., Appl. Phys.
Lett., vol.94, pp.022106, (2009)[5] H. J. Ryu, Z. Aksamija, et al., Phys. Rev.
Lett., vol.15, pp.256601, (2010)

Fabrication of Si nanowire thermopile by focused ion beam
Yuhei Suzuki1,2*, Faiz Salleh1,3,4, Masaru Shimomura2,3, Akihiro Ishida2,3 and Hiroya Ikeda1,2,3
1Research Institute of Electronics, Shizuoka University,
Johoku 3-5-1, Naka-ku, Hamamatsu, 432-8011, Japan
2
Graduate School of Engineering, Shizuoka University,
Johoku 3-5-1, Naka-ku, Hamamatsu, 432-8561, Japan
3
Graduate School of Science and Technology, Shizuoka University,
Johoku 3-5-1, Naka-ku, Hamamatsu, 432-8011, Japan
4
Fellow of the Japan Society for the Promotion of Science,
Ichibancho 8, Chiyoda-ku, Tokyo, 102-8472, Japan
TEL&FAX: +81-53-478-1335*
e-mail: f0330181@ipc.shizuoka.ac.jp

Abstract
With the aim of fabricating a thermopile infrared photodetector using Si nanowire
thermocouples, we propose a method of alternate Ga-ion implantation by focused ion beam
(FIB) into a P-doped Si-on-insulator (SOI) nanowire array. In the present study,
rectangular-patterned n-type SOI layers with a P concentration of 5.1 × 1017 cm-3was prepared,
and then Ga ions were implanted into them with a dose of 4.0 × 1013 cm-2 and 3.5 × 1014 cm-2.
From the simulation of Ga ion implantation, the average Ga concentration of the Ga-implanted
SOI layers was estimated to be 3 × 1018 cm-3 and 3 × 1019 cm-3, respectively. The Seebeckcoefficient of the Ga-implanted SOI layer has an opposite sign as compared with that of theoriginal P-doped SOI layer. This result indicates that p-type Si was formed by the Ga
implantation. The relationship between the average Ga concentration and Seebeck coefficient
was consistent with a general characteristic of semiconductor materials.

1. Introduction
Thermoelectric semiconductors are oneof the candidates for an energy harvestingsince they have an ability to interconvert
between thermal energy and electricalenergy directly. They are usually used in theform of the thermopile that is an array of nand p-type semiconductor connected bymetal electrodes electrically in series but
thermally in parallel. A thermopile infrared(IR) photodetector has been studied as anapplication of a thermopile [1-3]. It is
composed of a thermopile equipped with anabsorption film at an end. When an infrared
ray is irradiated to a thermopile IR
photodetector, a temperature of a thermopile
on the part of an absorption film rises, and atemperature difference between the thermopile occurs. Thermopile IR photodetector
observe a thermoelectromotive force corresponding to the temperature difference. The
voltage sensitivity RV of a thermopile IR
photodetector is given by

RV =

nSη
G ,

(1)

where n is the number of thermocouples, S is
the Seebeck coefficient, η is the absorption
efficiency€ of the IR absorption film and G is
the thermal conductance between the highand low-temperature regions, which is
proportional to the thermal conductivity. It is
obvious from eq. (1) that a high-sensitive
thermopile IR photodetector requires athermopile material with a high Seebeckcoefficient and a low thermal conductivity.
Thermoelectric properties of semiconductor materials depend on its impurity
concentration. With increase in the impurity
concentration, the Seebeck coefficientdecreases and the thermal conductivity
increases. The Fermi level should becontrolled for high RV because Fermi level ofsemiconductor determines its impurity
concentration. In parallel, it is needed to enhance thermoelectric parameters themselves.

Introduction of nanowire thermopile toinfrared photodetector is expected toenhance RV through an increase in thenumber of thermocouples by miniaturization,
an increase in the Seebeck coefficient by acarrier confinement effect [4,5], and adecrease in the thermal conductivity byincrease in the rate of the boundary scattering of phonons [6,7]. These are why we haveinvestigated thermoelectric characteristics ofnanostructured Si [8-15]. Therefore, theconstruction of process techniques for a Si
nanowire thermopile is essential for practicaldevice fabrication. We propose a method for
preparing a nanowire thermopile in a
Si-on-insulator (SOI) wafer as shown in Fig.
1. First of all, the array of n-type Si
nanowires is patterned in an n-type SOI
wafer, in which one end of the wires isconnected by an n-type Si pad. Then, Ga
ions are alternately implanted into the n-type
Si nanowires by focused ion beam (FIB) toform p-type Si nanowires. FIB enables us toprecisely control the implantation to Si
nanowires with high reproducibility because
it has high spatial resolution and fineadjustment of the dose of implanted ions.
Finally, the n- and p-type nanowires areelectrically connected by depositing an Al
electrode on the Si pad.
In the present study, in order to makesure the forming p-type Si by FIB
implantation, we performed the implantationof Ga ions by FIB into P-doped n-type Si
wires. Then, the Seebeck coefficient wasmeasured for the original P-doped Si and theGa-implanted Si.

2. FIB conditions based on Ga
implantation simulation
The depth profile of implanted Ga ion
within the SOI layer was simulated todetermine the FIB implantation conditions
such as the dose of Ga ions and theacceleration voltage of Ga ions. Since
crystalline defects are produced in the SOI
layer during the ion implantation process, theSOI wafer has to be annealed after the Ga
implantation for recovering the crystallinity.
Hence, the distribution of thermally-diffused
Ga ions by post-annealing was also taken
into account.

Fig. 1 Fabrication procedure of a Si nanowire
thermopile by using FIB Ga ion implantation.

Fig. 2 Ga concentration profile in the SOI
layer after annealing.

The simulation of implanting Ga ion
accelerated through 30 kV into the SOI layer
with a thickness of 85 nm was performed fora dose of 4.0 × 1013 cm-2 and 3.5 × 1014 cm-2.
Then, the thermal diffusion of Ga ions by theannealing at 830 oC for 20 minutes wascalculated. The final Ga concentration
profile in the SOI layer is shown in Fig. 2.
The original P concentration NP, which was
determined to be 5.1× 1017 cm-3, is also
shown in the figure. Because it is found from
Fig. 2 that the Ga concentration is larger than
the P concentration at the whole position of
the SOI layer, it is predicted that p-type SOI
layers are formed without p-n junctions forthese implantation conditions. In addition,
the average Ga concentration in SOI layers
was estimated from the Ga concentration
profile, which is 3× 1018 cm-3 and 3 × 1019
cm-3, respectively. The dose of Ga ion andthe average Ga concentration NGa in the SOI
layer used in this study are summarized inTable 1.

Table 1 P concentration NP, Ga dose and average Ga concentration NGa for SOI
samples prepared in this study. Measured Seebeck coefficients are also listed.

NP [cm-3]n-typep1
P2

5.1 × 1017

5.1 × 1017

5.1 × 10

17

Dose [cm-2]

NGa [cm-3]

S [mV/K]---0.357

4.0 × 1013

3.5 × 1014

3 × 1018

0.331

3 × 1019

0.245

3. Sample preparation andmeasurement procedure
SOI wafer with the area of 1 × 1 cm2 wascut, and the thickness of SOI layer was 340
nm. SOI layer was thinned to a thickness of
85 nm by repeating thermal oxidation and
HF etching, and then P atoms were thermally
doped into the SOI layer at 830 oC for 20
minutes. The P concentration in the SOI
layer was estimated to be 5.1 × 1017 cm-3
from four-probe resistance measurement atroom temperature. A number of n-type Si
wires with a width of 500 nm and a length of
1 mm were patterned by SiO2 mask, electron
beam lithography and KOH etching. Si pads
with an area of 1.2 × 0.5 mm2 were also patterned at the both ends of Si wires for TEMF
measurement. Ga ions were implanted
according to the FIB conditions determinedin Sec. 2 with the implantation angle of 7 o
for avoiding a channeling effect [16]. Al
electrodes with an area of 0.4 × 0.4 mm2
were deposited on the Si pads. The thicknessof buried oxide layer was 400 nm and thep-type Si substrate had an impurity
concentration of 1015 cm-3.
Seebeck coefficient was measured using
the same method as that used in our previous
investigations [8,9,13]. The measurement
setup is depicted in Fig. 3. The two gilded
Cu plates were arranged with a gap of 1 mm.
The SOI sample was bridged on the gap. A
resistive heater was attached to one of a pairof Cu plates. Therefore, by heating one sideof the sample, the temperature difference
between both ends of the wire was produced,
and the TEMF corresponding to thetemperature difference was generated. The
time evolutions of the TEMF and thetemperature of the SOI surface at the lowtemperature side TL and the high temperatureside TH were simultaneously measured. Fromthese data, we can obtain the Seebeck coefficient as a function of average temperature.

Fig. 3 Schematic diagram ofcoefficient measurement system.

SeebeckFig. 4 Measured TEMF of original P-doped andGa-implanted SOI layers as a function oftemperature difference.

During the TEMF and temperature
measurement, the Seebeck coefficient wasalmost constant in an average temperature
range from 340 K to 380 K. Therefore, theaverage Seebeck coefficient S was evaluatedfrom the gradient of relationship between the
TEMF ΔV and the temperature difference
ΔT=TH-TL as given by S=-ΔV/ΔT.

4. Result and discussion
Figure 4 shows the TEMF ΔV measured
as a function of temperature difference ΔT
for the original n-type and implanted SOI
samples. It is clearly found that the TEMF ofthe implanted SOIs shows negative values
while the n-type SOI has positive TEMF.
This sign inversion indicates that p-type Si
was successfully formed by the Ga
implantation. The relationship between the
TEMF and the temperature difference can be
fitted by a straight line, which evidences that
the SOI layer has a constant Seebeckcoefficient in the measurement temperature
range. The Seebeck coefficients for original
P-doped and Ga implanted SOI layer are
listed in Table 1. The Seebeck coefficient of
the p2 sample with a larger dose is found tobe smaller than that of the p1 sample. In
general, absolute value of Seebeck coefficient in semiconductor materials decreases
with increase in its carrier concentration[17-19]. Hence, the relationship between themeasured Seebeck coefficients of implanted
SOI layers and the dose amount is consistent
with the general property, and we cancontrol the Seebeck coefficient of p-type Si
nanowire through the dose of implanted Ga
ions.

5. Conclusion
We proposed and investigated a methodof alternately implanting Ga ion by FIB into
the array of P-doped Si wires for fabricatinga Si nanowire thermopile. The measured
TEMF of Ga-implanted SOI layer shows an
opposite sign as compared with the TEMF oforiginal P-doped SOI layer. From this result,
it is found that the SOI layer changed fromn-type to p-type by Ga implantation into
P-doped SOI layer. The measured Seebeckcoefficient of highly Ga-implanted SOI layer
was smaller than that of SOI layer with alower dose. This result is consistent with thetendency of characteristics in generalsemiconductor materials. Consequently,
Ga-ion implantation by FIB can be applied
to fabrication of a Si nanowire thermopile.
Acknowledgment
This work was financially supported by aGrant-in-Aid for Challenging Exploratory
Research (No. 24651168) and a Grant-in-Aid
for JSPS Fellows (No. 23·6002) from the
JSPS.

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and Actuators A: Phys. 46 (1995) 422[3]M.
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New Measurement Method of Thermal Conductivity for Si Nanowire
1,2Shoma Yoshida, 1,3Yuhei Suzuki, 1,4Veerappan Manimuth,
3,4 Masaru Simomura, 3,4Kenji Murakami, 1,3,4 Hiroya Ikeya,

1Research Institute of Electronics, Shizuoka University
2Faculty of Engineering, Shizuoka University
3Graduate School of Engineering, Shizuoka University
4Graduate School of Science and Technology, Shizuoka University

Abstract
This paper introduces a new measurement method of heat transport characteristics
for Si nanowire. In this method, we can measure the thermal diffusivity without
touching and destroying the sample. Therefore we can expect diversification of
measurement environment as well as easy measurement procedure and highthroughput.

1. Background
Thermoelectric materials can interconvert between heat and electric power andhave some benefits which are no movableparts, maintenance free, miniaturization andenvironment-friendly energy conversion.
However there are some problems forpractical use. They are insufficientefficiency and low thermoelectromotive
force. Therefore, we want to enhance
efficiency of thermoelectric conversion.
Figure 1 shows relationship between
the power generation efficiency and ZT.
The figure-of-merit Z is expressed as
Z = S2 / 

free path is average distance between
phonon scatterings. Specifically, ph isthe parameter characterizing phonon
scatterings. Thermal conductivity decreases
when the diameter of wire is smaller thanphonon mean free path. Figure 2 shows thecomparison between two wires withlarge and small diameter. ph is thephonon mean free path. If nanowire
diameter is smaller than ph, theprobability of phonon collision at
nanowire wall increases.

(1)

where S is the seebeck coefficient,  is theelectric conductivity and is the thermal
conductivity. From Fig. 1, we have toincrease Z to enhance the efficiency.
Equation 1 shows Z is inversely
proportional to thermal conductivity.
Therefore, one of the enhancement of theefficiency is reduction of thermal
conductivity.
Heat
transport
phenomenon is conducted by lattice
vibration. From the point of view ofquantum mechanics, it is regarded as theparticle called for phonon. Phonon mean

Fig.1 Relationship between the power
generation efficiency and ZT

Fig. 2 Relationship between the wire diameterand phonon mean free path
Therefore, phonon momentum or velocity
gets lower. This leads to the decrease inthermal conductivity.
For this purpose, we use nanostructured materials such as nanowire.

temperature by lock-in amplifier withrespect to the reference AC source.

Light irradiation

2. Measurement Method
Measurement of thermal conductivity
for nanowires is very difficult. We havean idea about a new method based on theprinciple of AC calorimetry method tomeasure thermal diffusivity [1]. Figure 3
is a schematic diagram of AC calorimetry.
Laser is irradiated to the surface ofsample to provide heat flux and measure
the time delivering heat to thermocouple(TC). Then, we measure the change ofheat travelling time(t) with changing L.
Finally, we evaluate thermal diffusivity
from the relationship between L and t.
First of all, we are constructing ameasurement
systemofthermaldiffusivity for bulk materials as shown in
Fig. 4. In this system, we put anoscillation signal into halogen lamp fromfunction generator, by which the sample
is periodically heated. The heatpropagates through the sample and we
detect temperature by pyrometer. At thattime, we observe the delay time of ACTC
SAMPLE

Fig. 3

3.

L
Schematic diagram of AC
calorimetry method.

Future Plan

Now, we are constructing the systemon measurement using a lock-in
amplifier for detecting t and its
analysis procedure. In the next step, we
will observe the thermal diffusivity of Si
wafer and SiO2 layer as a reference.

Vacuum Chamber

Z-Direction

SampleX-Direction

Y-Direction

Halogen Lamp

Optical Fiber
Pyrometer

Lock-in Amplifier

Function Generator

AC Source
AC Temperature

Fig. 4

Schematic diagram of measurement system

Reference[1]I. Hatta, Y. Sasuga, R. Kato and A.
Maezono, Rev. Sci. Instrum. 56(1985) 1643

1

Improvement of DSSC through surface
modification
Rangga Winantyo1 and Kenji Murakami2
1

Graduate School of Science and Technology, Shizuoka University, Japan
2

Graduate School of Engineering, Shizuoka University, Japan

Abstract
One dimensional nano-wire ZnO has been studied to enhance electrical properties of Dye
Sensitized Solar Cells. Different ZnO nano-structures were synthesized by using different ZnO
salts through water bath deposition. After calcination, nano-structures on FTO surface wereobserved. The nano-structure profile was evaluated to see the effect on cell’s electricalproperties. Better aligned nano-structure was formed by using zinc acetate as precursor solution.

2. Experimental

1. Introduction
DSSC is one potential candidate fornext-generation solar cells, because it does
not require a material with high purity so thatthe cost of the production process isrelatively low.

a. Preparation of ZnO thin film
The ZnO thin films were deposited on the
FTO

coatedglass

substrate

by

using

hydrothermal method. All the substrates werecleaned with ethanol and acetone using

DSSC’s performance can be improved byultrasonic cleaner. The ZnO nanostructured

decreasing the electron-hole recombination

thin film was grown through hydrothermal

rate and by increasing the light-harvesting

method by immersing the substrate in acapability. ZnO is a semiconductor material

mixturewith a band gap around 3.37 eV at room

(CH3COO)2Zn.2H2O,

zinc

temperature, which is similar to TiO2

hexahydrateand(3.2eV)[1]. In the present study, we arehexamethylene tetramine (CH2)6N4

investigating the method to enhance the cell’sfor 2 hours.

of

5

mM

Zn(NO3)2

zinc

acetatenitrate5mM
at 900C

performance through surface modification.
b. DSC preparation
Platinum (Pt) coated glass was used as acounter electrode. The ZnO thin film wasdipped into Ruthenium dye (N719) solution

2

for 24 hours. The DMPII electrolyte wasused in this experiment.
c. Characterization method
Bunkoh Keiki-JASCO solar simulator
was used to measure the solar cell efficiency.

Fig.2 Nanostructure of DSC based on zinc nitrate(left) and zinc acetate( right).

The measurement was done under 1.5AM.
4. Conclusion

3. Result and Discussion
A well-aligned ZnO nanowire array was
From the measurement result in Fig.1, it

synthesizedfor

improving

thecell’scan be seen that the efficiency is slightly

efficiency.

Our

preliminary

conclusion

increased.

suggests that the efficiency can be increased
by changing the form of the nanostructure.

References[1]

Phattalung, Sutassana, Na., Smith,Michael, F., and Limpijumnong, Sukit,“First Principles Study of NativeDefects

Fig.1 I-V measurement of ZnO DSC based on

in Anatase TiO2,” Suranaree University ofzinc acetate and based on zinc nitrate.

Technology (SUT), 2006

By using the zinc acetate to growvertically aligned nanowire, the efficiency ofthe cell raises from 1.9% to 2.4%. From Fig.2,
we also can see that there are changes in themorphology of the nanostructure.

Crystal Structures of K2PtCl6 and related materials
Yuya Shirahashi
Graduate School of Engineering, Shizuoka University,
3-5-1 Johoku, Naka-ku, Hamamatsu 432-8561, Japan*e-mail: f0330175@ipc.shizuoka.ac.jp

Abstract
Crystal structure of K2PtCl6 and related materials were analyzed. They are having a crystal
structure similar, but they have different characteristics. The structural properties of the resultant
composites were characterized by X-ray diffraction (XRD).

1. Introduction
Platinum and its compounds has been widelyused because platinum has characteristics ofchemically very stable and hardly-oxidized.
For example, platinum has been applied tofuel cell and electrode, and medical, catalyst
purification of exhaust gas of automobile
because of platinum has the high melting pointand catalytic activities, anti-microbialactivities. [1]However, the structure of the platinumcompounds isn’t known in many cases.
Therefore, knowing the structure of theplatinum compound is important in order toexpand the range of applications.
This paper shows growth mechanism anddetailed structure of K2PtCl6 and relatedmaterials.

consist of Cs2PtI6. This compounds consists ofa single-phase Cs2PtI6.
Figure2 shows XRD spectra of the resultant
powder by using K2PtCl6 as the source
material. Figure2 shows that the compounds
consists of three materials. The peak of
K2PtCl6 is caused by residual source material
because K2PtCl6 is not soluble. The peak of
Cs2PtI6 appeared by the synthesis of K2PtCl6
and Cs, I2. The peak of Cs2PtCl6 appeared as
an intermediate product because reaction was
not sufficient. The amount of the resulting
compounds is more in the order of K2PtCl6,
Cs2PtI6, and Cs2PtCl6 from the peak intensity.

2. Experiments

The Peak appears in the order of Cs2PtI6,

Platinum compounds has been synthesized
as a source material K2PtCl6 and H2PtCl6. First,
Platinum compounds were synthesized bysource material, CsCl and I/KI. Second,
platinum compounds were obtained by drying
the solution. The structural properties of thecompounds were characterized by X-ray
diffraction (XRD).

Cs2PtCl6 and K2PtCl6 in the same Miller
indices from the difference in lattice constant.

3. Results and Discussion

Figure1 shows XRD spectra of the resultant
powder by using H2PtCl6 as the source
material. Figure1 shows that the compounds

Figure1. XRD spectra of the resultant
powder by using H2PtCl6 as the sourcematerial.

Figure2. XRD spectra of the resultant powder
by using H2PtCl6 as the source material.

4. Conclusion

Cs2PtI6 was synthesized by source material(K2PtCl6 or H2PtCl6) and CsCl, I2/KI. As aresults, the crystal structure of K2PtCl6 andrelated materials can be estimated.
In future, the crystal structure of anotherrelated materials of K2PtCl6 will be analyzed.

5. Reference[1]Kyung-Hwan Cho, Jong-Eun Park, Tetsuya

Osaka, Soo-Gil Park, Volume 51, Issue 5, 10
November 2005, Pages 956–960