The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Organizing Committee
Symposium Chair
Prof. Hideshi Miura (Kyushu University, JAPAN)
General Chair
Prof. Masaaki Otsu (University of Fukui, JAPAN)
Dr. Gandjar Kiswanto (University of Indonesia, INDONESIA)
Co-General Chair
Prof. Ming Yang (Tokyo Metropolitan University, JAPAN)
Dr. Ario Sunar Baskoro (University of Indonesia, INDONESIA)International Advisory Committee
Prof. Naoto Ohtake (Tokyo Institute of Technology, JAPAN)
Prof. Hiroshi Asanuma (Chiba University, JAPAN)
Dr. Gandjar Kiswanto (University of Indonesia, INDONESIA)
Prof. Ahmad Fauzi Mohd Noor (USM, MALAYSIA)
Dr. Zainuddin Sajuri (UKM, MALAYSIA)
Prof. Chaosuan Kanchanomai (Thammasat University, THAILAND)
Dr. Zainuddin Sajuri (Universiti Kebangsaan Malaysia, MALAYSIA)
Prof. Syoni Soepriyanto (Institute of Technology Bandung, INDONESIA)
Scientific & Program Committee
Prof. Yukio Miyashita (Nagaoka University of Technology, JAPAN)(Chair)Dr. Sugeng Supriadi (University of Indonesia, INDONESIA)(Chair)
Dr. Yudan Whulanza (University of Indonesia, INDONESIA)
Dr. Jos Istiyanto (University of Indonesia, INDONESIA)
Prof. Nobuo Takeda (University of Tokyo, JAPAN)
Prof. Kenji Kaneko (Tokyo University of Science, JAPAN)
Prof. Masaaki Itabashi (Tokyo University of Science, Suwa, JAPAN)
Prof. Hideki Kyogoku (Kinki University, JAPAN)
Prof. Ikuo Ihara (Nagaoka University of Technology, JAPAN)
Prof. Masanobu Kubota (Kyusyu University, JAPAN)
Prof. Hiroyuki Kawada (Waseda Univeristy, JAPAN)
Prof. Toshio Hattori (Shizuoka Institute of Science and Technology, JAPAN)
Prof. Hiroshi Asanuma (Chiba University, JAPAN)
Prof. Koji Fujimoto (The University of Tokyo, JAPAN)
Dr. Tsutomu Murai (Japan Science and Technology Agency, JAPAN)
Prof. Masahiro Fukumoto (Toyohashi University of Technology, JAPAN)
Prof. Hironori Tohmyo (Tohoku Univeristy, JAPAN)
Prof. Shinji Ogiwara (Tokyo University of Science, JAPAN)
Prof. Satoshi Kobayashi (Tokyo Metropolitan University, JAPAN)
Prof. Shuichi Wakayama (Tokyo Metropolitan University, JAPAN)
Prof. Hiroki Akasaka (Tokyo Institute of Technology, JAPAN)
Prof. Ryo Matsumoto (Osaka University, JAPAN)
Prof. Hidemitsu Furukawa (Yamagata University, JAPAN)
Prof. Mikio Muraoka (Akita University, JAPAN)
Prof. Ming Yang (Tokyo Metropolitan University, JAPAN)
Dr. Miki Yamazaki (Hitachi, JAPAN)
Prof. Wataru Nakao (Yokohama National University, JAPAN)
Symposium Secretary
Prof. Masato Okada (University of Fukui, JAPAN)
Ms. Herra Asta (University of Indonesia, INDONESIA)

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ACTIVITY LOCATOR
AND FLOOR PLANS

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ACTIVITY LOCATOR
AND FLOOR PLANS

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ACTIVITY LOCATOR
AND FLOOR PLANSv

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ACTIVITY LOCATOR
AND FLOOR PLANS

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

PROGRAMAT GLANCE
Date

Time

10 August2015

04.0006.00 pm

11 August2015

07.3008.00 am
08.0008.40 am
08.4009.00 am
09.0009.30 am
09.3010.30 am
10.3010.45 am
10.4512.00 am

13 August
2015

Venue

Registration and Welcome Drink

Pre-function Hall

Registration

Pre-function Hall

Opening Ceremony
Photo Session
Keynote Speech 1

Rinjani Room
I, II, III

Talk show: Serve the Country
Coffee break
Keynote Speech 2 and 3

12.0001.00 pm

Lunch
Poster Session
Exhibition

Restaurant

01.0003.00 pm

Parallel session

Meeting Rooms

03.0003.30 pm

Coffee Break
Poster Session
Exhibition

Pre-function Hall

Parallel session

Meeting Rooms

Poster Session
Exhibition

Pre-function Hall

Banquette Dinner

Rinjani Room
I, II, III

Parallel session

Meeting Rooms

10.0010.30 am

Coffee Break
Poster Session
Exhibition

Pre-function Hall

10.3012.00 am

Parallel session

Meeting Rooms

12.0001.00 pm

Lunch
Poster Session
Exhibition

Restaurant

01.0003.00 pm

Parallel session

Meeting Rooms

03.0003.30 pm

Coffee Break
Poster Session
Exhibition

Pre-function Hall

Parallel session

Meeting Rooms

Closing Ceremony

Selaparang Room

03.3005.00 pm
05.0007.00 pm
07.0009.00 pm
08.0010.00 am

12 August2015

Program

03.3005.00 pm
05.00 06.00 pm
08.00 am08.00 pm

Pre-function Hall

Pre-function Hall

Social Tour Lombok

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Table of Contents
IN-SITU NONINVASIVE HEAT FLUX EVALUATION BY ULTRASOUND FOR HIGHTEMPERATURE MATERIAL PROCESS MONITORING

1

ELASTIC-PLASTIC FRACTURE TOUGHNESS OF EXTRUDED Mg-Al-Zn
ALLOYS

3

FABRICATION AND EVALUATION OFMICRO-STRUCTURED REACTION FIELD
WITH VERTICALLYALIGNED CARBON NANOTUBES FOR MICRO BIOANALYSIS DEVICE

5

IMPROVEMENT OF MECHANICAL PROPERTIES IN UNTWISTED CNT YARNBY PAA/DMSO DENSIFICATION TREATMENT

7

EFFECTS OF PROCESSING CONDITIONS IN UNIDIRECTIONAL CARBONFIBER THERMOPLASTIC PREPREG TAPE LAYING USING NEAR INFRAREDHEATER

9

ESTIMATION OF ELECTRICAL CONDUCTIVITY OF TIB2/AL COMPOSITESWITH IMAGE ANALYSIS

11

NEW MANUFACTUTING METHOD OF PARTICLE INTERMETALLICCOMPOUND REINFORCED COMPSOITES

13

THE EFFECT OF ZrO2 ADDITION ON THE ELECTROCHEMICAL BEHAVIOR
OF LOW-CARBON STEELS IN CrO3–H2SO4 ELECTROLYTE

15

PRODUCTION AND IMPROVEMENT OF CHARACTERISTICS OF SINTEREDPOROUS TITANIUM USING CAPSULE-FREE HIP METHOD17PROPOSAL FOR NEW FABRICATION METHOD OF GREEN COMPOSITE BY
USING FRICTION STIR PROCESS19

INFLUENCE OF GRAIN SIZE ON PROCESS AFFECTED ZONE IN MICROPIERCING

21

REPLICATION OFANTIREFLECTION STRUCTURE FOR SOLAR CELLS BY
LASER-ASSISTED IMPRINTING

23

SELECTIVE PATTERN-HARDENING OF AISI-SUS420/SUS430 MOLDS FORMICRO-TEXTURING

25

EVALUATION OF MAGNETIZATION FOR CANTILEVER SAMPLE USINGMEMS DEVICE27

FABRICATION OF MICRO-CHANNEL FROM THE POWDER BY PULSEELECTRIC CURRENT SINTERING METHOD AND MARKED GRAFITE DIE

29

NUMERICAL STUDY ON EFFECTS OF MATERIAL INHOMOGENEITY AND
GRAIN SIZE ON NECKING AND WRINKLING BEHAVIORS IN MICRO DEEPDRAWING OF METAL FOILS31

TECHNICAL MILESTONE TO APPLY SELF-HEALING CERAMICS

33

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BLANKING VIBRATION AND NOISE OF DIFFERENT TENSILE STRENGTH
MATERIALS IN SCREW DRIVE SERVO PRESS BLANKING

35

BEHAVIOR OF ASYMMETRIC DEFORMATION OF ROD IN LOCALLYLATERAL UPSETTING

37

DIRECT MEASUREMENT OF FRICTION FORCE DURING FORMING

39MEASUREMENT OF TOOL TEMPERATURE IN BURNISHING USINGDIAMOND TIP

41

HIGH DENSITY OXYGEN PLASMA ASHING OF CVD-DIAMOND COATED
END-MILLING TOOLS43

FRONT AND REAR SIDES SIMULTANEOUS FORMING BY FRICTION STIR
INCREMENTAL FORMING

45

IN SITU OBSERVATIONS ON FRETTING WEAR BEHAVIOR OF PLASMASPRAYED HYDROXYAPATITE COATING47

WETTABILITY BETWEEN GRAPHITE AND HIGH- TEMPERATURE LEADFREE SOLDER ALLOY

49

THE ROLE OF HYDROGEN IN HIGH RATE PLASMA NITRIDING OF
MARTENSITE STAINLESS STEEL51

DEVELOPMENT OF TITANIUM FIBER/ALUMINUM MULTIFUNCTIONALCOMPOSITES

53

PHASE-FILED SIMULATION OF ELEMENTARY PROCESS IN POWDER BED
FUSION55

ESTIMATION OF FRICTIONAL PERFORMANCE OF LUBRICANTS FOR HOTFORGING OF STEEL BY CARIBRATION CURVES CONSIDERING HEATTRANSFER IN RING COMPRRESSION TEST57

DEVELOPMENT OF HIGH SPEED TORSIONAL FATIGUE TESTING MACHINEFOR MULTIPLE SPECIMENSDEVELOPMENT OF HIGH SPEED TORSIONALFATIGUE TESTING MACHINE FOR MULTIPLE SPECIMENS

59

GUIDED WAVE PROPAGATION IN WIRE CABLE UNDER TENSILE STRESS61

BEHAVIOR ON GENERATION AND PROPAGATION OF FATIGUE CRACK
FOR WC-Co CEMENTED CARBIDES63

A DEVICE OF GAS TRANSMISSION RATE FOR THIN FILMS IN A HYDROGEN
STORAGE SYSTEM

65

THE STUDY OF VARIANT SELECTION FOR SHAPE MEMORY EFFECT ON
FE-MN-SI BASED ALLOY WITH EBSD

67

CHARACTERISTICS OF METAL MATRIX COMPOSITES REINFORCED WITH
PALM OIL FLY ASH FABRICATED BY USING STIR CASTING AND
CENTRIFUGAL CASTING METHODS

69
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MECHANICAL PROPERTIES OF AMORPHOUS CARBON FILMS DEPOSITEDFROM CARBON ISOTOPE MATERIALS71

EVALUATION OF ADHESIVE STRENGTH BETWEEN Ti–29Nb–13Ta–4.6Zr
ALLOY AND SOL-GEL FABRICATED HYDROXYAPATITE FILMS

73

ADHESION BEHAVIOR OF MECHANICALLY JOINED PARTS OF ALUMINUM
ALLOYS75

DEVELOPMENT FOR HYDROGEN STORAGE Mg-Ti SOLID SOLUTION BY
MECHANICAL ALLOYING77

DEVELOPMENT OF DAMAGE DETECTION TECHNIQUE FOR SHEET-TYPE
LITHIUM-ION BATTERIES USING ACOUSTIC EMISSIONS

79

IMPACT RESISTANCE OF THE GLASSES INSTALLED NANO-LAMIRATEDFILMS AND ARAMID FIBER CLOTH

81

INFLUENCE OF ANNEALING ON SHAPE MEMORY PROPERTY IN FE-MN-SI
BASED ALLOY

83

FRICTION STIR INCREMENTAL FORMING OF ALUMINUM ALLOY SHEETSWITH BACK SUPPORT DIE

85

DEVELOPMENT OF CU-SN ALLOY WITH MECHANICAL ALLOYING87

FEM SIMULATION FOR TWO-DIMENSION CUTTING OF TITANIUM-ALLOY

89

CORROSION RESISTANCE OF Si-DOPED AMORPHOUS CARBON FILMS
AGAINST NITRIC ACID91

TENSILE AND FATIGUE BEHAVIORS FOR LIP–SEAL WELDS OF LIQUID
LITHIUM PIPES AT ELEVATED TEMPERATURE

93

EFFECT OF LOCAL PLASTIC PROCESSING ON FATIGUE STRENGTH OFSTEEL PLATE WITH A HOLE

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

IN-SITU NONINVASIVE HEAT FLUX EVALUATION BY ULTRASOUND FOR
HIGH TEMPERATURE MATERIAL PROCESS MONITORING
Shingo Isobe1, Ikuo Ihara2*, and Akira kosugi1
1
2

Graduate Student, Nagaoka University of Technology, Nagaoka, 940-2188, Japan
Dept. Mech. Eng., Nagaoka University of Technology, Nagaoka, 940-2188, Japan* ihara@mech.nagaokaut.ac.jp

Keywords: Nondestructive measurement, Heat flux, Process monitoring, Ultrasound, Hightemperature
In materials processes at high temperature, it is often required to estimate temperature and/or
heat flux of the processed materials and processing machines because such thermal parameters areclosely related to the quality and reliability of final products. Such parameters are also indispensable tounderstand physical phenomena related to the thermal conduction and thermal resistance of heatedmaterials. In addition, quantitative and real-time information on such parameters during materials
processing are required to make accurate theoretical or computational analyses for material processes athigh temperature. It is, therefore, strongly required to conduct in-situ measurements of temperature andheat flux for heated materials during high temperature processing. In particular, in-situ noninvasive
monitoring of heat flux through the interface between two materials is quite attractive and could be
beneficial for basic and practical researches of high temperature materials processing.
Ultrasound, due to its high sensitivity to temperature, has the potential to be an effective meansfor estimating thermal parameters of materials. Because of the advantages of ultrasonic measurements
such as noninvasive and faster time response, ultrasound is expected to be a promising tool formeasuring heat flux as well as temperature of heated materials. In this work, an effective ultrasonicthermometry providing simultaneous measurements of temperature and heat flux in a heated material isdeveloped and its feasibility of conducting in-situ noninvasive monitoring of heat flux through theinterface between two materials is examined through experiments with a single-side heated steel plate.
Temperature and heat flux determination by ultrasoundConsidering that a material whose single side is heated by contacting with a hot material andassuming a one-dimensional temperature distribution in the single-side heated material, the transit timeof ultrasound in the direction of the temperature distribution can be given by
L
1
tL 
dx
(1)
0 vT  x where L is the propagation distance of ultrasound, v(T(x)) is the ultrasonic wave velocity which is afunction of the temperature T at a location x. Based on equation (1), the temperature distribution T(x)near the heated surface can quantitatively be determined by the ultrasonic thermometry we had
developed[1]-[3]. This method basically consists of ultrasonic pulse echo measurements and a finite
difference analysis. It is noted that the temperature at the heated surface can be obtained by the method
without any boundary conditions at the heating surface.
Heat flux q, that is a heat transfer rate per unit area in the x direction, is generally expressedbyq  kTx(2)

where k is the thermal conductivity and T x is temperature gradient. If the thermal conductivity is
known, the heat flux can be determined from the temperature gradient near the heating surface. Since
the temperature gradient can be obtained from the temperature distribution determined by the ultrasonicthermometry, the heat flux at the heating surface is easily determined from the equation (2).
Experiment and result
Figure 1 shows a schematic of the experimental setup used. A steel plate of 30 mm thickness
is used as the specimen and its bottom surface is heated by contacting with a copper plate of 30 mm
thickness whose bottom is heated by a gas burner. To obtain a reference vale of heat flux at the bottom

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

surface of the steel, a conventional thin film heat flux sensor is inserted into the boundary between thesteel and copper plates. A longitudinal ultrasonic transducer of 5 MHz is installed on the top surface ofthe steel and ultrasonic pulse-echo measurements are made during heating. The transit times of theultrasonic pulse echoes from the steel plate are precisely determined by taking a cross correlation of thepulse echoes and used for estimating the temperature distribution in the steel. Once the temperature
distribution are determined by the ultrasonic method, heat fluxes are then determined from equation (2).
Figure 2 shows variations of estimated heat fluxes with elapsed time. It can be seen that heat flux
increases just after heating starts and decreases after the heating stops. It is noted that the heat flux
estimated by ultrasound agrees well with that estimated using the conventional heat flux sensor. A certain
random error in the ultrasonic result is observed. It is basically caused by the fluctuation in the estimatedtemperature. In this paper, further discussion on the advantage of using the ultrasonic method formaterial process monitoring is made.
Ultrasonictransducer

Ultrasonicpulser/receiver

A/D board

30 mm

Copperthin plate

Steel plate

30 mm

PC

Copper plate

Heat flux sensor

Gas burner

Figure 1 Schematic of the experimental setup.

Heat flux (kW/m2)

50
40
30
20
10
Ultrasound(h=5mm)
Heat flux sensor

0
-10
0

50
100
150
Elapsed time (s)Figure 2 Variation of the estimated heat flux at the heating surface with the elapsed time.

Acknowledgment
Support from the JSPS Grant-in-Aid for Scientific Research (B) 25289238 are appreciated.
References[1] M. Takahashi and I. Ihara, Ultrasonic Determination of Temperature Distribution in Thick Plates during Single
Sided Heating, Modern Physics Letters B, Vol. 22, No. 11(2008), pp.971- 976.
[2] I. Ihara and T. Tomomatsu, In-Situ Measurement of Internal Temperature Distribution of Sintered Materials
Using Ultrasonic Technique, IOP Conf. Series: Materials Science and Engineering, 18 (2011) 022008.
[3] H. Yamada, A. Kosugi and I. Ihara, Non-Contact Monitoring of Surface Temperature Distribution by Laser
Ultrasound Scanning, Japanese Journal of Applied Physics, Vol.50, No.7 (2011), 07HC06.
[4] G. E. Myers: Analytical Methods in Conduction Heat Transfer (McGraw-Hill, New York, 1971) p. 262.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ELASTIC-PLASTIC FRACTURE TOUGHNESS OF EXTRUDED Mg-Al-Zn ALLOYS
Zainuddin Sajuri1,*, Noradila Abdul Latif1, Junaidi Syarif2,
Yukio Miyashita3 and Yoshiharu Mutoh3
1

Department of Mechanical and Materials Engineering, Faculty of Engineering and Built Environment,
Universiti Kebangsaan Malaysia, 43600 UKM Bangi, Selangor, Malaysia
2
Mechanical Engineering Dept., College of Engineering, University of Sharjah, PO Box:27272, Sharjah, UAE
3
Department of Mechanical Engineering, Nagaoka University of Technology,1603-1, Kamitomioka, Nagaoka, Niigata 940-2188, Japan*sajuri@ukm.edu.my

Keywords: Elastic-plastic fracture toughness, ductile, J-integral, stretch zone, Mg-Al-Zn alloys
Introduction
Extruded Mg-Al-Zn alloys are categorized as ductile materials since the elongation to fracture is morethan 15% [1]. Therefore, conducting -integral test is preferable rather than the KIC linear elastic
fracture toughness test in investigating the fracture behavior and evaluating the plane strain fracture
toughness value of these alloys. The objective of this study is to determine the elastic-plastic fracture
toughness JIC of extruded AZ31 and AZ61 magnesium alloys.
Experimental Procedures
The materials used in this study were extruded AZ31 and AZ61 magnesium alloys. A standard multiple
specimens JIC test method was performed to determine the JIC-value of both alloys. Four fatigue
pre-cracked single edge notched bend (SENB) specimens were used to develop a resistance curve(R-curve) for identifying JIC-value of each material according to the ASTM E1820 [2]. The method todetermination of JIC-value was based on the critical stretch zone that measured from the fracturesurface as recommended by the JSME S001 for obtaining JIC-value [3].
Results and Discussion
The microstructures and nominal stress-strain curves of AZ31 and AZ61 are shown in Figs. 1 and 2,
respectively. In Figure 1, AZ61 indicates smaller grain size than that of AZ31. Further, a higher
twinning density has been observed in AZ31. The density of the twins in AZ61 was lower whencompared with the AZ31. This might indicate an increased resistance to twinning deformation in AZ61.
Meanwhile, the yield stress and tensile strength of AZ61 were higher compared to AZ31 as shown in
Fig. 2. This was most probably due to high aluminum content in AZ61 which resulted in decreasing
the grain size and increasing the  -phase volume that pinned the movement of dislocations and slips[4]. Figure 3 shows the load-load line displacement curves at different interrupted displacements of
AZ31 and AZ61. The J-value was evaluated using the load-load line displacement and geometricalsize of the specimen as 𝐽 = 2𝐴⁄𝐵𝑏, where, A is the area under the curve, B is specimen thickness andb is the remaining crack ligament. The ductile crack growth size (a) and the critical stretch zone
width (SZWC) of the break-opened specimen were then measured from the SEM micrograph. The
J-value, a and SZWC were summarized in Table 1. All these values were then used to develop theNominal stress, MPa

400
AZ61

300
AZ31

200
100
Strain rate of 1 x 10-3 s-1

0
0

Figure 1 Microstructures of (a) AZ31 and (b) AZ61

5

10
15
Nominal strain, %20

Figure 2 Nominal stress-strain
curves of AZ31 and AZ61

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

R-line and stretch zone line on the R-curve as shown in Figure 4. From the R-curve, the JIC-values of
extruded AZ31 and AZ61 magnesium alloys were found at 18.8 and 24.7kJ/m2, respectively.
Conclusion
J-integral was used to determine the elastic-plastic fracture toughness JIC of extruded AZ31 and AZ61.
The determination of JIC for both Mg-Al-Zn alloys was based on the interception of stretch zone line
and R-curve. The JIC of extruded AZ31 and AZ61 Mg alloys were 18.8 and 24.7kJ/m2, respectively.
2.5

2.5

AZ61

2.0

Load, kN

Load, kN

2.0

AZ31

1.5
1.0
0.6 mm
0.8 mm
1.0 mm
1.2 mm

0.5
0
0

1.5
1.0
0.6 mm
0.8 mm
1.0 mm
1.2 mm

0.5

0
0.5
1.0
1.5
0
0.5
1.0
1.5
Load line displacement, mm
Load line displacement, mm
Figure 3 Load-load line displacement curves at different interrupted displacements for AZ31 and AZ61
Table 1 J-value, crack growth size and critical stretch zone width of AZ31 and AZ61
Interrupted
displacement (mm)J (kJ/m2)
a (mm)SZWC(mm)

0.6
20.8
0.122
0.081

AZ31
0.8
1.0
25.1
41.7
0.227
0.329
0.089
0.092

1.2
47.8
0.428
0.092

AZ31

70

R-line30
2000Stretchzone line

AZ61

60

2

J-integral [kJ/m ]J-integral [kJ/m2]

Stretchzone line

40

10

1.2
67.1
0.404
0.054

80

60
50

AZ61
0.8
1.0
39.7
53.6
0.178
0.286
0.051
0.056

0.6
25.1
0.066
0.047

50

30
20

JIC-valuea SZWC

0.1
0.2
0.3
0.4
Crack growth,  a [mm]R-line40

JIC-valuea SZWC

10
0.5

0
0

0.1
0.2
0.3
0.4
Crack growth,  a [mm]0.5

Figure 4 JIC determinations for AZ31 and AZ61 using R-curves

Acknowledgement
The authors acknowledge the supports from the Universiti Kebangsaan Malaysia, Ministry of Higher
Education Malaysia (ERGS/1/2011/TK/UKM/02/9) and Nagaoka University of Technology.
References[1] T. Yokoyama, Impact tensile stress–strain characteristics of wrought magnesium alloys, Strain, 39-4 (2003),
167-175.
[2] ASTM E1820 Standard Test Method for Measurement of Fracture Toughness (2007)[3] JSME S001 Standard Method of Test for Elastic-Plastic Fracture Tougness, JIC (1981)(In Japanese)[4] M. S. Dargusch, K. Pettersen, K. Nogita, M. D. Nave, G. L. Dunlop, The effect of aluminium content on themechanical properties and microstructure of die cast binary magnesium-aluminium alloys, Materials
Transactions, 47-4 (2006), 977-982.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

FABRICATION AND EVALUATION OF MICRO-STRUCTURED REACTION FIELD
WITH VERTICALLY ALIGNED CARBON NANOTUBES FOR MICROBIO-ANALYSIS DEVICE
Yuma Suzuki1*, Ewelina Pabjańczyk-Wlazlo2, Tetsuhide Shimizu1, Ming Yang1
1

Advanced Materials and Processing Engineering Laboratory (AMPEL),
Graduate School of System Design, Tokyo Metropolitan University,2
Department of Material and Commodity Sciences and Textile Metrology, Lodz University of Technology* yuma-suzuki@ed.tmu.ac.jp

Keywords: Micro bio-analysis device, Carbon Nanotubes, Photolithography, Protein adsorption,
Ultraviolet Spectrometry
Micro bio-analysis device (MBD), such as micro Total Analysis Systems (-TAS) andLab-on-a-chip (LOC), is demanded to realize a rapid and high sensitivity diagnosis of biomolecules(for instance, viruses and proteins) [1]. The reaction field with micro structures in MBD can offer
lager specific surface area (surface areas per volume) and reduction of diffusion distance which
improves detection sensitivity and reaction efficiency [2]. Carbon nanotubes (CNTs) are one ofnanomaterials with chemical stability and high aspect ratio which brings in high specific surface area[3]. Furthermore, for distinct control of micro structural dimensions in the reaction filed, thefabrication of micro pattern array by combining CNTs synthesis and photolithography has been
attracting significant attention. However, designs of the structure and the surface of micro-structured
reaction field have not been investigated. In this work, we fabricated the reaction field with micro
structure array by patterning metal catalyst with photolithography and synthesizing vertically aligned
CNTs (VACNTs). For high detection sensitivity and reaction efficiency in MBD, the influence ofstructure dimensions and zeta potential on protein adsorption and diffusion time was investigated.
Patterned-VACNTs structures were synthesized by ultraviolet photolithography and catalytic
thermal CVD (Chemical Vapor Deposition). Fig.1 shows the fabrication process for patterned
VACNTs. Silicon (Si) wafer was used as substrates. First, positive photoresist was coated on thesubstrate by spin coater and ultraviolet light was exposed to resist-coated substrate covering aphotomask. After developing photoresist, metal catalysts, aluminum (Al) and iron (Fe), was depositedin sequence by sputtering. Coated photoresist was removed by organic solvent and VACNTs was
synthesized on metal catalysts by CVD.
The protein adsorption on the reaction field with patterned-VACNTs structures was evaluated byultraviolet spectrometry. First, 20l bovine serum albumin (BSA) solution was dropped on thereaction field. Then, the sample was incubated 15 and 30 minutes for protein diffusion and adsorption
on the reaction field. In this method, the reflectivity of protein-adsorbed reaction field was measured
and we compared reflectivity the wavelength of 205 nm which protein adsorbs [4] for various
dimension of patterned-VACNTs structure.
Fig. 2 shows the patterned-VACNTs structures on the substrate. VACNTs structures withdifferent pattern dimensions were fabricated. Protein adsorption on VACNTs structures wascharacterized by ultraviolet spectrometry (See Fig. 3). The results indicated that the adsorbed amountand time of protein depend on the structure dimensions and the zeta potential of VACNTs structures.
Wider structural pitch tends to adsorb more proteins with shorter time. On the other hand, narrower
pitch tends to adsorb more proteins longer time. The balance between structural dimensions and zeta
potential is one of important issues for designs of the reaction field with high detection sensitivity andreaction efficiency.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

1. Coating resist

2. Exposing UV

3. Developing

4. Depositing metal

Power

5. Removing resist

6. Synthesizing VACNTs

Reaction fieldwith patterned-VACNTs structures

Figure 1 Schematic image of the fabrication process for patterned VACNTs

a

20μm

b

20μm

c

20μm

d

20μm

Figure 2 SEM images of patterned-VACNTs structures with various dimensions.
(a) Structural diameter (D) is 10 m and structural pitch (P) is 10 m.
(b) D = 10 m, P = 30 m (c) D = 10 m, P = 50 m (d) D = 10 m, P = 70 m

BSA reflectivity at 205 nm[a.u.]

1
0.95

0.9
0.85
0.8
0.75

0.7

pitch 10μm

0.65pitch 30μm

0.6pitch 50μm

0.55pitch 70μm

0.5
10

15

20
25
Diffusion time [min]30

35

Figure 3 BSA reflectivity at the wavelength of 205 nm
in the reaction field with patterned-VACNTs structures

References[1] Emmanuel Delamarche, David Juncker, Heinz Schmid, Microfluidics for Processing Surfaces andMiniaturizing Biological Assays, Advanced Materials, 17 (2005), pp. 2911-2933.
[2] Sang Won Han, Seulah Lee, Juree Hong, Eunji Jang, Taeyoon Lee, Won-Gun Koh, Mutiscale substrates
based on hydrogel-incorporated silicon nanowires for protein patterning and microarray-based immunoassays,
Biosensors and Bioelectronics, 45 (2013), pp. 129–135.
[3] Ming Yang, Yakuya Yabe, Katsumi Uchiyama, Fabricarion of Micro Device for Rapid and High-Sensitive
Bio-Analysis, Journal of Solid Mechanics and Materials Engineering, 7-2 (2013), pp. 199-205.
[4] Robert K. Scopes, Measurement of protein by spectrophotometry at 205 nm, Analytical Biochemistry, 59-1
(1974), pp. 277-282.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

IMPROVEMENT OF MECHANICAL PROPERTIES IN UNTWISTED CNT YARNBY PAA/DMSO DENSIFICATION TREATMENT
Tae Sung Kim1, Kaoru Sugano1, Atsushi Hosoi2, Hiroyuki Kawada3,*
1

Graduate student, Department of Applied Mechanics, Graduate School of Fundamental Science andEngineering, Waseda University, 3-4-1 Okubo, Shinjuku, Tokyo 169-8555, Japan
2
Assistant Professor, Department of applied Mechanics and Aerospace Engineering, Waseda University, 3-4-1
Okubo, Shinjuku, Tokyo 169-8555, Japan
3
Professor, Department of applied Mechanics and Aerospace Engineering, Waseda University, 3-4-1 Okubo,
Shinjuku, Tokyo 169-8555, Japan*Corresponding.author: kawada@waseda.jp, Tel:+81-3-5286-3261, Fax:+81-3-5273-2667

Keywords: CNT yarn, tensile properties, densification treatment, heat stretching, polymer solution
INTRODUCTION
Carbon nanotube (CNT) is increasingly applied as a reinforcement of polymer matrix composite
because of extremely high mechanical properties. Among several forms of CNT reinforcement,
continuous CNT yarn, which is fabricated by drawing Multiwall Carbon Nanotubes (MWCNTs)
continuously from vertically aligned MWCNT arrays, can be a next-generation reinforcement whichenables the use of CNTs in the macro-scale. Strength development mechanism of CNT yarn is thought
to depend on the interactive force such as the van der Waals forces between CNTs. Therefore,
improvement of stress transfer capability is required for the enhancement of CNT yarn. There havebeen various methods to enrich the interaction between CNTs. Especially densification treatment usinga polymer solution is thought to be an effective and simple method. However, the tendency of theeffective polymer for densification treatment is not fully understood. There is still a lot of room for
improvement in selecting polymers. In this study, we fabricated untwisted CNT yarn which has higher
mechanical properties than other CNT yarns. Untwisted CNT yarn was polymer treated for the
purpose of densification. Then, we selected polymer considering its molecular structure and physicalproperty. Furthermore, applying heat stretching treatment in addition to densification treatment,
mechanical properties of untwisted CNT yarn was improved.
MATERIALS
Untwisted CNT yarn is fabricated by passing CNT sheets through a die. Figure 1 shows a schematic
view of drawing untwisted CNT yarn. CNT sheets are drawn continuously from vertically aligned
MWCNT arrays. MWCNTs are synthesized on the Si wafer substrate by atmosphere pressurechemical vapor deposition (AP-CVD). The diameter of fabricated CNT yarn is about 35 m.
Aligned MWCNT array
Untwisted CNT yarn

CNT sheetDie

Si wafer substrate
Figure 1 Schematic view of drawing untwisted CNT yarn.

EXPERIMENTAL
Polymer treatments are conducted to CNT yarns for the purpose of densification. First, CNT yarn isimmersed in Polyacrylic acid/ Dimethyl sulfoxide (PAA/DMSO) solution with the condition of 60°C
for 3 hours. After that, polymer treated CNT yarn is dried for 1 hour at 150°C to evaporate DMSO in it.
Then, CNT yarn is applied 150mN load, which is less than 10 % of braking force, and heat stretchingtreatment is conducted.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Cumulative probability offiber failure lnln(1/(1-Pf))

RESULTS AND DISCUSSIONS
Figure 2 shows Weibull distributions of CNT yarn strength obtained by single fiber tensile test. Here,‘Untreated’, ‘PAA’ and ‘PAA-T’ denotes untreated CNT yarn, PAA/DMSO treated CNT yarn and heat
stretching treated CNT yarn after PAA/DMSO treated, respectively. In the case of 5 wt% PAA solution
concentration, tensile strength of CNT yarn is improved from 0.90 GPa up to 2.35 GPa bydensification and heat stretching treatment. Young’s modulus is increased from 60.9 GPa to 172 GPain 7 wt% PAA solution concentration. Figure 3 shows SEM images of CNT yarn surface. There issome disarray of constituent CNTs on the surface of polymer treated CNT yarn (Figure 3 (a)). On theother hand, the orientations of constituent CNTs are by heat stretching treatment (Figure 3 (b)).
Contact area between CNTs was increased, and interactive force of CNTs and mechanical properties of
CNT yarn was highly improved.

2
1

PAA-7wt.%-T
PAA-7wt.%
PAA-5wt.%-T
PAA-5wt.%
DMSO only
Untreated

0
-1
-2
-3
-4
-5
6

6.5

7

7.5

Tensile strength ln MPa

8

Figure 2 Comparison of mechanical properties for treated fibers.

1.0 m(a) PAA

1.0 m(b) PAA -Tension

Figure 3 SEM images of the surface morphology of PAA/DMSO treated fibers.

Conclusion
In this study, untwisted CNT yarn was polymer treated using Poly acrylic acid/Dimethyl sulfoxide(PAA/DMSO) solution for the purpose of densification. As a result of single fiber tensile test, strengthand Young’s modulus of PAA/DMSO treated CNT yarn were drastically increased. Furthermore, as aresult of heat stretching treatment in addition to densification treatment, strength and Young’s modulusof CNT yarn was improved up to 2.35 GPa, 172 GPa respectively. After observation of CNT yarn
surfaces using SEM, disarray of CNTs with heat stretching treatment was remedied. Contact area
between CNTs was increased, and interactive force of CNTs and mechanical properties of CNT yarn
was highly improved.
References[1] Jiang KL, Li QQ, Fan SS. Spinning continuous carbon nanotube yarns. Nature 2002; 419 (6909): pp. 801.
[2] Sugano K, Kurata M, Kawada H. Evaluation of mechanical properties of untwisted carbon nanotube yarn for
application to composite materials. Carbon 2014; 78: pp. 356-365.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

EFFECTS OF PROCESSING CONDITIONS IN UNIDIRECTIONAL CARBONFIBER THERMOPLASTIC PREPREG TAPE LAYING USING NEAR INFRAREDHEATER
Daiki TANABE1,*, Kazuaki NISHIYABU2, Tetsusei KURASHIKI1
1

Dept. of Management of Industry and technology, Graduate School of Eng., Osaka University2
Dept. of Mechanical Engineering, Faculty of Science and Eng., Kinki University* d-tanabe@mit.eng.osaka-u.ac.jp

Keywords: ATL, CF/PA6, CFRTP, Near infrared heating, Interlaminar shear strength
Automated thermoplastic tape laying (ATL) is a high-potential manufacturing method forcontinuous fiber reinforced thermoplastic composites which can ensure operational safety even inhighly loaded components such as aircraft wing skins and fuselages or construction large scaled
structures [1]. This method often uses computer-controlled robotics to lay one or several layers ofcarbon fiber thermoplastic tape onto a mold, and should be optimized some processing parameterssuch as heating, pressing and feed speed and so on. Though the final aim of this study is to
manufacture composite laminates made from tapes pre-impregnated with unidirectional carbon fiber
reinforcement and a thermoplastic matrix, this study focuses on thermal distribution and joining
strength in a laying process of unidirectional carbon fiber reinforced thermoplastic tape. The tape arecommonly heated up to the melting temperature by various heating sources including heated roller, hot
gas torch, near-infrared light lump and diode laser and so on [1, 2]. A near-infrared heater has someadvantages such as rapid start-up speed, suitable for heating of carbon fiber and low cost. However,
the heating characteristic is influenced significantly by position of the heater and feed rate of prepreg
tape, thus it is necessary to optimize these heating parameters. Therefore, it is important for thethermoplastic prepreg tape laying process to investigate the heating and consolidation behavior. This
study aimed to predict the optimum processing condition for thermoplastic tape laying. The material
used for the experiment was carbon fiber reinforced polyamide 6 prepreg tape (TenCate, CETEX®
TC910). This prepreg tape has unidirectional construction with a resin content of Vf=58vol.% and a
thickness of t=0.16mm.
Figure 1 shows appearance of the tape laying machine authors made originally. This tapelaying machine has the near infrared heater (Heraues, ZKB600/80G) as source of heating and thepressure rollers. The heater has two tungsten filament as heating source, and it is suitable as the heaterof CFRTP prepreg tape because the absorbed fraction is high into black body as carbon fiber. The
pressure rollers are heated rapidly to any temperature by a induction heating equipment (Hidec Co.,
Ltd.). This tape laying process is divided into the three steps. Firstly, the prepreg tapes are heated by
the near infrared heater. Then, the matrix polymer of prepreg tapes is melted rapidly. Secondly, two
prepreg tapes are welded by pressing using pressure rollers. Finally, the laminated tape is cooled down.
Two fine thermocouple (ANBE SMT Co., Ltd., KSG-40-100-100, type K) were bonded on the prepreg
tape surface to investigate the fusion behavior. Two prepreg tapes with 50mm in width were inserted topressure roller at the angle of θCF=60º. The fiber direction of two prepreg tapes and the feed direction
is same direction. The angle of the near infrared heater was fixed at θH=30º, and the feed speed waschanged from v=14mm/s from to v=28mm/s. The heating behavior of the surface of prepreg tape wasmeasured by fine wire thermocouples using a temperature measuring instrument (Graphtech Co., Ltd.,midi LOGGER GL200A). The images of joint surfaces peeled off after joining were imported with ascanner device (Epson Co., Ltd., ES-7000H). The welding surfaces were also observed with amicroscope to investigate the fusion joining interface. The single lap shear strength test was carriedout to evaluate a joint strength by using universal testing machine (SHIMAZU Co., Ltd., AG-50kN
XDplus). Figure 3 shows appearance of single lap joining test specimen. Before the testing, aluminumtabs were bonded to end of specimens with epoxy adhesive. The cross-head speed was v=0.5mm/min.
Figure 2 shows the change of surface temperature of CF/PA6 prepreg tape at v=14mm/s. The
temperature of pressure roller was TR=30ºC. The temperature of prepreg tapes were increased withincreasing the x-coordinate. When the prepreg tapes were cooled down quickly, the tapes were
contacted pressure rollers. Also, when the distance of heater was xH=50mm and xH=60mm, thePage 9

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

temperature was over T=225ºC. This temperature is melting temperature of PA6 polymer. The
maximum temperature was increased with decreasing the distance between heater and pressure rollers(xH). It was also revealed that the range of heating was increased with decreasing the distance of heater.
From these results, it is considered that the proper distance of the heater is from 50mm to 60mm.
Figure 3 shows the result of single lap shear strength test. The lap shear strength (τap) and the
maximum load (Pmax) were increased with decreasing the distance between heater and pressure roller(xH). The feed speed was v=14mm/s, the lap shear strength and maximum load were higher thanv=24mm/s because the matrix polymer of prepreg tapes melted sufficiently.
From these results, the effects of processing conditions such as number and position of infrared
heaters, feed speed and temperature of pressure roller on the laying behavior of prepreg tape wererevealed to obtain the proper processing parameters.
Table and Figure
UD-CF/PA6 prepreg
Thermocouple
Near-IR heater

Induction heating coilFeed speed, vzySupport rollerx

Thermocouple
Pressure roller
xH

Figure 1

Appearance of ATL machine.

Roller contact zone

300
250
200
150
100
50
0
-150

-100
-50
0
x-coordinate, x (mm)x𝑥𝐻=50mm
H=50mmx𝑥𝐻=60mm
H=60mmx𝑥𝐻=70mm
H=70mm

300
250
200
150
100
50
0
-150

50

(a) Top surface

-100
-50
0
x-coordinate, x (mm)50(b) Lower surface

Change of surface temperature of CF/PA6 prepreg tape at v=14mm/s.
800

8
6
4
2
0

Maximum load, Pmax (N)Lap shear strength, τap (MPa)Figure 2

Roller contact zone

350

x𝑥𝐻=50mm
H=50mmx𝑥𝐻=60mm
H=60mmx𝑥𝐻=70mm
H=70mmSurface temperature, T (ºC)

Surface temperature, T (ºC)

350v=14mm/s

600
400v=28mm/s

200
0
40

50

60

70

DistanceDistance
between heaterand consolidation
betweenheaterroller,
andxH (mm)consolidation roller, xH (mm)Figure 3

Result of single lap shear strength test of CF/PA6 laminated tapes.

References[1] C.M. Stokes-Griffin, P. Compston, A combined optical-thermal model for near-infrared laser heating ofthermoplastic composites in an automated tape placement process, Composites Part A: Applied Science andManufacturing, In Press, Corrected Proof, Available online 14 August 2014.
[2] Z. Qureshi, T. Swait, R. Scaife, H.M. El-Dessouky, In situ consolidation of thermoplastic prepreg tape using
automated tape placement technology: Potential and possibilities, Composites: Part B 66 (2014) 255–267.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ESTIMATION OF ELECTRICAL CONDUCTIVITY OF TiB2/Al COMPOSITESWITH IMAGE ANALYSIS
Kenjiro Sugio1*, Nariaki Kawano2, Takaaki Hirose2 , Yong-Bum Choi1 and Gen Sasaki1
1

Division of Materials and Production Engineering, Institute of Engineering, Hiroshima University,
1-4-1, Kagamiyama, Higashi-Hiroshima, Hiroshima 739-8527, Japan2
Department of Mechanical Science and Engineering, Graduate School of Engineering, Hiroshima University,
1-4-1, Kagamiyama, Higashi-Hiroshima, Hiroshima 739-8527, Japan*Corresponding ksugio@hiroshima-u.ac.jp

Keywords: electrical conductivity, image analysis, particle dispersibility
In the present study, it was aimed at the development of the software that is possible to calculate theeffective electrical conductivity of aluminum-based particle-dispersed composites from the actual
microstructure image by performing conduction simulation. However, the electrical conductivity
calculated with the microstructure image is supposed to be not equivalent to the electrical conductivity
measured with experimental method, such as four-terminal method, because microstructure image is a
two-dimensional information though the experimental conductivity is obtained from three-dimensional
microstructure. It is not well studied on the relationship between two-dimensional andthree-dimensional electrical conductivity. Therefore, statistical relationship between three-dimensional
and two-dimensional electrical conductivity were investigated by performing conduction simulation.
First, a three-dimensional particle distribution was created in the simulation cell, and the electricalconductivity in three-dimension was calculated. Then, the sequential two-dimensional cross-sections
were cut out from the three-dimensional particle distribution. The electrical conductivity for eachcross-section was calculated. Electrical conductivity of Al, TiB2, SiC and Al2O3 were set to 40.0, 11.1,
0.01 and 1.0×10-18 (106/Ω/m), respectively. Three-dimensional cell was divided into 100 × 100 × 100
elements, and electrodes whose size is 100 × 100 × 5 are added to the both sides. The size of elements
was 2 × 2 × 2 (μm). Diameter of reinforcing particles was 10μm, and their dispersibility were changed
by using uniform random number and changing standard deviation of normalized random number.
Two-dimensional cross-section which cut out from the three-dimensional cell was divided into 100 ×
100 elements, and electrodes whose size is 100 × 5 were added to the both sides. The degree ofdispersion of the particles in the cross-section was evaluated by two-dimensional local number(LN2DR) [1]. The potential difference between left and right sides set as 0.05V in three and two
dimensions. Finite volume method was used to calculate the steady state of electrical potential
distributions. Calculations were stopped when the average variation of potential in each elements issmaller than 10-13. Electrical conductivity was calculated from the potential distribution of the steadystate.
Fig. 1 shows the electrical conductivity of
three-dimensional cell and two-dimensionalcross-sections in 20vol.%TiB2/Al composite where
TiB2 particles were uniformly distributed. Opencircles indicates the electrical conductivity of thetwo-dimensional cross section which is extracted
from the three-dimensional cell. The horizontal axis
represents the position of the cross section. The
particle arrangement of typical two-dimensional
sections are shown in the figure. Broken line
indicates the average electrical conductivity of thetwo-dimensional cross-sections and each electricalconductivity (open circle) is not apart from theaverage (broken line) in the case of uniform
random distribution. Solid line indicates the Fig. 1 Relation of electrical conductivities ofelectrical conductivity of the three-dimensional cell, 3D and 2D cross section in uniform random
and the average value of two-dimensional distribution of 20vol.%TiB particles.
2

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Electrical conductivity(10-6m-1Ω)-Electrical conductivity(10-6m-1Ω-1)cross-sections is slightly lower than the40
three-dimensional value. Series of calculations with35
uniform random distribution were carried out by
30changing the kind of the reinforcement (TiB2, SiC,
Al2O3) and the volume fraction (5% to 50%). Fig. 2
25shows a summary of those calculations. Closed
20
circle indicates electrical conductivity ofTiB2 3D
15
three-dimensional cell dispersing the TiB2 particles,
TiB2 2D
SiC 3Dand open circles indicates the average electrical10
SiC 2D
conductivity of the two-dimensional cross-sections.
Al2O3 3D
5The
averageelectricalconductivityof
Al2O3 2D0
two-dimensional cross-sections is lower than that of
0
20
40
60
three-dimensional cell, and the difference increase
Volume Fraction (%)
with increasing volume fraction. Electricalconductivity of Al and TiB2 were 40.0 and 11.1 Fig. 2 Relation of electrical conductivities of(106/Ω/m), respectively, and the difference is not so 3D and 2D cross section in uniform random
large. Therefore, there is no large difference distribution of TiB2, SiC and Al2O3 particles.
between three-dimensional and two-dimensional
5vol.% 3D
40
5vol.% 2Delectrical conductivity in case of TiB2 particles.
10vol.% 3D
10vol.% 2D
Meanwhile, electrical conductivity of SiC and
35
15vol.% 3D-18
615vol.% 2D
Al2O3 were 0.01 and 1.0×10(10 /Ω/m),20vol.% 3D
respectively, and difference against Al is too large.
20vol.% 2D
30
25vol.% 3DThis possibly enhanced difference between
25vol.% 2D
30vol.% 3D
three-dimensional and two-dimensional electrical25
30vol.% 2D
conductivity in case of SiC and Al2O3. Fig. 3 shows
35vol.% 3D
35vol.% 2D
results of the calculation by changing the
40vol.% 3D
20
40vol.% 2D
dispersibility of the TiB2 particles. The horizontal
45vol.% 3D
axis is LN2DRvar which can represent dispersibility
45vol.% 2D
15
50vol.% 3D
0
10
20
30
40
quantitatively. If LN2DRvar is 7, gravity centers of
50vol.% 2D
LN2DR
dispersed particles are arranged in uniform random
distribution. If LN2DRvar increases, clustering
tendency of the particles increases. Filled markers Fig. 3 Relationship between electricalindicatesthethree-dimensional
electrical conductivities and dispersibility of TiB2
conductivity, and open markers indicates the particles.
average of the two-dimensional electricalconductivity. The difference between two-dimensional and three-dimensional electrical conductivity isconstant in each volume fraction even when dispersibility is changed. This means that it is notnecessary to consider the dispersibility when converting the two-dimensional conductivity tothree-dimensional conductivity. In addition, the same trend was seen in the case of SiC and Al2O3
particles. Table 1 shows list of correction value for converting two-dimensional electrical conductivity
to three-dimensional electrical conductivity. Although maximum correction value for TiB2 is 105.45%,
maximum correction value for SiC and Al2O3 is about 171%.
var

Table 1 List of correction value (%) for converting 2D electrical conductivity to 3D electricalconductivity.
σp/σAl5vol.%

10vol.%

15vol.%

20vol.%

25vol.%

30vol.%

35vol.%

40vol.%

45vol.%

50vol.%

TiB2

0.2775100.80101.49102.10102.72103.32103.95104.45104.70105.07105.45SiC

2.5×10-4

102.87106.27109.67112.65117.27122.41129.74135.29145.20171.04Al2O3

2.5×10-20

102.85105.89108.78112.53116.84122.94129.87135.47145.64171.83

References[1] K. Sugio et. al., Evaluating Spatial Distribution Randomness of Overlap Permissive Second Phase in Threeand Two-Dimensions, Mater. Trans., 48(2007) pp. 2762-2767.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

NEW MANUFACTUTING METHOD OF PARTICLE INTERMETALLICCOMPOUND REINFORCED COMPSOITES
Yongbum Choi 1*, Zhefeng Xu1, Kenjiro Sugio1, Kazuhiro Matsugi1, Gen Sasaki 1
1

Division of Materials and Manufacturing, Graduate School of Engineering, Hiroshima University, 1-4-1
Kagamiyama, Higashi-Hiroshima, Hiroshima 739-8527, Japan*Corresponding. ybchoi@hiroshima-u.ac.jp

Keywords: Intermetallic, Composites, Porosity, Infiltration and reaction method.
Conversion to aluminum alloy from cast iron materials has been increasing. And advanced featuresof the aluminum alloy are demanded. Due to this, composites which are strengthened with ceramicparticles in aluminum alloy are developed. It is applied in practical for brake disk and piston in mobile
parts in the industrial field [1]. Generally, manufacturing methods of a ceramic particles dispersed
composite include Powder metallurgy process [2], Stir casting [3] and squeeze casting method [4], etc.
A new process in this study is proposed to fabricate an aluminum alloy matrix composite dispersedwith refining intermetallic compound particles by using the reaction between porous nickel and molten
aluminum alloy. The objective of this study is to investigate the effects of the specific surface area on
the intermetallic compounds formed in the composites by infiltration and reaction method. Therefore,
the research investigated on the reaction of porous nickel and molten Al alloy, the amounts of theintermetallic compounds by different specific surface area of porous nickel and the aspect ratio ofintermetallic compounds under 1250m2/m3, 2800m2/m3, 5800m2/m3and > 5800m2/m3, respectively.
A366 alloy in ASTM, which composition of Al-12Si-1Ni-1Cu-1Mg (mass %), was used as matrix.
Preform was porous nickel (Toyama Sumitomo Electric Co., Ltd.). Volume ratio of porous nickel is
4∼6vol.%. Porous nickels with four different kinds of specific surface areas (a: 1250 m2/m3, b: 2800
m2/m3, c: 5800 m2/m3and d: > 5800 m2/m3) were used in the experiment to examine the reactive behaviorof the intermetallic compound. Infiltration and reaction method was used to fabricate the composites.
Temperature of molten Al alloy and applied pressure to molten Al alloy were 973K and 0.1MPa.
Figure 1 shows microstructure of Al matrix composites dispersed intermetallic compound with specificsurface area of porous: a) 1250 m2/m3, b) 2800 m2/m3, c) 5800 m2/m3 and d) >5800 m2/m3. As a resultof observing the microstructure of the composites manufactured using four kinds of porous nickel, on
the whole, the intermetallic compound were distributed inside matrix by the reaction of a porous nickel
and aluminum. However, the microstructure observed from the composites fabricated with low specificsurface area (in figure 1(a) 1250 m2/m3) that almost all porous nickel were changed to intermetallic

Figure 1 Microstructure of Al matrix composites dispersed intermetallic compound composites

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

compound. Area fraction of the generated intermetallic compound is about 15%. In addition, the portion
which remains as porous nickel body without delamination from the porous surface was also observed.
However, the area fraction of the un-reacted nickel was 2.8%. And most of the fine Al3Ni intermetallic
compounds (below 0.1mm2 calculated area) were homogeneously dispersed inside matrix. But shape ofthe needle-like Al3Ni was observed too. And numerous pores were observed in the composites. The
microstructure when specific surface area is high (in figure 1(d) >5800 m2/m3), more fine intermetallic
compounds were distributed as compared with other materials in figure 1(a), (b) and (c). The fineintermetallic compounds increased with increasing specific surface area of porous nickel. In addition,
based on the figure 1(d). Area fraction (%) of intermetallic compound Al3Ni was about 28.2%, which isalmost twice comparison with figure 1(a). Furthermore, the area fraction of un-reacted nickel reached aminimum value of 0 % when the specific surface was >5800 m2/m3. The results indicate that higher the
specific surface area of porous nickel is, larger the contact surface with molten Al is, and it provided
further reaction of nickel with Al.
As shown, figure 2 is the porosity in Al matrix composites dispersed intermetallic compound
according to the variation of the specific surface area of porous nickel. Porosity in porous nickel, it was
occurred by initial polymer foam to fabrication of porous nickel. Polymer foam can be removed fromthe metal/polymer by heat-treatment. Pore was remained by the hollow struts inside porous nickel.
Porosity in porous nickel is different by specific surface area. The porosity inside porous nickel
decreased by the increasing a specific surface area of porous
1.8

nickel. Although there was almost similar porosity in the

Pore in porous nickel

1.6

composites infiltrated with porous nickel of three different

Pore in Matrix

1.4

specific surface area, pore size in the composites infiltrated withsurface area. The general porosity trend in both porous nickel
and matrix is reducing with the increasing of the specific surface
area at the same fabrication conditions, and the minimum
porosity is about 0.17% in the matrix and 0.19% in the porous
2

3

nickel of >5800 m /m since with the reduction of the porous
thickness accelerated the reaction between Ni and Al anddelamination caused by the difference of thermal expiation
coefficient with nickel, which is a great reduction of defect in

1.2

Porosity(%)

high specific surface are is smaller than those with low specific

1
0.8
0.6
0.4
0.2
0
8000
>5800 10000
Specific surface area, S/m2/m3

020002800 4000
1250

6000
5800

Figure 2 Porosity of Al3Ni, intermetallic
compound and Un-reacted nickel.

the composites as well.
References[1] Y. B. Choi, G. Sasaki, K. Matsugi, S. Kondoh, O.Yanagisawa, Low pressure casting process of FeCrSi/A366.0
alloy composites and their characterization. Key Eng. Mater. 326-328(2006), pp.1741-1744[2] M.H. Lee, Y.B. Choi, K. Sugio, K. Matsugi, G. Sasaki, Effect of aluminum carbide on thermal conductivity of
the unidirectional CF/Al composites fabricated by low pressure infiltration process. Compos. Sci. Technol.,97(2014), pp.1-5.
[3] H. Izadi, A. Nolting, C. Munro, D. P. Bshop, K. P. Plucknett, A. P. Gerlich, Friction stir processing of Al/SiC
composites fabricated by powder metallurgy. J. Mater. Proc. Technol., 213. 11(2013), pp.1900-1907.
[4] V. Bharath, N. Madev, V. Auradi, S.A. Kori, Preparation of 6061Al-Al2O3 MMCs by stir casting and evaluationof mechanical and wear properties. Proc. Mater. Sci., 6(2014), pp. 1658-1667

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

THE EFFECT OF ZRO2 ADDITION ON THE ELECTROCHEMICAL BEHAVIOR
OF LOW-CARBON STEELS IN CRO3–H2SO4 ELECTROLYTE
Kazuma Hashimoto1, Kenjiro Sugio1*, Gen Sasaki1, Asep Ridwan2, Aditianto Ramelan2
1

Division of Materials and Manufacturing, Graduate School of Engineering, Hiroshima University, 1-4-1
Kagamiyama, Higashi-Hiroshima, Hiroshima 739-8527, Japan2

Department of Materials Engineering, Faculty of Mechanical and Aerospace Engineering,
Bandung Institute of Technology, Bandung, Indonesia*Corresponding. ksugio@hiroshima-u.ac.jp

Keywords
Electrodeposited Cr, Nanocomposite coating, Zirconia particle, Polarization behavior,Electrophoretic deposition
Hard chromium coating is mainly used to improve the surface performance of pistons, brake disks,and gears. In the future, better abrasion resistance and better corrosion resistance of hard Cr coating aredemanded. Previous studies showed that co-deposition of ceramic particles such as Al2O3/SiC in Cr coating
improved these characteristics1). The Cr-ZrO2 composite coating is one of the prospective candidates for thispurpose. However, the effect of ZrO2 addition on the electrochemical behavior of carbon steels in
CrO3–H2SO4 electrolyte is unknown. In this study, the polarization behavior of low carbon steel in
CrO3-H2SO4 electrolyte with ZrO2 addition was investigated.
Low-carbon steels with the size of 1 cm×1 cm were prepared for the substrate. The specimens werecovered with thermoset polymer and were polished by waterproof abrasive paper from #200 to #2000. The
electrolytes consisted of CrO3 (250g/L) and H2SO4 (2.5g/L), and the additive amount of ZrO2 particles wasvaried from 0 g/L to 2.0 g/L. The average size of ZrO2 particles was 30 nm. The cyclic voltammetry
measurements were carried out by VersaSTAT-300 (Princeton™ Applied Research). A saturated calomelelectrode was used as the reference electrode. The potential scanning rate was 0.5mV/s from -1.5V to 0.6V.
Fig. 1(a) shows the polarization curves of each ZrO2 amount. The shape of polarization curve waschanged by the addition of ZrO2 particles. Polarization point can be confirmed by the plot of the potential
versus logarithmic value of the absolute current. Polarization point was shifted higher by the addition ofZrO2 particles. Fig. 1(b) summarizes the polarization potential at different ZrO2 density. The polarization
potential approaches to zero with increasing the amount of ZrO2.
was improved by the addition of ZrO2 particles. Fig.1

This suggests that the corrosion resistance(c) shows the cathodic polarization curves for thedeposition of Cr-ZrO2. The vertical axis is current value between -0.1 and 0A and the horizontal axis ispotential value between -0.6 and -0.4V. The slope of the plot indicates electrical conductivity of thecathode including the substrate and the plated film. The slope become gentler with increasing theamount of ZrO2, and this means that addition of ZrO2 particle decrease electrical conductivity of the

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plated film. Reductive reaction without ZrO2 stopped rapidly around -0.52V as shown in Fig 1(c). In
contrast, reductive reaction stopped gently around -0.45V, when ZrO2 was included in solution. Thissuggests that reductive reaction continued after the end of reductive reaction of Cr. Fig.2 shows the
specimen after each experiment. The gross of the surface decreased with the amount of ZrO2 particles.
This suggests that ZrO2 particles were introduced into the coating. Previous study showed thathydrogen ion was adsorbed on ZrO2 particles and they were positive charged in strong acidity. The
positive charged ZrO2 particle were reacted at the cathode. Fig.3 shows optical microscope images of thespecimens after each experiment. The particle size was bigger by the addition of ZrO2. And lattice defectconcentration decreased by the addition of ZrO2. But it did not depend on the amount of ZrO2. This suggeststhat ZrO2 improve mechanical characteristic of Cr coating.

Figure 1: (a) Polarization curves on each ZrO2 density, (b)polarization potential versus ZrO2 density and (c)cathodic
polarization curves on each ZrO2 density.

Figure 2: Specimens after plating. (a) 0 g/L, (b) 0.5 g/L, (c) 1.0 g/L and (d) 2.0 g/L ZrO2 particles were included in theelectrolytes.

Figure 3: Optical microscope images of the specimens after plating. (a) 0 g/L, (b) 0.5 g/L, (c) 1.0 g/L and (d) 2.0 g/L
ZrO2 particles were included in the electrolytes.

Reference
1.

Jifeng Gao, Jinping Suo, Preparation and characterization of the electrodeposited Cr-Al2O3/SiC composite
coating, Applied Surface Science 257 (2011), pp.9643– 9648

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PRODUCTION AND IMPROVEMENT OF CHARACTERISTICS OF SINTEREDPOROUS TITANIUM USING CAPSULE-FREE HIP METHOD
Shuichi MASHINO 1, Takuji OKABE 2*
1

Graduate Student, Hiroshima Institute of Technology, 2-1-1 Miyake, Saeki-ku, Hiroshima 731-5193, Japan2
Department of Mechanical Systems Engineering, Hiroshima Institute of Technology,
2-1-1 Miyake, Saeki-ku, Hiroshima 731-5193, Japan*okabe@cc.it-hiroshima.ac.jp

Keywords: Capsule-free hot isostatic pressing, Sintering, Porous material, Porosity, Titanium
1. Introduction
Porous materials are seen as new lightweight, functional materials with properties such as low
density and large surface area [1]. However, these materials also have problems such as low
mechanical strength and inconsistent porosity. The production of sintered porous materials by
capsule-free hot isostatic pressing (HIP) is capable of simultaneous control of both mechanicalstrength and porosity [2]. An additional feature of this method is that it does not necessarily require afoaming agent. In this study, we produced sintered porous titanium using the capsule-free HIP method
and attempted to improve material characteristics by investigating the effects of capsule-free HIP on
the sintered compacts.
2. Experimental Procedure
First (Experiment 1), commercial titanium powders (99.5% purity, -350 mesh, 4 g) were placed ina 15-mm diameter die and compacted using a hydraulic hand pressing machine (uniaxial compaction
loads: 4.9, 9.8, 14.7, 19.6, and 24.5 kN). The green compacts were then sintered using the capsule-free
HIP method at a holding temperature of 1523 K, under an argon gas atmosphere, at a holding pressureof 150 MPa, for a holding time of 1 h. In addition, the green compacts were sintered by electric
furnace at the relatively low temperature of 873 K, under an argon gas atmosphere, at normal pressure,for a holding time of 1 h. The pre-sintered compacts were then subjected to the capsule-free HIPtreatment under the same conditions as described above. The characteristics of the sintered compacts
were evaluated by microstructure observation (optical microscope; OM), component analysis (X-ray
diffraction), porosity measurement (image analysis software), density measurement (Archimedes
method), and compression test (universal testing machine).
Next (Experiment 2), commercial titanium powders (99.5% purity, -150 mesh, 4 g) wereclassified by sieves having meshes of 106 and 150 μm. Using the sieved powders (≤106 μm and 106－
150 μm), experiments and evaluations were performed as described for Experiment 1.
3. Results and Discussion
Experiment 1: Examples of OM images of the sintered compact surfaces are shown in Fig. 1.
Measurement of mean porosity at the surface and relative density (i.e., bulk density to true density) for
Ti: 4.51 Mg/m3 are shown in Fig. 2. These figures show that the sintered compact that was subjectedto the direct capsule-free HIP treatment definitely has some pores, but the porosity is much lower thanthat of the pre-sintered compacts. The relative densities of the direct HIPed compacts are more than
70%, indicating that these compacts have generally low porosity. Conversely, the compacts that werepre-sintered have porosities that are more than twice those of the direct-HIPed ones, and the relative
densities are 55 %－65 %, indicating a significant density reduction. The experimental results showthat pre-sintering is effective in the production of sintered porous compacts using the capsule-free HIP
method. The shapes of most pores are rounded, which indicates that the pores were formed
isotopically by the high-temperature, high-pressure gas of the HIP treatment, leading to an
improvement of mechanical strength by the relaxation of stress concentrations. In such processing
conditions, however, large variation of characteristics was confirmed. Therefore, it was necessary toinvestigate its improvement.
Experiment 2: The results of the capsule-free HIP treatment after pre-sintering are shown in Fig.

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50 μm

Direct Capsule-free HIP
Direct Capsule-free HIP

Pre-sintered+HIP
Pre-sintered+HIP

80
70
60
50
40
30
20
10
0

80
70
60
50
40
30
20
10
0
4.9

(b) Pre-sintered＋HIP

9.8

14.7

19.6

Relative Density / %(a) Direct Capsule-free HIPPorosity / %Compaction Load : 4.9 kN

Relative Density
Porosity

24.5

Compaction Load / kN
Fig. 2. Porosity at surface and relative density ofsintered porous compacts (Experiment 1).

Fig. 1. OM images of sintered porouscompacts at surface (Experiment 1).

－106 μm－106 μm

106－150 μm
106－150 μm
60

50

50

40

40

30

30

20

20

10

10

0

0
4.9

9.8
14.7
19.6
24.5
Compaction Load / kN
Fig. 3. Porosity at surface and relative densityof sintered porous compacts (Experiment 2).

Porosity / %Porosity / %Relative Density
Porosity
60

Relative Density / %3 for various compaction loads. This result shows that these compacts have porosity that is equivalent
to the pre-sintered ones in Experiment 1 (Fig. 2). Moreover, porosity decreases and density increases
with increasing compaction load. The variation of porosity with processing conditions is relatively
subdued compared with Experiment 1. This trend is attributed to the effect of homogenization of theparticle size of raw powders. Therefore, homogenization of the green compact is expected to lead to an
improvement of characteristics of the sintered compacts. Mean porosity measurements at the crosssection are shown in Fig. 4. The measurement points were at three locations (top, side and center) on
the cross section. From this result, the locational variation of porosity is large inside the sinteredcompact. In particular, the center portion is densified, and its porosity is reduced to about half that of
the surface. This trend was confirmed in all processing conditions. However, it is necessary toinvestigate the cause and countermeasures.
50
45
40
35
30
25
20
15
10
5
0

Particle size of raw powders ≤106 μm

top
side
center
4.9

9.8
14.7
19.6
24.5
Compaction Load / kN
Fig. 4. Porosity on cross section of sinteredporous compacts (Experiment 2).

4．Conclusion
Sintered porous titanium produced by capsule-free HIP method after pre-sintering at a relatively
low temperature with a short holding time, had high porosity. These sintered compacts showed some
improvement of characteristics by homogenization of the green compact, but the center portion of thesintered compact was densified, and its porosity was reduced.
References[1] S. Suzuki, J. L. Martin, H. Utsunomiya, H. Nakajima, Plastic Deformation Processes of Lotus-Type Porous
Metals, Journal of the JSTP, 52(2011), pp. 206-211.
[2] T. Takahashi, K. Matsumaru and K. Ishizaki, Production and Sintering Mechanism of Open-Pore Porous
Materials by Using Hot Isostatic Press (HIP), J. Jpn. Soc. Powder Powder Metallurgy, 55(2008), pp. 170-176.

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PROPOSAL FOR NEW FABRICATION METHOD OF GREEN COMPOSITE BY
USING FRICTION STIR PROCESS
Tairyu Yamashita1, Yuhei Kurabe2, Yukio Miyashita3,*, Yuichi Otsuka3 and Hisashi Hori4
1

Graduate Student, Graduate School of Engineering, Nagaoka University of Technology, Japan
2

3

National Institute of Technology, Ishikawa College, Japan

Nagaoka University of Technology, 1603-1 Kamitomioka, Nagaoka, Niigata, 940-2188, Japan
４

Nippon Light Metal Co., Ltd., Japan)*Corresponding Author: miyayuki@mech.nagaokaut.ac.jp, Tel: +81-258-47-9704

Keywords: Green composite, Friction stir process, Bamboo fiber, HDPE, Tool
Introduction
Composite material by using reinforcement of natural fiber is called green composite. Green
composite has advantages in recyclability and environmental friendly. However, the mechanicalproperty can be easily degraded during the fabrication process by thermal degradation in natural fiberused as reinforcement [1]. In the present study, friction stir process was applied as a lower heat input
process to fabricating green composite material in order to avoid occurrence of thermal degradation in
reinforcement of natural fiber. Bamboo fiber and high-density polyethylene (HDPE) were used asreinforcement and matrix in the present study. A new method to providing the reinforcement into thematrix without pre-machining is proposed.
Experimental Procedures
Materials used for the matrix and the reinforcement was HDPE and bamboo fiber. Physical andmechanical property for HDPE is shown in Table 1. Bamboo fiber used in the present study is shownin Fig.1. Specimen size of HDPE was 100 mm × 40 mm with thickness of 5 mm. Bamboo fibers weremixed with resin to form a cylindrical shape with 4.1 mm in the diameter and 15 mm in the height,
then it was inserted into the tool. Friction stir process was used to mixing the matrix and thereinforcement in order to fabricate green composite in the present study. Preparing groove or hole isnecessary as a pre-machining process in friction stir process in usually. However, those machining
processes decrease productivity and result in increasing production cost. In this study, a new method is
proposed as a friction stir process without pre-machining process. It is expected in the present process
proposed that reinforcement is pushed out from the tool during the process and is broken to be finewith tool traveling. The fine fiber broken was stirred with the matrix by rotating of the tool and thengreen composite is fabricated. Schematic of the tool is shown in Fig.2. The tool has pin with M10
screw. The reinforcement was filled into a hole in the pin. In the friction stir process, tool tilt angle
was 0°and rotating direction of the tool was counterclockwise. Friction stir process experiments werecarried out at conditions of 600 rpm - 15 mm/min (rotating speed - traverse speed), 600 rpm - 45
mm/min and 1200 rpm - 15 mm/min. In the present process, rotating tool was inserted into the matrix
until 4 mm depth from surface of the specimen with plunging speed of 15 mm/min, then the tool
moved with certain traverse speed．
Table 1 Physical and mechanical property of HDPE.

Fig.2 Tool used for friction
Fig.1 Bamboo fiber used. stir process.

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(a)

(b)

(c)

Fig.3 Observation for cross section of the stir zone
applied with (a) 600 rpm – 45 mm/min, (b) 600 rpm –15 mm/min and (c) 1200 rpm – 15 mm/min.

Fig.5 Axial force acting in the tool during the process.

Fig.4 Temperature change during the process.
Step 1

Step 2

Step 3

Step 4

Fig.6 Schematic illustration of fabricatingprocess of green composite by using friction
stir process proposed in the present study.

Result and discussions
Figure 3 shows observation of cross section of the stir zone. According to the figures, thereinforcement was stirred into the matrix in all three conditions applied. According to observation of
the stir zone from the top surface, bamboo fibers were oriented following rotation of the tool. Shape ofthe bamboo fiber bundle changed from 1200 to 350 µm in length and 350 to 15 µm in diameter by thestirring process. Figure 4 shows results of temperature measurement during the process. Temperature
was measured by using a thermocouple attached inside the tool at 1 mm from tip of the tool.
According to Fig. 4, the maximum temperatures measured in each condition were lower than themelting temperature of HDPE and also lower than the temperature inducing thermal degradation of
mechanical property in bamboo fiber. Figure 5 shows change in axial force acting in the tool during
the process. The axial force increased at the beginning of the process, then that decreased and reachedto the constant value. The degradation of the axial force occurred due to softening and flowing of thematerial due to friction stirring. In Fig.5, spring force for pushing the reinforcement out is indicated asa dashed line. In case of that the friction stir process was carried out with condition of 1200 rpm - 15mm/min, the axial force at the steady state was lower than the spring force. On the other hands, inother two conditions, the axial forces at the steady state were higher than the spring force. However,
the reinforcement was pushed out from the tool into the matrix, even the axial force was higher than
the spring force. According to the results shown in above, friction stir process conducted in the present
study was considered as schematically shown in Fig. 6. The material was started to move by friction
stirring and then flowed into the tool and broke the fiber. The fiber was flowed out from the tool andmixed in the matrix with traverse of the tool, then resulted in fabricating green composite.
Summary
Green composite of HDPE matrix and bamboo fiber reinforcement was successfully produced by thefriction stir process proposed in the present study without pre-machining process. Bamboo fiber
bundle was broken and mixed with the matrix by flowing of HDPE during the process. The maximum
temperature observed during the process was lower than the melting temperature and the temperatureinducing thermal degradation of mechanical property in bamboo fiber.
Reference[1] S. Ochi, H. Takagi and R. Niki, Mechanical Properties of Heat-Treated Natural Fibers, J. Soc. Mat.
Sci., Japan, 51, 10 (2002) pp.1164-1168.

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INFLUENCE OF GRAIN SIZE ON PROCESS AFFECTED ZONE IN MICROPIERCING
Tomomi Shiratori1*, Shizuka Nakano2, Yohei Suzuki1, Masahito Katoh2,
Naoko Sato2, Takafumi Komatsu1, Ming Yang3.
1

Komatsuseiki Kosakusho Co. LTD., 2 National Institute of Advanced Industrial Science and Technology,
3
Tokyo Metropolitan University*shiratori@komatsuseiki.co.jp

Keywords: Piercing, Punching, Fine-grain, Sheared surface, EBSD
Control technologies of material grain size had been developed for austenitic stainless steel
SUS304. It is well known that the material has a tendency to increase the tensile strength anddecrease the elongation for smaller grain size. There have been a number of studies on micropiercing and basic stamping characteristics for austenitic stainless steel SUS304 with fine-grained
micro structure [1,2]. There are few investigations on the correlation of stability of sheared surface andgrain size in micro piercing [3]. We have been investigating that stability of starting point of fractured
surface becoming smaller when choosing the average grain size of 3.0μm [3]. However, themechanism has not yet been elucidated. Purpose of this study is that understanding the effects of grain
size on process affected zone from middle step of micro piercing.
To reveal phenomena, we selected the experimental conditions: punch diameter of φ0.08mm,
material thickness of 0.10mm and as a result, aspect ratio of 1.25, clearance between punch and die are
set by 2.5μm, which is 2.5% of the material thickness. Average grain sizes of material are 7.5μm fornormal material, 1.5μm and 3.0μm for fine-grained material. Mechanical properties are shown in Table
1. Desk top stamping unit (SSI-01) is produced with BISAIKAKOH KENKYUSHO shown in figure
1 (a), and the motions are controlled by servo motor and the maximum stamping force is 10kN. Micropiercing die was produced for the experiment shown in Figure 1 (b). The die has a nano-meter
accuracy stage embedded in under of the die tool. It is possible to control the relative position
between punch and die in nano-meter scale to obtain concentric position. Process affected zones
of middle step of micro holes were observed by FE-SEM and measured by EBSD (Electron Back
Scattered Diffraction). EBSD investigation area of middle step of hole is 65μm (x axis) and 120μm (yaxis) with 0.15μm pitch analysis. Process affected zones are investigated by EBSD phase map and
EBSD KAM (Kernel Average Misorientation) map.
Figure 2 shows the SEM BSE image of cross section of middle step of micro piercing.
Process affected zones are presented along a line from punch edge to die edge for each material.
Figure 3 (a) shows the EBSD Phase maps from grain sizes 1.5μm to 7.5μm. During micro piercing,austenitic stainless SUS304 is able to transform the phase to strain-induced martensitic. Whentransform is progressing, the phase changes from face-centered cubic (fcc, γ phase) tobody-centered cubic (bcc, α phase). Martensitic phase in grain size 7.5μm looks hackly and slanted
to 45 degrees direction, long and wide grains are appeared. Slip line of 45 degrees direction is inducedby shear strain. Martensitic phase in grain size 3.0μm is smoother in comparison to grain size 7.5μm.
It is considered that grain size is strongly related to strain-induced martensitic transformation.
Martensitic phase in grain size 1.5μm is more expansive and smoother than grain size 3.0μm. It relates
tensile strength. When the grain size is smaller, strain-induced martensitic is becoming smooth andexpand expansive. Figure 3 (b) shows the EBSD KAM map from grain size 1.5μm to 7.5μm.
Misorientation angle is investigated in this map and this angle is known to correspond to plastic strain[4,5]. Misorientation angle 5 degrees area (Color : Red) for grain size 7.5μm expand from punch edge
to die edge. 45 degrees slip line and 5 degrees misorientaion areas are clearly correlated. It is said thatmisorientation angle 5 degrees area is related to local plastic strain. Misorientation angle 5 degrees
area in grain size 3.0μm looks more expansive than 7.5μm. Misorientation angle 5 degrees area ingrain size 1.5μm is expender than 3.0μm.Area for misorientation angle 5 degrees depends on grain
size and tensile strength. If grain sizes are smaller, area for misorientation angle 5 degrees becomes
more expansive with stronger tensile strength.
Investigated results show that the process affected zones in micro piercing of middle step ofPage 21

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

piercing, strain-induced martensitic phase and misorientation angle are strongly related to grain size.
Table and Figure
Table 1 Characteristics of material
Grain Size [μm]
1.5
3.0

7.5

Tensile strength [MPa]

875

845

803

0.2%Proof stress [MPa]

599

504

433

Elongation [%]

45.8

49.5

56.6

n Value

0.31

0.37

0.43

Vickers hardness [HV]

261

227

191

Grain size 1.5μm

SEM BSE image
Punch
penetration

Grain
Size
(a)
EBSD
phase
map

(a) Desk top press unit(b) Micro piercing tool

Figure 1 Experiment tools

1.5μm

3.0μm

7.5μm

Phase
α phaseGrain size 7.5μm

Grain size 3.0μm

γ PhaseFigure 2 SEM BSE image

(b)
EBSD
KAM
map
Misorientation
angle
degree05Figure 3 EBSD Phase map and KAM map

References[1] T. Komatsu, T. Matsumura & S. Torizuka, Effect of grain size in stainless steel on cutting performance inmicro-scale cutting, Int. J. Autom. Tech., 5-3(2011), pp. 334-341.
[2] T. Komatsu, H. Kobayashi, S. Torizuka, & S. Nagayama, Micro hole piercing for ultra fine grained steel, Int.
J. Autom. Tech., 6-6(2011), pp. 802-808.
[3] T. Shiratori, S. Nakano, Y. Suzuki, M. Katoh, M. Yang, T. Komatsu, Influence of grain size on sheared
surface in micropiercing, j. JSTP, 56-652(2015), pp. 401-506[4] M. Kamaya, A. J. Wilkinson, J. M. Titchmarsh, Measurement of plastic strain of polycrystalline material byelectron backscatter diffraction, Nucl. Eng. Des., 235 (2005), pp. 713–725.
[5] Masayuki Kamaya, Measurement of local plastic strain distribution of stainless steel by electron backscatter
diffraction, Materials characterization, 60 (2009),pp. 125-132.

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REPLICATION OF ANTIREFLECTION STRUCTURE FOR SOLAR CELLS BY
LASER-ASSISTED IMPRINTING
Keisuke Nagato1,2*, Yuki Yajima1, Ken Takahashi1, Tetsuya Hamaguchi1, Masayuki Nakao1
1

The University of Tokyo, 2JST PRESTO*nagato@hnl.t.u-tokyo.ac.jp

Keywords: Replication, Solar cells, Laser-assisted imprinting
Nanoimprinting is a promising candidate, in which large-area polymer surfaces are replicated byusing nanostructured mold with high throughput and low cost. However, in the conventional thermal
nanoimprinting, entire of polymer and mold should be heated and cooled across the glass transition
temperature of the polymer. Therefore, the cycle time is limited [1]. To date, we have proposed a laserassisted replication technology as a breakthrough the cycle time limitation [2]. The laser is irradiated tothe surface of the mold and only the surface of the polymer is heated and filled into the nanostructureof the mold, and then the replicated polymer is cooled by heat conduction to the inside of the mold andthe polymer immediately after the laser irradiation.
The laser-assisted roller imprinting is applied to the roller imprint method [3]. A polymer film ispressed by a glass roller and a Ni roll mold. The laser is irradiated to the Ni mold surface through theglass roller and the polymer film. The surface of the mold is directly heated by the laser and thenanostructure is thermally replicated to the polymer surface. The laser is scanned onto the contactedarea and the continuous replication of large-area polymer is performed.
On the other hand, the surface of solar cell reflects the light with Fresnel reflectance principle [4].
We propose that the light-incorporation efficiency is increased with covering a polymer film with anantireflection structure. The increase of efficiency leads to the enhancement of the actual generation
efficiency.
A subwavelength-sized antireflection structure was replicated on a 75-m-thick polymethylmethacrylate (PMMA) film by the laser-assisted roller imprinting. The Ni mold of the structure wasfabricated by electroplating using an anodic-alumina nanohole array [5]. The average pitch and depthof the inverted-conical structure of Ni mold were about 150 and 300 nm, respectively. The image
inserted in Fig. 1 shows a scanning electron microscopy (SEM) image of the replicated conical structure.
We measured the reflectivity of the imprinted surface using a spectrophotometer (CM-2600d, Konica
Minolta Sensing, Inc.). Figure 1 shows the reflectivities of the flat PMMA surface and the imprinted
surface. The reflectivity of the imprinted surface, about 0.5%, was drastically lower than that of the flat
surface, about 4%, over the entire range of visible wavelength.
The enhancement of solar cells using the antireflection-structured film was demonstrated. The solar
cell used was LR0GC02 from Sharp Co. (40×60 mm2). When measuring the property, the reverse sideof the PMMA film was attached on the cell using glycerin with the same refractive index as PMMA
(1.4―1.5). A fluorescent light, FL20SSW/18 from TOSHIBA CO., was set at 110 mm above the solar

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cells. As shown in Fig. 2, the maximum generation power was increased for 4% when the antireflectionstructured film was attached. This result can be simply discussed with the effect of the antireflection
properties of the surface of the attached film. Because the antireflection structure has an advantage thatthe reflectivity is low throughout the various incident angle of the light, the attachment of theantireflection structure on solar cells leads to the enhancement of the power generation.
In summary, we demonstrated the replication of the antireflection structure on PMMA film using
laser-assisted roller imprinting technique and enhancement of the power generation of solar cells. The
increase, 4%, of the generation corresponded to the decrease of the reflectivity of the PMMA film. The
laser-assisted roller imprinting is a promising method of the mass fabrication of antireflection film ofsolar cells because of the advantages of high throughput and large area.

Figure 1. Reflectivity of flat and antireflection-structured PMMA films. Inserted figure shows a tilted-view
SEM image of antireflection structure.

Figure 2. Current-voltage and power-voltage properties of solar cells covered by flat and antireflectionstructured PMMA films.

References[1] K. Takahashi, K. Nagato, T. Hamaguchi, M. Nakao, High-speed replication of light-extraction structure withthermal roller nanoimprinting, Microelectron. Eng. in press.
[2] K. Nagato, K. Takahashi, T. Sato, J. Choi, T. Hamaguchi, M. Nakao, Laser-assisted replication of large-areananostructures, J. Mater. Process. Technol. 214 (2014), pp.2444-2449.
[3] K. Nagato, K. Takahashi, M. Nakao, Thermal roller nanoimprinting –multilayered structures, laser-assisted
replication–, 14th International Conference on Nanoimprint and Nanoprint Technology (NNT).
[4] S.-J. Liu and C.-T. Liaw, Fast fabrication of nano-structured anti-reflection layers for enhancement of solar
cells performance using plasma sputtering and infrared assisted roller embossing techniques, Opt. Express 20(2012), pp. 5143-5150.
[5] T. Yanagishita, K. Yasui, T. Kondo, Y. Kawaomoto, K. Nishio, H. Masuda, Antireflection polymer surface
using anodic alumina molds with tapered holes, Chem. Lett. 36 (2007), pp.530-531.

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SELECTIVE PATTERN-HARDENING OF AISI-SUS420/SUS430
MOLDS FOR MICRO-TEXTURING
Tatsuhiko Aizawa1*, Istiroyah2, Yoshio Sugita3
1

Shibaura Institute of Technology, 2Brawijaya University, 3YS-Electric Industry, Co. Ltd.
*taizawa@sic.shibaura-it.ac.jp.

Keywords: Micro-texturing, Plasma nitriding, Stainless steels, Solid solution hardening
INTRODUCTION
Micro-texturing has grown to be one of the most important die-and-mold technologies to improve
the surface properties and functions of metallic and plastic products. The conventional micro-millingand micro-EDM [1, 2] require for huge amount of leading time to fabricate the micro-textured molds
and dies. Non-traditional micro-texturing method is proposed with aid of the high density plasma
nitriding process [3-5]. Its key principle in manufacturing lies in the selective pattern-hardening ofmolds, where the unprinted part of mold surface is selectively nitrided and hardened with other printed
parts remaining un-nitrided. In the present study, this key technology is experimentally demonstrated
by using the patterned AISI-SUS420 and SUS430 specimens. Their un-masked surfaces are nitrided
to have the hardness higher than 1200 HV; while, the hardness of masked surfaces, which are notessentially nitrided, remain as they are before nitriding. Besides the hardness testing, SEM and EDX
are utilized to discuss the relationship between the selective pattern-hardening and the inner nitriding
process.
EXPERIMENTAL PROCEDURE
This plasma-nitriding assisted micro-texturing method consists of three processes as shown in Fig.
1. At first, an original mask-pattern is printed on the surface of mold-die by using the dispenser.
Selection of plastic primer inks to endure in the nitriding condition becomes essential to make micropatterning onto the surface of tool steel and stainless steel substrates. In second, the substrate materials
are plasma-nitrided to have sufficient hardness without change in the contents of chromium and ironelements and with less damage onto the substrate material surface. Finally, this nitrided surface ismechanically polished to remove the un-nitrided patterns and the residual inks and to form a negative
micro-texture against the initially printed micro-pattern.
1. Ink Jet Printing

2. Plasma Nitriding

3. Sand Blasting

4. Injection Molding

Dispenser/Digital control/Ink-jet printing/Spatial
resolution
1 mm

High Density
Plasma Nitriding
System/RF-plasma/DC-plasma/Heating/Low temperaturenitriding
Sand-Blasting/Controlling items- Gas flow rate
- Sand materials- Sand diameter- SettingFig. 1: Plasma nitriding assisted micro-texturing method.
EXPERIMENTAL RESULTS and DISCUSSION
In order to demonstrate that the original masking pattern should be homogeneously transformed into

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

the hardness profile pattern, a single square region was only left unprinted at the center of AISI-SUS420
type J2 specimen as shown in Figure 2 a). That is, the whole surface except for a square region at thecenter of specimen was printed as a mask. Figure 2b) compares the hardness profiles measured bothin the lateral and longitudinal directions across the mask. The applied indentation load was constant, 1
N or 100 g in weight. Less significant difference was seen in both hardness profiles; i.e. the hardness
in the un-masked regions is 1400 Hv, and, it remains to be 200 Hv in the masked region. The steep
change of hardness across the edge of masks reveals that the unmasked regions are nitrided to havemuch higher hardness than matrix hardness of AISI-SUS420 stainless steels. On the other, the masked
regions are free from infiltration of nitrogen atoms into matrix. Hence, the hardness map in the nitrided
specimen corresponds to the masking pattern. Hence, the masked regions are easy to be removed anddug into the depth by the sand blasting; the un-masked regions have sufficient hardness not to bemechanically removed. Owing to this selective pattern hardening, the micro-textured AISI420/430
stainless steel molds are fabricated by the procedures in Fig. 1 even without use of mechanical millingor micro-EDM.

X

a)

1600

Hardness / HV

MaskedRegion

1400
1200

Y

X

b)
HV : 1400(Ave)

1000
800
600
400

HV : 250(Ave)

200
0

Y

軸ラベル

10 mm

Fig. 2: Selective pattern hardening via the high density plasma nitriding onto the un-masked surfaces.
a) Masked specimen, and, b) Hardness profile in the x- and y-directions on the specimen.
CONCLUSION
Non-traditional micro-texturing method is proposed to fabricate the micro-textured molds and diesfor injection molding of micro-patterned plastic parts and for stamping to micro-patterned metallic parts.
The final micro-patterns on the products are controllable by the initial CAD data for ink-jet printing onto
the surface of stainless steel molds and dies. This initial micro-pattern is transformed by the lowtemperature plasma nitriding to the hardness profile. The patterned regions have lower hardness ornearly the same hardness of matrix. The bare surface except for this initial patterns has higher hardness
than 1400 Hv. This hardness profile results in the microthe sand-blasting. Then, the initial micro-pattern turns to be a micro-texture into the molds and dies.
References[1] B. Denkena, J. Koehler, J. Laestner, “Efficient machining of micro-dimples for friction reduction,” Proc. 7th
ICOMM, 2012; 85-89.
[2] Y. Jiang, W.S. Zhao, X.M. Kang, L. Gu, “Adaptive control for micro-hole EDM process with wavelet transform
detecting method,” Proc. 6th ICOMM, 2011; 207-211.
[3] T. Aizawa, T. Yamaguchi, “High-density plasma nitriding assisted micro-texturing onto martensitic stainlesssteel mold-die,” Proc. IWMF2014, 2014; 31-37.
[4] T. Katoh, T. Aizawa, T. Yamaguchi, “Plasma assisted nitriding for micro-texturing onto martensitic stainlesssteels,” Manufacturing Review. 2 (2015).
[5] T. Aizawa, T. Yamaguchi, “Plasma nitriding assisted micro-texturing into martensitic stainless steel molds forinjection molding,” Proc. 4M/ICOMM (2015) 454-459.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

EVALUATION OF MAGNETIZATION FOR CANTILEVER SAMPLE USING MEMS
DEVICE
Takuya Maetani1, Junpei Sakurai1, Mizue Mizoshiri1, Seiichi Hata1,*
1

Graduate School of Nagoya University, Japan*Corresponding Author: hata@mech.nagoya-u.ac.jp

Keywords: MEMS, magnetic material, relative permeability, cantilever, combinatorial method
INTRODUCTION
The combinatorial method [1] is a useful method for the searching of the novel materials or optimize
the composition of alloy materials. In this method, it is studied how to fabricate the library and how toevaluate the characteristics. The library that has different compositions or fabrication conditions could
synthesize by using sputtering or laser molecular beam epitaxy with the moving masks to limit thedeposition area. In the recent work, we proposed the Combinatorial New Facing Targets Sputteringand the Combinatorial Arc Plasma Deposition [2]. These thin film deposition methods could fabricatecomposition gradation library. There are several high throughput evaluation methods such as fatiguehydrogen storage characteristics [3].
Magnetostrictive materials have gained attention because of the features that are high power, fast
response, and contactless actuation. Magnetostriction and relative permeability are importantcharacteristics to determine the performance of the sensors and actuators with magnetostrictive
materials. However, conventional evaluation method of the magnetostriction and relative permeability
took a lot of time. In this research, we proposed the high throughput evaluation method of relative
permeability using with the MEMS device.
FABRICATION PROCESSThe MEMS device for evaluation of relative permeability consists of the cantilever samples,vibrating electrode and pickup coil as shown in figure 1. The evaluation substrate has vibrating
electrode and pickup coil. The sample substrate and the evaluation substrate are fabricated differently,
and then bonded these two parts.
Figure 2 shows the fabricated evaluation substrate. The detail of the fabrication process is shown infigure 3. The fabrication process of evaluation substrate starts with cutting Si substrate in dimension of
30 mm ×30 mm. The Si substrates are ultrasonically rinsed with acetone, ethanol and DI water andbaked on a hot plate. Then a bottom electrode layer of Au/Cr is sputtered on the patterned negative
photoresist (ZPN-1150, Zeon Inc.) using RF magnetron sputtering machine. This is shown in figure3(a). Next, lift-off process is done with acetone using an ultrasonic shaker for one minute. This isshown in figure 3(b). Next, polyimide based positive photoresist (PW-1500, Toray Industries Inc.) ispatterned as an insulating layer. Then polyimide is nitride for the 340°C in the nitride atmosphere asshown in figure 3(c). Next, top electrodes of Au/Cr are sputtered on the patterned negative photoresist
using RF magnetron sputtering machine. This is shown in figure 3(d). Next, lift-off process is donewith acetone using an ultrasonic shaker for one minute. Then, acryl based negative photoresist(THB-151N, JSR Inc.) is patterned as the mold for the electroforming as shown in figure 3(e). Next,
25 μm thick Au is electroformed as shown in figure 3(f). The condition of current density is 0.8A/dm2.
Next, 5 μm thick In is electroformed as shown in figure 3(g). The condition of current density is sameas Au electroforming. Finally, electroforming mold is removed using THB-S2 (JSR Inc.) on a hot plate(80 °C for 10 minutes) as shown in figure 3(h).
EXPEIMENT AND RESULT
To confirm the principle of measurement using evaluation substrate, magnetization of magnetic
material is evaluated. Figure 4 shows the experimental setup. The sample is vibrated 20 μm amplitude
by using a piezo actuator in the proximity of the pickup coil in static magnetic field, which is appliedby using a Helmholtz coil. Then, induced electromotive force is produced to the pickup coil.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

30 mm

1 mm

Figure 1. The overview of the device
Au/Cr

Figure 2. Photo of evaluation substrate
Au/Cr

Insulating layer

(a)

(b)

Electroformed Au

(c)

(d)

Electroformed In(f)(h)
Figure 3. The fabrication step of evaluation substrate

Helmholtz coil

Coil

Evaluation substrate
SamplePiezo-electric
deviceVoltageVibration follower

Voltage
amplifier

Voltage
Amplifier(×10)

Function
generator

Figure 4. The experimental setup

Oscilloscope

0.8
0.6
0.4
0.2
0.0
-0.2
-0.4
-0.6
-0.8

VSM

This method

50
40
30
20
-1000

Output voltage [mV]Power supply

Density of magnetic flux [T]

(e)

10 mm

-500
0
500 1000
Magnetic field [A/m]Figure 5. The magnetization curves measured byusing with VSM and this method

To compare the measurement result of magnetization, evaluation substrate and vibrating sample
magnetometer (VSM) were used. The evaluation sample (Fe32.0Ni55.3Cr12.7) is fabricated by using New
Facing Targets Sputtering on the 10 mm ×10 mm glass substrate, then annealed at 773 K in thevacuum. Figure 5 shows the measurement result. The result of using evaluation substrate is similar to
the result of using VSM. This result suggests the possibility of high throughput measurement formagnetism. However, it could not transfer the voltage to the magnetism in this method, it is necessary
compared the output voltage to the reference material such as Ni.
CONCLUSIONSIn this research, the evaluation method of magnetism is studied. The evaluation substrate is fabricatedsuccessfully. The magnetization of Fe32.0Ni55.3Cr12.7 was evaluated by using VSM, which is aconventional measurement method, and evaluation substrate. The magnetization curves measured by
two methods are similar. There is possibility that this measurement method could apply to the
combinatorial material searching method.
References[1]X.-D. Xiang, X. –D. Sun, G. Briceho, Y. Lou, K.-A Wang, H. Chang, W. G. Wallace Freedman, S.-W. Chen, P.
G. Schultz, A Combinatorial Approach to Materials Discovery, Science 268 (1995), pp.1738-1740[2]S. Hata, R. Yamauchi, J. Sakurai and A.Shimokohbe, Combinatorial Arc Plasma Deposition of Thin Films,
Jpn. J. Appl. Phys. 45 (2006), pp.2708-2713[3]A. Ludwig, J. Cao, A. Savan, M. Ehmann, High-throughput characterization of hydrogenstorage materials using thin films on micromachined Si substrates, J. Alloys andCompounds, 446-447 (2007), pp.516-521

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

FABRICATION OF MICRO-CHANNEL FROM THE POWDER BY PULSEELECTRIC CURRENT SINTERING METHOD AND MARKED GRAFITE DIE
Satoshi Kishimoto1and Yutaka Kagawa2
1

National Institute for Materials Science, 2National Institute for Materials Science*KISHIMOTO.Satoshi@nims.go.jp

Keywords: Spark plasma sintering method, micro-channel, micro-pattern, Ion milling, Low
melting point glass.

There are a lot of investigations to fabricate the micro-channel for sensors. To fabricate themicro-channel, the electron beam lithography [1], photo-lithography [2], micro-milling andnano-imprint [3] have been utilized. However, these methods need expensive equipment and longprocess. In this study, a rapid fabricating method using pulse electric current sintering (Spark plasma
sintering: SPS) method has been developed.
One surface of a graphite punch (10mm diameter) was polished and patterns of micro-channel were
fabricated on the polished surface by an ion milling. Using this punch, low melting point glass powder
is sintered by SPS (Sumitomo, SPS-515). The surface of the punch with micro-channel and fabricatedspecimen was observed by a scanning electron microscope (SEM), and an optical micro-scope.
Figure 1 (a) shows a fabrication process of graphite punches by ion milling. A mask of same metal
wires of 100micro-meter diameter is placed on the polished surface of graphite punch. Argon ion
shower is exposed unto the mask. After removing the mask, micro-pattern is fabricated. Figure 1 (b)shows a substrate with micro pattern fabrication process, low melting point glass (ZnO･Bi2O3･B2O3)
and Al powder was sintered by pulsed electric current sintering (temperature: 460oC, holding time: 10
minutes).
(a)(b)

Pressure
Graphite punch
Die

Micro-pattern wire
Ar ion
Mask

Graphite Die

Electric current

Micro Pattern

Specimen

Figure 1 Micro-pattern fabrication process, (a): Graphite punch with micro-pattern fabrication process
and (b): Substrate with micro-pattern fabrication process.

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Figures 2 show the photographs of the scanning electron microscope images of graphite punch oflower magnification: (a) and the higher magnification: (b). Figures 3 show the photographs of thescanning electron microscope images of sintered low melting point glass surfaces. Figure 3 (a) showsa lower magnification image and (b) shows a higher magnification image. These photographs indicate
a micro-cannel patterns with 100 micro meter width can be fabricated by this method.

(b)

(a)

Ridge

Ridge

50µm

200µm

Figure 2 Scanning electron microscope images of graphite punch, (a) : lower magnification and (b) :higher magnification.

(a)

(b)
GrooveGroove

200µm

50µm

Figure 3 Scanning electron microscope image of sintered low melting point glass surface, (a) : lower
magnification and (b) : higher magnification.

References[1] Kishimoto, S., Egashira, M., and Shinya, N., “Micro-creep deformation measurement by a Moiré method
using electron beam lithography and electron beam scan”, Opt. Eng., Vol. 32, (1993), pp. 522-526.
[2] Nam, Y., Kim, M. and Kim T., “Pneumatically controlled multi-level microchannel for separation andaxtraction of microparticles”, Sensors and Actuators B, Vol. 190, (2014), pp. 86-92.
[3] Tang, M. J., Xie, H. M., Zhu, J. G., Li, X.J., and Li, Y. J., “Study of moiré grating fabrication on metal
samples using nanoimprint lithography”, Opt. Express., Vol. 20, (2012), pp. 2942-2955.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

NUMERICAL STUDY ON EFFECTS OF MATERIAL INHOMOGENEITY ANDDRAWING OF METAL FOILS
Tsuyoshi Furushima1*, Kenta Takahashi1, Ken-ichi Manabe11Dept. Mechanical Engineering, Tokyo Metropolitan Univ., Japan*furushima-tsuyoshi@tmu.ac.jp

Keywords: Material inhomogeneity, Micro deep drawing, Necking, Wrinkling
1. Introduction
Micro metal forming with metal foils has been focused on in technical applications such as
electronic devices, medical equipment, sensor technologies, and optoelectronics. The miniaturization ofcomponents takes place problem of size effects due to grain size and surface roughness on
inhomogeneous deformation. In particular, the surface roughening of metal foils may affect not only
local problem such as accuracy of products and frictional condition but also global deformation behavior
such as necking behavior [1]. Therefore, prediction of inhomogeneous deformation such as free surface
roughening, friction and necking behaviors is very important in micro metal forming for metal foils withultra thin thickness. From these backgrounds, we have proposed mesoscopical inhomogeneous finite
element method considering mesoscopical material inhomogeneity for prediction of free surface
roughening, necking and wrinkling in micro deep drawing process.
2. Mesoscopical inhomogeneous finite element method
In this model, grain size and variation in flow stress of each grain as aspect of mesoscopic material
information were considered. A commercial FE code of LS-DYNA ver. 971 was used. Figure 1 shows
the mesoscopical inhomogeneous FE model of micro deep drawing. To model grain shape, Voronoi
Tessellation was utilized in this model [2]. In the micro deep drawing, punch diameter of 944μm andthickness of 50μm are used. As boundary conditions of micro deep drawing, quarter model were given.
For each grain, different flow stress was input. The modified Swift's flow stress equation as shown in
Eq. (1) considering material inhomogeneous parameter αi was used for each grain.
(1) eq   i F  0   eq n (i=1,2,...k)The probability distribution of the αi value in the FE model is divided into k classes formesoscopical inhomogeneous FE material modeling. In this model, the class k was 7. Table 1 and 2
show the mechanical properties and αi values and their standard deviations of blank used in this analysis.
In this study, a homogeneous model and inhomogeneous models with grain size of 2, 10 and 20μm were
given to investigate the effects of inhomogeneity and grain size on micro deep drawability. In addition,
the effects of drawing ratio DR=1.80, 1.91, 2.01 and 2.12 and blank holder force on micro deep
drawability were investigated numerically.
Punch(Shell/Rigid)

Blank holder(Shell/Rigid)(a) Homogeneous model(b) d=2μm

Inhomogeneous blank(Solid/Isotropic elasto-plastic)z
Die(Shell/Rigid)(c) d=10μm(d) d=20μmxy(e) Micro deep drawing model

Figure 1 Mesoscopical inhomogeneous FE model of micro deep drawing.

Page 31The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015’s modulus

E /GPa197

Table 1 Mechanical properties of blank in FE analysis
Poisson’s
Yield stressK-valuen-valueratio y /MPa
K /MPa
n
0.3

526

1570

Initialequivalent strain o

0.29

0.03

Table 2 α-values and their standard deviations


1

2

3

4

5

6

7



Condition

0.7

0.8

0.9

1.0

1.1

1.2

1.3

0.11

3. Results and discussion
Figure 2 shows the comparison of strain distribution in thickness direction between homogeneousmodel and inhomogeneous model with grain size of 2 and 20μm under condition of DR=1.91 and
BHF=48N. The necking did not occur under condition of grain size of 2μm and homogeneous model.
However, the necking caused by local deformation occurs at cup radius under grain size of 20μm. Figure
3 shows the effect of grain size on limiting drawing ratio (LDR) diagram obtained from various DR and
BHF. As a result, it is found that the LDR and wrinkling limit decrease with increasing grain size. Thesetendency obtained from the inhomogeneous simulation are in good agreement with that of previous
experiment [3]. From these results, we clarify the importance of the simulation considering material
inhomogeneity in micro metal forming.
(a)

(b)

(a) Homogeneous model(c)(b) d=2μm(c) d=20μm

Figure 2 Comparison of strain distribution in thickness direction under fracture limit condition.

Blank holder force /N

70.0

Fracture limit

60.0
50.0
40.0

30.0

d=2mm d=10mm

Homogeneous(=0)

20.0
10.0

0.0

Wrinkling limit
1.70

1.80

d=20mm
1.90

2.00

2.10

2.20

Drawing ratio

Figure 3 Effect of grain size on limiting drawing ratio (LDR) diagram obtained from various DR and BHF.

4. Conclusion
In this study, we have proposed mesoscopical inhomogeneous finite element method considering
mesoscopical material inhomogeneity for prediction of free surface roughening, necking and wrinklingin micro deep drawing process. As a result, the LDR and wrinkling limit decrease with increasing grain
size. Therefore, the importance of the simulation considering material inhomogeneity in micro metal
forming was clarified.
References[1] T. Furushima, H. Tsunezaki, K. Manabe, S. Alexandrov, Ductile fracture and free surface roughening behaviorsof pure copper foils for micro/meso-scale forming, Int. J. Mach. Tools. Manuf., 76 (2014), pp. 34-48.
[2] T. Furushima, K. Manabe, S. Alexandrov, Mesoscopical inhomogeneous material modeling for micro metal
forming, AIP Conf. Proc., 1618 (2014), pp. 331-334.
[3] F. Vollertsen, Effect of the deep drawing diagram in micro forming, Prod. Eng. Res. Devel., 6 (2012), pp11-18.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

TECHNICAL MILESTONE TO APPLY SELF-HEALING CERAMICS
Wataru NAKAO1,2*
1

Institute of Advanced Sciences, Yokohama National University, 79-5, Tokiwadai, Hodogaya-ku, Yokohama,240-8501, Japan
2
Faculty of Engineering, Yokohama National University*wnakao@ynu.ac.jp

Keywords: High Temperature, Structural Materials, Jet Engine, Ceramics based Composite, Car
EngineSelf-healing materials are designed by using novel structural material concept, in which the chemical
reaction is exploited as a function like creature. The novel concept has received large interest; many
researchers have made a large effort to develop the advanced self-healing materials, e.g., self-healing
polymer [1], self-healing concrete [2], and self-healing alloys [3].
Fiber reinforced self-healing ceramic (shFRC) [4] is one of attractive self-healing materials and has
been anticipated to be the candidate materials for turbine blade. As the turbine blades are employed inhigh temperatures from 600 oC to 1500 oC, the shFRCs should be developed by adjusting the chemical
reaction to be suitable to the service conditions. Moreover, more complex structure compared withordinary composites must be controlled to maximize the chemical reaction to the strength recovery
effects. Therefore, the actualization of turbine blades using shFRC will not be attained without theadequate strategy to solve the retained technical issues.
Considering the situation, the present author established the roadmap including to the technicalmilestones, as shown in Fig.1, in which the technical readiness level [5] is employed as the verticalaxis. The first stage is to actualize the concept of fiber-reinforced self-healing ceramics as the realmaterial. Furthermore, in the stage, the development of the advanced shFRC which satisfied therequired performances for the application, i.e., the material design concept of shFRC is established.
The second stage is to summarize the process factors of shFRC which exhibit the excellent
performance with small scatter. In order to reach the goal of the second stage, it is necessary to
determine the basic processing of shFRC, and to investigate the shFRC performance with the variationof the processing parameters. The subsequent stages are evaluation the performance of the shFRC
components, and to enhance the accuracy of the processing.

Figure 1 Technical readiness level roadmap of shFRC turbine blade for jet engines

Recently the present author and co-workers succeeded to the first milestone in the first stage, where
the first mile stone was defined that the self-healing rate and high temperature strength are satisfied to
the required performances of the middle pressure turbine blade used at temperatures from 1000 oC to

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1200 oC. Some critical material design strategies to enhance self-healing ability and mechanicalproperties were found through the development of shFRC with the required performance. One of thestrategies is employment of the heterogeneous interlayer. 50wt% TiSi2-SiC heterogeneous interlayer
leads to the adequate self-healing behavior, by which the developed shFRC can heal completely the
introduced crack at 1000 oC for 10 min. The healing rate is superior to the every crack propagation
except critical fracture expected in the middle pressure turbine application. Also it was found thatpre-stressed fiber bundle coating enhances the total mechanical performance of shFRC, because thedefects in fiber bundle became the critical fracture resource in the proto-type shFRC. Employing theprocess, the stiff interlayer, such as 50wt% TiSi2-SiC heterogeneous interlayer, can be used. As a result,it was succeeded that the new type of shFRC having 7 times higher strength than that of proto-type
shFRC was developed. As the strength is higher than the design stress of Ni superalloy turbine blade,the developed it is confirmed that the developed shFRC has the adequate performance to the middle
pressure turbine blade.
References[1] S. R. White, N. R. Sottos, P. H. Geubelle, J. S. Moore, M. R. Kessler, S. R. Sriram, E. N. Brown and S.
Viswanathan, Autonomic healing of polymer composite, Nature, 409 (2001), pp. 794-797.
[2] A. García, E. Schlangen, M. Van De Ven, and D. Van Vliet, Induction heating of mastic containing
conductive fibers and fillers, Materials and Structures/Materiaux et Constructions, 44[2], 2011, pp. 499-508.
[3] S. M. He, N. H. van Dijk, M. Paladugu, H. Schut, J. Kohlbrecher, F. D. Tichelaar, and S. van der Zwaag, In
situ determination of aging precipitation in deformed Fe-Cu and Fe-Cu-B-N alloys by time-resolved
small-angle neutron scattering, Physical Review B, 82 (2010), pp.174111-174129[4] W. Nakao, D. Maruoka, S. Ozaki, M. Nanko, and T. Osada, Advanced Ceramic Composite Using
Self-Healing and Fiber-Reinforcement, Mechanical Properties and Performance of Engineering Ceramics andComposites IX, (2015), pp. 187-193.
[5] Department of Defence, Technology Readiness Assessment Deskbook, https:// www.skatelescope.org /public/
2011-11-18_WBS-SOW_Development_Reference_Documents/DoD_TRA_July_2009_Read_Version.pdf

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

BLANKING VIBRATION AND NOISE OF DIFFERENT TENSILE STRENGTH
MATERIALS IN SCREW DRIVE SERVO PRESS BLANKING
Tomohiro MURAKAMI1,*, Ming YANG2
1

Polytecnich University, 2Tokyo Metropolitan University*murakami@uitec.ac.jp

Keywords: tensile strength, blanking noise, break through, screw drive servo press
1. INTRODUCTION
There is a noise reduction effect in low speed blanking by screw drive servo press, especially in lowstrength metals [1]. But it is said that the noise reduction effect is not admitted by stainless steel and hightensile steel [2]. A blanking noise is occurred to blanking vibration [3]. And a blanking vibration isrelated to blanking tool life. Therefore it is necessary to grasp how the materials influences blanking
noise and blanking vibration. In this study, we investigated the relationship between the tensile strengthof materials, the noise and blanking vibration, in screw drive servo press blanking.
2. Experimental procedure
We used a screw drive servo press (KOMATSU HCP3000: 800kN) in this experiment. The size ofblanking is 40mm×40mm, Corner radius is 1mm, and die clearance is 5% to the thickness of 1mm. We
defined the blanking speed as the punch speed just before breaking. Under the blanking speed
V=15mm/s as the usual blanking speed and V=0.5mm/s as the low-speed blanking, we compared copper(C1100), mild steel (SPC), stainless steel (SUS304) and high tensile steel (SPFC). Specimen’s thickness
is 1mm and the size is 80mm×80mm.Blanking noise and the location of the punch was measured bymeasurement devices indicated in Figure 1.
3. Results
Figure 2 indicates the tensile strength and the breaking load at the blanking. The Breaking load is
proportional to the tensile strength. Figure 3 indicates the location of the punch and the sound pressure
when blanking SPC. t2 is the breaking time. The distance of p2-p3 is the break through distance.

Figure 1: Position of measurement devices

Figure 2: Breaking load in the Blanking

Figure 3: Punch trajectory and Sound pressure in the blanking speed V=15/mm/s and V=0.5mm/s

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Figure 4: Break through distance and Sound level

Figure 6: Sheared edge of SPC

Figure 5: Relation between breaking load
and Sound level

Figure 7: Relation between breaking load
and Break through distance

From Figure 3, we can confirm that the blanking noise was caused by break through. Figure 4 indicatesthe break through distance and the sound level. Both the break through distance and the sound level are
big at V=15mm/s. The break through distance is proportional to the material strength. And the soundlevel is proportional to the break through distance. The correlation coefficient of the break throughdistance and the sound level is 0.95.
4. Consideration
Figure 5 indicates a relation between the breaking load and the sound level in blanking. The soundlevel is proportional to the breaking load. But the sound level of V=15mm/s was bigger than V=0.5mm/s.
As shown in Figure 6, in spite of different blanking speed and different breaking load, the shearing
distance of SPC were almost same 0.78mm. It was the same with the other materials. It suggeststhat a factor besides the breaking load influences blanking noise. Figure 7 indicates a relation between
the blanking load and the break through distance. Even if the breaking load decreases, a line of V=15mm/s doesn't pass 0mm. The difference between the line of V=15 mm/s and the line of V=0.5mm/s ismore than 0.2 mm. In the break through distance, press slide is descended along planned trajectory. We
call it “planned descent”. Planned descent is the descent distance between t2-t3 in a planned trajectory
on figure 3. The broken line on figure 7 indicates the figure which deducted plan descent from the lineof V=15 mm/s. We considered that unknown factor hides in blanking vibration of V=15 mm/s. And thebreak through distance of V=15mm/s became big by this unknown factor. As shown in Figure 3, before
deviation of punch position is canceled, breaking of V=15 mm/s has occurred. Therefore we're inferringthat this unknown factor is related to control-system of screw drive servo press.
5. Conclusion
We confirmed that the breaking load is proportional to the tensile strength of materials, and the breakthrough distance is proportional to the breaking load. And we confirmed that the blanking noise wascaused by break through. But, in spite of the same breaking load, the sound level and the break throughdistance of V=15mm/s was bigger than V=0.5mm/s. It suggests that unknown factor hides in blanking
vibration. We're inferring that this unknown factor is related to control-system of servo press.
References[1] Chieko Yamagata, Masaaki Otsu, Kouzou Osakada, JSTM spring meet of 2002, (2002), pp.161-162.
[2] Tomohiro Murakami, Press Technology, vol41-4 (2003), pp. 28-33.
[3] Kazuki Aoi, Journal of JSTM, vol22-251 (1981), pp. 1188-1193.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

BEHAVIOR OF ASYMMETRIC DEFORMATION OF ROD IN LOCALLY-LATERALUPSETTING
Takashi Nomura1, Nguyen Sinh Con1, Kazuhito Asai1, Kazuhiko Kitamura1*
1

Nagoya Institute of technology*Corresponding. kitamura.kazuhiko@nitech.ac.jp

Keywords: Lateral upsetting, Material flow, Tooling, Asymmetry
1. Introduction
Remarkable progress in cold forging has recently brought manufacture of various products with anasymmetric and complicated shape. To achieve higher accuracy of these products efficiently,
appropriate preforms at a previous step are required in multiple-step cold forging. Ishikawa et al.1) has
proposed how to make an asymmetric product using synchronized a motion of servo-press with arobot handling a billet. Kuboki et al.2) has bended a band in plain by incremental stamping using aflexibly inclined and triangle punch. In both sophisticated methods, lower forming load using asimpler shaped tool extremely contributes to economic manufacturing. In this paper, a punch withappropriate nose shape and punch arrangement are investigated experimentally and numerically topromote asymmetric material flowing under lower forming load in locally-lateral upsetting of a bar.
2. Lateral upsetting outlines
Figure 1 shows the diagram of the locally-lateral upsetting. In this test, a punch compressed theintermediate position of a round billet. This billet is set to 0°, 15°, 30°, 45° of the angle between thebillet axis and the punch width direction.
Material of a punch and a die is die steel for cold forging (quenched and tempered, 60 HRC). This
punch has a cylindrical nose with the thicknesses of 6, 8, and 12 mm and the nose radii of 3, 4, and 6mm, respectively. The width of the punch is 50 mm. The punch is finished by lapping and the surface
roughness is 0.2 μm Rz. Compression ratios are 20%, 30%, 40%, and 50%. The stroke is 10 mm when
the compression ratio is 50%.
The punch pushes the billet at a speed of 1.0 mm/s using a hydronic press that is supported with four
rigid columns. Compression load measured load cell of a back punch.
The billet material is 99.7% aluminum as annealed. The billet has diameter of 20 mm and length of
40 mm. It has a yield point of 45 [MPa] and the yield stress can be approximated by Y = 115ε0.26 [MPa](0.1 ≦ ε ≦ 0.7). Lubrication used beef fat (75%) and graphite (25%). The coulomb’s frictionalcoefficient μ of 0.06 were measured by ring compression test under the lubricated condition.
Angle
θ = 0, 15, 30, 45°Compression direction
t

Punch

50 mm

Groove onlower dieLower die

y

40 mm

Dia. 20 mm

Billetzxx
Figure 1 The diagram of the locally-lateral upsetting.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

In this test, the billet end surfaces are shifted in the y-axis direction. This shift ratio Sy is calculated
y-distance yd is divided by 20 mm of initial diameter. The shift ratio Sy as meant asymmetry after
forging is estimated by measuring the shift distance of the billet after the test.
3. The punch-setting angle θ and billet on locally-lateral upsettingThis test was investigated the influence on the material behavior by angles. The billet axis was x-axis,
and vertex axis of the punch was y-axis. A punch used 8 mm of thickness and 4 mm of nose radius.
Figure 2(a) shows shift ratio by punch-setting angle θ and stroke. θ = 0° didn’t shift of any stroke. θ =15° and 30°, the shift ratio was increased gradually with increasing stroke. When stroke is 10 mm,
shift ratio is 4 and 10%, respectively. When θ is 45° and 4 mm stroke, shift ratio reached 5%.
Furthermore, when it pushed punch, rapidly shift ratio is increased. When the stroke is 10 mm, shift ratio
reached 29%. As the result, larger punch-setting angle θ is effective in shifting the both end of the billet.
The material can flow to the punch thickness direction when the punch-setting angle θ increases.
When punch-setting angle θ is 45°, it was investigated the influence on the material behavior to thedifferent radii of curvature. The punch thicknesses used 6, 8 and 12 mm and nose radii of 3, 4 and 6mm, respectively. Figure 2(b) shows the shift ratio by different punch thickness and stroke. The shift
ratio indicates until stroke is 5 mm. The shift ratio was 15% at the stroke 10 mm when the punch-nose
radius. The shift ratio is large without regard to stroke when the punch with nose radius of 6 mm is used.
Finally, the shift ratio reaches 36% at stroke 10 mm. As these results, the shift ratio Sy can increase
when the larger punch-nose radius and the larger punch-setting angle θ is used. Thus, the asymmetricpre-form can be obtained in the locally-lateral rod upsetting by these simple punches.
30%

40%

θ

S y = yd / d 030%

Shift ratio : Sy

Shift ratio : Sy

Shift : yd
20%

20%

s = 10 mm billetof θ = 45°

10%

10%

θ = 0°

0%
0

2

4

6
8
Stroke [mm]

(a)

10

0%
12

0

2

4
6
8
Stroke [mm]

10

12

(b)

Figure 2 Shift ratios after tests. (a) Different the punch-setting angle θ in the punch thickness of 8 mm. (b)Different radii of curvature in the punch-setting angle θ = 45°.

4. Conclusion
Influence of the punch-nose radius and the punch-setting angle on the behavior of asymmetricdeformation of rod was experimentally investigated in locally-lateral upsetting of the rod. The
following results were obtained:
1) The shift ratio Sy of the asymmetric pre-form increased with punch-setting angle increasing. The
maximum Sy reached 29% when the compression ratio was 50% using the cylindrical punch with thepunch-setting angle of 45°.
2) The shift ratio Sy of the asymmetric pre-form increased with the punch-nose radius increasing. The
maximum Sy reached 36% when the compression ratio was 50% using the cylindrical punch with thepunch-nose radius of 6 mm.
References[1] Takashi Ishikawa, Wang Xin-yun, Nobuki Yukawa, Yoshinori Yoshida, Takuya Sukeda, Incremental forming
using servo-press and robot, Proc. 5th JSTP-ISPF, (2009), pp. 41-46.
[2] Takashi Kuboki, Armad Azrie, Yingjun Jin, A new incremental in-plane bending of thin sheet metals for
micro machine components by using a tiltable punch, CIRP, Manufacturing Technology, 63(2014), pp.
249-252.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

DIRECT MEASUREMENT OF FRICTION FORCE DURING FORMING
Takehiko Makino1*, Yuya Yogo2, Toshinari Michimoto2, Yohei Takamori2
1

Department of Mechanical Engineering,
Nagoya Institute of Technology, Gokiso-cho, Showa-ku, Nagoya, 466-8555 Japan
2
Nagoya Institute of Technology, Gokiso-cho, Showa-ku, Nagoya, 466-8555 Japan* makino.takehiko@nitech.ac.jp

Keywords: Forming, Friction force, Direct measurement, Laser reflection
Abstract
Information about distribution and change in friction during forming is required to construct a
prediction model. The small deflection of a thin part fabricated at the other side of a flat tool surface is
detected using laser reflection to decompose the normal and horizontal (frictional) forces. The locus oflaser reflection on a detector showed both components. The values estimated from the locus werefairly close to the measured forces.
Introduction
The forming process, especially small-scale forging [1], is strongly affected by friction between thetool and material. Usually, the friction state is estimated by comparing the results of the load required
to perform the forming process and the load calculated by the finite element method with constantfriction coefficients. The information about distribution and change in friction during forming isrequired to construct a prediction model. Various methods for direct measurement of friction force
during forming have been reported. The pressure-pin method [2] detects the friction force using an
inclined pin placed in the cavity of the tool. In this method, the material tends to occupy the cavity,
and the friction force between the pin and tool cavity (commonly unknown) must be subtracted fromthe measured force. In the friction sensor method [3], the deflection of a thin part of the tool surface ismeasured by the strain gauges attached to the flat plate connected to the thin part. However, the thinpart of this method becomes relatively large in area because the deflection is still detected by the strain
gauges. The objective of the present study is to develop a method for direct friction force measurement
during small-scale forming using laser reflection.
Apparatus and method
In the proposed method, a small thin part (diameter: 2.0 mm, thickness: 0.42 mm) with a protrusion
was fabricated at the other side of a flat tool surface by electric discharge machining. The small
deflection of the thin part during forming is detected by two laser beams reflected by two mirrors
attached to the pin glued to the protrusion, as shown in Figure 1. The loci of the laser enlarged by the
optical lever were detected by two position sensitive detectors (PSDs, Hamamatsu Photonics
S5990-01) placed at both sides of the tool. The calibrations for each direction are performed using thetwo-axis press designed for this study (Figure 2). For calibration, the small pin (diameter: 2.0 mm,
height: 3.2 mm, material: SKH) is placed on the thin part and compressed and tensioned
simultaneously by the vertical and horizontal presses, respectively, as shown in Figure 3.

Figure 1 Principle of direct
friction measurement.

Figure 2 Two-axis press and two PSDs.

Figure 3 Schema of the calibration.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Results
Figure 4 shows the results of the calibration (1) with loading normal and horizontal forces applied
separately. In this case, (a) after the normal force reached the maximum value of 400 N, (b) the
horizontal force is applied while retaining the normal force. After the horizontal force reached thepredetermined value of 100 N, the horizontal force is released. The horizontal axis scale (time) for (a)and (b) is the same. After the horizontal force becomes zero, the normal force is released. Fig. 4 (c)shows that the locus on the left PSD moves in an upper-right direction from the origin (indicated by“O” in the figure) and then turns almost 90° to a lower-right direction at “P.” This movement
corresponds to the sequence of loading until the horizontal force reaches the predetermined value.
During unloading, the locus retraces the route conversely. Fig. 4 (d) indicates that the locus on theright PSD moves from the origin to the lower-left. However, the loci due to the normal and horizontal
forces overlap each other.
Figure 5 shows the results of the calibration (2) with loading normal and horizontal forces applied
simultaneously. In this case, as the normal force increases, the horizontal force is applied. In Fig. 5 (c),the turning appears at “P” when the horizontal force is applied. While the horizontal force is at aconstant value, the displacement in the locus due to the horizontal force continues.
(a)

(a)

(c)

(c)

P
O

(b)

O P

(b)

(d)

O

O

Figure 4 Results of the calibration (1). (a) Normal
force, (b) horizontal force, (c) locus on the left PSD,(d) locus on the right PSD.

(d)Figure 5 Results of the calibration (2). (a) Normal
force, (b) horizontal force, (c) locus on the left PSD,(d) locus on the right PSD.

Discussion
Logically speaking, the normal and horizontal forces need to be decomposed by the vertical loci of thelaser using two PSDs. However, the locus on the left PSD represents both components that can be
decomposed. Analysis of Fig. 5 (c) showed that the normal and horizontal forces estimated from thelocus (35 N and 311 N) were fairly close to the measured forces (36 N and 279 N, respectively).
Conclusion
The direct friction force measurement apparatus developed in this study decomposes the normal andfriction forces applied at a flat tool surface.
Acknowledgement
This work was supported by JSPS KAKENHI Grant Number 24560130.

References[1] Makino, T., Dohda, K., Ishitani, A., Zhang, H., “Anisotropy of plastic deformation in micro/meso-scale metal
forming -development of testing method”, Transactions of NAMRI/SME, 37(2009), pp. 330-340.
[2] Van Rooyen, G. T., Backofen, W. A., “Friction in cold rolling”, Journal of the Iron and Steel Institute, 186(1957), pp. 235-244.
[3] Yoneyama, T., Takahashi, M., “Development of a Friction Sensor and Measurement in Aluminum Hot
Extrusion”, Journal of Japan Society for Technology of Plasticity, 521(2004), pp. 82-470.

Page 40

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

MEASUREMENT OF TOOL TEMPERATURE
IN BURNISHING USING DIAMOND TIP
Masato Okada1*, Masayoshi Shinya2, Hidetake Tanaka3, Naoki Asakawa4, Masaaki Otsu1
1

University of Fukui, 2Graduate School of Natural Science and Technology, Kanazawa University
3
Sophia University, 4 Kanazawa University*Corresponding author: okada_m@u-fukui.ac.jp

Keywords: Surface finishing, Burnishing, Diamond tipped tool, Tool temperature, Two-color pyrometer
Introduction
Diamond burnishing is a micro-plastic forming method which uses a diamond-tipped tool to generatesmooth surfaces and work-hardened layers [1]. It offers a high level of processing efficiency relative togrinding and does not require dedicated equipment because it can be performed using the same machine
tools as those that are commonly used for cutting. In diamond burnishing, the tool temperature at theburnishing point is an important factor determining the degree of diamond tip wear and the target surface
quality. With the burnishing process, however, it is difficult to measure the tool temperature because theamount of heat and the size of the heated portion are very small. Furthermore, because the diamond tip
is an electrical insulator, a tool-work thermocouple configuration cannot be applied. To overcome this,a method of measuring the temperature of the diamond burnishing tool using a two-color pyrometerwith an optical fiber is proposed. In addition, the influence of the burnishing conditions on tool
temperature is investigated.
Experimental method and conditions
The setup for measuring the tool temperature is shown in
Fig. 1. The inner circumferential surface of the cylindrical workpiece, rotated by the main spindle of thelathe, was targeted. The tool temperature was measured using a two-color pyrometer with an opticalfiber, developed by the authors. This pyrometer offers a high response speed and a small measurement
area under non-contact conditions and eliminates the influence of the emissivity of the target by using
the ratio of the output voltages from two different infrared detectors [2]. The infrared rays emitted fromthe burnishing point of the diamond tip are accepted by the rotated probe, a small stainless tube housing
an optical fiber, which is fixed to the workpiece in the radial direction The collected rays are transmitted
to the other optical fiber, attached to the pyrometer, through a non-contact coupling. Thus, the infrared
rays are transmitted to the pyrometer at that instant when the burnishing point, the rotated probe, andthe optical fiber attached to the pyrometer are all aligned on the same axis. An accelerometer was
mounted on the shank of the burnishing tool to detect the micro vibrations caused by the diamond tip
passing the rotated probe. The experimental conditions are listed in Table 1.

Table 1 Experimental conditions
Tool
Workpiece

Burnishing speed
Thrust force
Feed rate
Lubrication

Diamond tipped tool
DT2D1, Sugino Machine Ltd.
Hardened stainless steel
JIS SUS420J2
53HRC
Ra = approx. 1.0 mm
v b= 100, 150, 200 m/min
F t = 90, 140, 180 N
f = 50 mm/rev
Dry

Fig. 1 Experimental setup

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Experimental results and discussions
In a case of the using a new tip, as shown in Fig. 2 (a), theoutput voltage from the pyrometer, for which the lower-limit measurable temperature is approximately
150°C, was not observed even when b = 200 m/min and Ft = 180 N. Subsequently, the output wasobserved in accordance with the increase in the burnishing length, and the output waveforms from thetwo-color detector, shown in Figs. 3 (a) and (b), were obtained when the tip wore as shown in Fig. 2 (b).
Thus, the tool was at a low temperature of 150°C or less in the primary stage, but the tool temperature
increased with the tool wear. Fig. 3 (c) also shows the output from the accelerometer. Given the output
from the accelerometer, it can be assumed that the burnishing point passes the rotated probe at
approximately  = 4.5 – 12 s. The temperature during this period is shown in Fig. 3 (d). The tool
temperature ranged from approximately 235°C to 305°C. It can be assumed that this distribution
occurred as a result of the slight change in the actual tool temperature during each revolution and thechange in the sliding condition as the diamond tip passes the stainless tube.
Figures 4 (a) and (b) show the influence of the thrust force and burnishing speed on the tool
temperature. The tool temperature was defined as the mean value of the five maximum temperatures
under each condition. These measured values were obtained when the tool wore as shown in Fig. 2 (b).
The tool temperature increased with an increase in the burnishing speed and thrust force, and the strong
influence of these burnishing conditions on the tool temperature was observed.
Conclusions
The infrared energy radiating from the contact area between the diamond tip andworkpiece can be obtained by fixing the optical fiber in the rotated probe. The tool temperature during
burnishing, which increases with the tool wear, can be measured by using the proposed method.
Moreover, it can be concluded that the burnishing speed and thrust force greatly influence the tool
temperature, and these results show that it is important to take the thermal effect on burnishing
characteristics into consideration.
Acknowledgements

The authors acknowledge the financial support for this study by the Machine Tool
Engineering Foundation and the Mazak Foundation. The authors also indebted to Sugino Machine Limited fortheir technical support.

References

Temperature °C[1] Mieczysƚaw Korzynski, Modeling and experimental validation of the force-surface roughness relation forsmoothing burnishing with a spherical tool, International Journal of Machine Tools and Manufacture, 47 (2007),
pp. 1956-1964.
[2] Masato Okada, Akira Hosokawa, Ryutaro Tanaka, Takashi Ueda, In-process monitoring of tool temperature inend milling by newly developed compact two-color pyrometer, Proceedings of the 8th International Conference
on Progress of Machining Technology, (2006), pp. 281-284.

0.5 mm(a) Virgin tip

250

vb = 200 m/min

200
150
50

100
Thrust force

150
Ft N

200

(a) Influence of Ft on temperature
Temperature °C

Wear region

0.5 mm(b)Worn tip
Fig. 2 Enlarged view of tip

300

300
250

Ft = 180 N

200
150

100
150
Circumferential speed

200
vb m/min(b) Influence of vb on temperature
Fig. 3 Output waveform
and temperature

Fig. 4 Influence of burnishingconditions on temperature

Page 42The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

HIGH DENSITY OXYGEN PLASMA ASHING
OF CVD-DIAMOND COATED END-MILLING TOOLS
E.E. Yunata1, K. Yamauchi1, T. Aizawa1*
1

Shibaura Institute of Technology*Corresponding. taizawa@sic.shibaura-it.ac.jp.

Keywords: Plasma technology, Oxygen plasmas, Ashing, Etching, CVD-diamond coated tools
INTRODUCTION
CVD-diamond coated end-milling tools have been utilized for dry machining of carbon fiber
reinforced plastic components and parts in the airplanes and automobiles [1]. Although they have highhardness even against the dry machining, they often suffer from chipping and wear in practice. Since
the tool substrate is made from WC (Co) and the designed tooth geometry is shaped into this tool, it
must be recycled and reused as a diamond-recoated tool by complete ashing the used diamond coatings
without significant loss of tool substrate materials [2-4]. In the present paper, high density oxygen
plasma ashing system is developed to make perfect ashing of CVD-diamond coatings within one hour.
The ashing rate of 10 to 15 m/h is much faster by 10 to 30 times than the conventional oxygen processes.
Optimum processing conditions are obtained to minimize the tool substrate damage and to reduce thetooth diameter shrinkage. The effect of plasma processing parameters on the ashing behavior is also
discussed by the on-line quantitative plasma diagnosis.
EXPERIMENTAL PROCEDURE
The present high density oxygen plasma ashing system consists of the vacuum chamber, the plasma
generator, the control unit, the carrier gas supply, and the plasma diagnosis instrument. The chamber isneutral in electricity; RF dipole electrodes and DC-bias work independently to generate RF and DC
plasmas, respectively. A standard experimental set up is summarized as follows. The base pressure isless than 5 x 10-3 Pa, and pure oxygen gas (purity; 99.99%) is only used as a carrier gas. RF voltage, DC
bias, and pressure are varied in a range from 100 V to 250 V, from -400 V to -600 V, and from 25 Pa to
100 Pa, respectively.
A hollow-cathode device is also employed in this system to control the external plasma condition toetching. As shown in Fig. 1 a), the generated RF-plasma is confined in this hollow so that higher ion
and electron densities are attained for ashing the DLC- and diamond films [5]. The diamond-coatedtool shown in Fig. 1 b) is placed in the center of hollow tube; the same DC-bias is applied to both thehollow and the tool. Under this set-up, the tool is rotated with the constant rotational speed of 6 to 10
rpm to make uniform ashing of diamonds on the whole tooth surfaces of tools. The above experimental
set-up of tools into the hollow has significant influence on the ashing behavior in practice.

a)b)Fig. 1: Experimental procedure for oxygen plasma ashing. a) Ashing system, and, b) Specimen.

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EXPERIMENTAL RESULTS and DISCUSSION
The RF-voltage has more importance on the oxygen plasma state, which is confined in the hollow
tube. In this experiment, this RF-voltage was reduced from 250 V to 100 V with keeping the DC biasby -500 V. As shown in Fig. 2 a), the whole tool surface just after ashing is covered by thin grey-coloredfilm, which is easy to be broken away by ultrasonic cleansing. Any oxide layers were never left on thepolished surface; as shown in Fig. 2 b), small tints were seen just after light cleansing. Figure 2 c) depicts
the microscopic image of top tooth edge. The original small holes were dotted on the tooth surface; they
were made by chemical treatment before diamond coating as a nucleation site of diamond films. The
initial WC (Co) tooth surfaces are recovered by the present plasma ashing condition.
The outer diameter (Dm) of top tooth after ashing was measured and compared to the calculated one(Dc) by subtracting the coating thickness from the outer diameter of coated top tooth. If Dm is muchless than Dc, the tooth edge and surface is significantly removed to lower the tool substrate life inrecycling. Hence, D (= Dc – Dm) plays an important parameter to evaluate the microscopic damage oftool tooth. Under this ashing conditions, D = 1 m. No significant defects and damages were seenin this condition.

a)

b)

c)

Fig. 2: Ashed specimen by the present method. a) Outlook of ashed CVD-diamond coated tool, b) A top
tooth of ashed tool after cleansing, and c) SEM image of tooth edge.

CONCLUSION
Perfect ashing of the used CVD-diamond coating without any damage and defects onto the WC (Co)substrates is completed by the present high density oxygen plasma ashing with use of the hollow cathodedevice. Direct control of the RF voltage as well as the DC-bias leads to optimization of ashing process
condition not to cause the over-oxidation of WC (Co) substrate but to make perfect removal of diamondcoatings. Short leading time by 3.6 ks or 1 h in the present ashing process is also attractive to toolingcompanies which often suffer from long ashing time over 20 hs.
References[1] R. Hasegawa, “Cutting tool for aircraft and applications,” Journal of Japan Society for Precision
Engineering.75 (8) (2009) pp. 953-957.
[2] E.E. Yunata, T. Aizawa, “High density oxygen plasma micro-texturing into CVD-diamond coated dies for
micro-embossing”, Proc. 10th 4M-ICOMM2015. (2015) pp. 460-465.
[3] E.E. Yunata, T. Aizawa, “Micro-texturing into DLC/diamond coated molds and dies via high density oxygen
plasma etching”, Manufacturing Review (2015) (in press).
[4] K. Yamauchi, E.E. Yunata, T. Aizawa, “High density oxygen plasma ashing of used CVD diamond coating forrecycling of WC (Co) tools,” Proc. IFMM-2015 (2015, May, Toyama) (in press).
[5] E.E. Yunata, T. Aizawa, D.J.H. Djoko Santjojo, “Characterization of hollow cathode plasma for etching andashing processes,” Proc. 8th SEATUC Conference (2014, UTM, Johor Balu, Malaysia) CD-ROM.

Page 44The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

FRONT AND REAR SIDES SIMULTANEOUS FORMING BY FRICTION STIR
INCREMENTAL FORMING
Masaaki Otsu1,＊, Seiya Ichikawa2, Takayuki Muranaka3, Hidenori Yoshimura4,
Ryo Matsumoto5, Masato Okada6
1

Professor, University of Fukui, 3-9-1 Bunkyo, Fukui, 910-8507, Japan
Graduate Student, University of Fukui, 3-9-1, Bunkyo, Fukui, 910-8507, Japan
3
Associate Professor, Fukui National College of Technology, Geshi, Sabae, 916-8507, Japan
4
Associate Professor, Kagawa University, Saiwaicho 1-1, Takamatsu, 760-8521, Japan
5
Associate Professor, Osaka University, 2-1 Yamadaoka, Suita, 565-0871, Japan
6
Lecturer, University of Fukui, 3-9-1 Bunkyo, Fukui, 910-8507, Japan*Corresponding Author: otsu@u-fukui.ac.jp
2

Keywords: Friction Stir Incremental Forming, Sheet Metal Forming, Aluminum Alloy Sheet, Front
and Rear Sides Simultaneous, Forming Accuracy
1. Introduction
The authors have developed friction stir incremental forming (FSIF) method for flexible forming
non-ferrous sheet metals without dies [1]. In conventional
incremental sheet metal forming, double side incremental
forming was developed to increase shape variety and improveforming accuracy [2]. In this study, double side friction stirincremental forming machine is developed. The formability,
formable working conditions, forming accuracy and sheet
thickness direction by single and double side friction stirincremental forming were investigated.
2. Experimental method
A forming equipment was developed originally. Appearanceof developed equipment is shown in Figure 1 and schematic
illustration is shown in Figure 2. Two sets of XYZ axes and aspindle are equipped. A schematic illustration of double side
friction stir incremental forming is drawn in Figure 3. A sheet
metal is clamped at the edge part by a blank holder, and formedby the two forming tools from upper and lower surface
simultaneously.
The diameter of the forming tools was 6 mm and the top
shape of the tool was hemispherical. AISI AA5052 aluminumalloy sheets were used for specimen. The size of the sheet was
100 mm x 100 mm and the thickness was 0.5 mm. The sheets
were formed into frustum of pyramidal shape with a height of 10mm. The tool rotation rate ω, feed rate v and wall angle θ werechanged and formable working conditions, thickness distribution
and forming accuracy were studied.
3. Experimental results
3.1 Formable working conditions
The tool feed rate was fixed to v = 1000 mm/min, and the tool
rotation rate ω and wall angle θ were changed to obtain theformable working conditions for both single and double side
friction stir incremental forming. The obtained formableworking conditions were plotted in Figure 4. Open circle marks
mean the forming was succeeded without fracture. Cross marks
indicate the sheet was fractured and forming was failed.

Figure 1 Appearance of forming
equipment

Figure 2 Schematic illustration of
double side friction stir incremental
forming equipment
Tool

Rotation direction

Sheet
Blank holder

Figure 3 Schematic illustration of
double side forming

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

50
45

o

Wall angle ,  [ ]

40
35
30
25
20
15
10
5

Tool feed rate v = 1000 mm/min

00 1 2 3 4 5 6 7 8 9 10 113Tool rotation rate ,  [x10 rpm](a) Single side
50
45

o

Wall angle ,  [ ]

40
35
30
25
20
15
10

Tool feed rate v = 1000 mm/min

50
0 1 2 3 4 5 6 7 8 3 9 10 11Tool rotation rate ,  [ x10 rpm](b) Double side
Figure 4 Formable working
conditions (v = 1000 mm/min)
5.5
5.0
4.5
Height [mm]The minimum wall angle by single side friction stirincremental forming was θ = 25 ° and that by double side friction
stir incremental forming was the same. This means theformability by single side forming and double side forming wasequal. However, the area of formable working conditions at low
tool rotation rate by double side forming was wider than that bysingle side forming. This reason is considerable that the friction
stirring phenomenon occurred easily due to clamping sheet by
two tool.
3.2 Cross-sectional shape of formed sheets
After forming with a height of 5 mm by both single and doubleside forming, the cross-sectional shape of formed sheets wasmeasured by a laser displacement sensor. Comparison among
ideal and formed by single and double tool is shown in Figure 5.
In a case of single side forming, actual wall angle was morethan 40 °, although the objective wall angle was θ = 40 °. This
reason is considerable that there are no counter supports, thesheet was elastically bent and the corner part between flange andconical part could not formed sharply.
In a case of double side forming, the corner part between
flange and conical part was formed more sharply than the single
side forming case and forming accuracy was improved
significantly.
3.3 Thickness distribution
Thickness distribution of formed sheets were measured. The
obtained thickness distributions were plotted in Figure 6. The
initial sheet thickness is 0.5 mm. Since the wall angle is θ = 40 °,the theoretical wall thickness at the conical part calculated bysine law is 0.321 mm.
In a case of single side forming, thickness changes graduallyand the minimum thickness was 0.403 mm. On the other hand, ina case of double side forming, the inclination of thickness change
is steeper than that of single side forming and the thickness is
closer to the ideal one. The minimum thickness was 0.215 mm
and this is smaller than the ideal one.

Single side
Double side
Ideal

4.0
3.5
3.0
2.5
2.0
1.5

References[1] Otsu, M., Ichikawa, T., Matsuda, M., Takashima, K., Friction Stir
Incremental Forming of Aluminum Alloy Sheets, Steel ResearchInternational, 81- 9(2010), pp. 942 - 945.
[2] Yongjun, W., Ying, H., Cao, J., Reddy, V.N., Experimental Study on
a New Method of Double Side Incremental Forming, Proceedings of
MSEC_ICM&P2008 (2008), 72279.

Thickness [mm]1.0
4. Conclusions
0.5
The original forming machine for double side friction stirincremental forming was developed. Aluminum alloy sheets were
0 2 4 6 8 10 12 14 16 18
Location [mm]formed by single and double side friction stir incremental Figure 5 Comparison
among formed
forming, and formable working conditions, formed shapes andand ideal cross-sectional shapesheet thickness distributions were compared. Formabilities ofsingle and double side forming were equal. However, in a case of
0.55
double side forming, the area of formable working conditions
0.50
was wider, the forming accuracy was higher and the sheet
0.45
thickness distribution is closer to the ideal one than those of
single side forming.
0.40
0.35
Ideal (Sine law)

0.30

Single side

0.25
0.20
0

Double side2

4

6 8 10 12 14 16 18
Location [mm]Figure 6 Distributions of formed
and ideal thickness

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

IN SITU OBSERVATIONS ON FRETTING WEAR BEHAVIOR OF
PLASMA-SPRAYED HYDROXYAPATITE COATING
Yuichi OTSUKA1*, Yukio MIYASHITA,2, Yoshiharu MUTOH3
1

Associate Professor, Department of System Safety Nagaoka University of Technology,1603-1 Kamitomioka Nagaoka , Niigata 940-2188 Japan
2
Associate Professor, Department of Mechanical engineering, Nagaoka University ofTechnology, 1603-1 Kamitomioka Nagaoka , Niigata 940-2188 Japan
3
Executive director / vice president, Nagaoka University of Technology, 1603-1
Kamitomioka Nagaoka , Niigata 940-2188 Japan*Corresponding Author:otsuka@vos.nagaokaut.ac.jp, Tel:+81-258-47-9575,
Fax+81-258-47-9573,

Keywords
Hydroxyapatite, plasma-spray, coating, interface mechanics, fretting fatigue, fretting wear
Introduction
This study aims at observing the effects of delaminations in Hydrocyapatite coating on wearing
behavior of artificial human bones in situ. Plasma-sprayed Hydroxyapatite (HAp) coating has been
widely used as bonding between human bone with artificial metallic implants in order to smoothlytransmit loads during gaits. However, the HAp coating is susceptible to be failed due to brittle fracture,
fatigue cracks, wears by fretting or corrosion etc. Wear particles made by fretting is also concerned
due to its possibility of activating inflammations at surrounding organs, which leads to loosening of
implants or subsequent failures of the implants. Delaminations in HAp coating interface can accelerate
the wearing behavior in fretting fatigue. Therefore, in-situ observation tests of fretting fatigue of HAp
coating were conducted.
Experimental Methods
Figure 1 shows the dimensions of testing specimen and contact pad. Testing specimen is made by
T-6Al-4V and contact pad is made by artificial cancellous bone (porous poly propylene 50 cpf, Avice,
Japan). Table 1 shows mechanical properties of the artificial cancellous bone. HAp coating isdeposited on the side surface of specimen by atmospheric plasma-spraying method. The side surfacesof HAp coating are subsequently polished by using emery papers from ＃120 to＃1200 in order toobserve the pass of delaminations with ease.
Fretting fatigue testing was conducted by using hydraulic fatigue testing machine(LMV3049, Saginomiya,
Japan）with confocal laser microscope (OLS1100, Olympus, Japan).Testing condition was;σmax = 150MPa，stress ratio R = 0.1, contact pressure Pc = 30MPa, respectively.
Table 1

Mechanical properties of artificial

cancellous bone(PP)Compressive
strength(MPa)
48

Young’smodulus(MPa)1469

Density(g/cc)
0,80

Fig. 1 Geometry of fretting testing specimen. (a)Specimen (b) contact pad

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Results and discussions.
Figure 2 shows the delamination pass after N= 2X 106 cycles. At initial stage vertical cracking at theedge of contact pads occurred due to stress concentrations by fretting contact. Delaminations initiatedfrom the root of vertical crack and then delamination grew to the center part of contact regions. Figure
3 shows the SEM picture of the edge of delaminations after testing. The wear debri were found at theresion adjacent to the edge of delaminations even in the cases of positive loading ratio. Figure 4 shows
delamination growth history and figure 5 shows delamination growth rate, respectively. Delamination
growth curve seems to be divided into two regions. In first region ( delamination length was less then
700 µm), growth rate decreases with increasing delamination length. The phenomenon can be
dominated by stress singularity at the edge of interface where intensity at the edge was higher than the
one at the subsequent interface cracks. In second region ( delamination length was more then
700 µm),the slope of growth rate was more decreased. Its mechanism is not clear but the wear particle
effects or changes in stress singularity are possible mechanism.
SummaryThis study aims at observing the effects of delaminations in Hydrocyapatite coating on wearing
behavior of artificial human bones in situ. HAp coating was delaminated and subsequently wear
particle was observed at the interface between contact pad with HAp coating. The changes in therigidity by delamination seems to promote wear behavior of HAp coating with contact pad. Fretting
fatigue behaviors in Simulated body fluid is going to be observed by the same system(references areomitted).
Ti-6Al-4V
Wear debri

Wear debri
Figure2

Observationsofdelamination

behavior of HAp coating during fretting fatigue

HAp

10μm

Figure 3 Wear debris at the tip of interface

Delamination length a, mdelamination10-3|σ|max=250MPa
f = 5Hz
R=0.1
Pc=30MPa

10-4
103

104
105
Numbber of cycles, N

Figure 4 Delamination growth plot in fretting
fatigue test. Dleamination initiated at N=8 X 103

Figure5

Delamination

growthrate

vsdelamination length. Delamination growth rate
decreases with increasing delamination length.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

WETTABILITY BETWEEN GRAPHITE AND HIGH- TEMPERATURE LEAD-FREE
SOLDER ALLOYGen Sasaki1*, Kunihiko Nakata1, Kenjiro Sugio1
1

Graduate School of Engineering, Hiroshima University*gen@hiroshima-u.ac.jp

Keywords: Composites, High-Temperature Lead-Free Solder, Wettability, Reactivity, Graphite
Although the development of lead-free solder metals at high-temperature over 473K is urgent study,
the alternative alloy used in conventional use is not yet developed. Sn, Bi, Zn base alloys are
considered as the candidate materials in previous studies. However, there alloys has many problems on
the ductility, the wettability, the long-term reliability (mechanical properties and thermal fatigue
properties) at the high-temperature, the reliability of the joining under the vibration environment andso on. On the other hand, some of the authors have developed the lead-free fuse [1, 2]. From the pointof view about easy control of the melting point,
electrical conductivity, the fusing, the corrosion
resistance and good joining ability, Zn-Sn alloy was selected as the base metal. In previous study [3],lead free Zn-Al-Sn alloy was developed using s-orbital energy level in order to satisfy both the tensile
strength of 200MPa and the elongation of 5%. One of suitable composition is Zn-10wt%Al-2wt%Sn(Zn-10Al-2Sn) alloy. But this planning value of strength and elongation for this alloy is still lower
than that of Pb base alloy. Especially, the low elongation and inferior properties on thermal fatigue
failure is fatal defects. Therefore, the addition of carbon short fiber to alloy is one of good idea for
improving these properties because of the effect of pull-out of fiber from matrix alloy and theimprovement of the thermal fatigue failure by suppressing the thermal expansion. Particularly, vapor
grown carbon fiber (VGCF) is promising materials for the dispersant in matrix alloy, which shows
good mechanical properties and high-cost performance. On the other hand, the chemical reaction
between alloy and carbon fiber during preparation of the composites may affect the microstructure andthe mechanical properties. Therefore, the wettability and the reactivity between the molten alloy andgraphite sheet with (002) plane in a surface, which is a same plane for surface of VGCF, isinvestigated in this study.
The Zn-10Al-2Sn alloy was obtained by mixing of the each pure element metals at 1023K, and thenpoured water cooled copper die. Graphite sheet was bonded to stainless steel plate with circle form
and was used for the substrate of dropping test. The condition of dropping test was at 1023K in Ar
atmosphere, and keeping time after dropping was 5, 10 and 30min. The contact angle was measured by
θ/2 method. Microstructure was observed by optical microscopy and EPMA.
Figure 1 show XRD analysis of graphite sheet used in this study. XRD analysis shows only (002) and(004) planes in graphite. It shows the surface of graphite consists from only (002) plane, which is very
similar with VGCF.
Figure 2 shows the droplet of molten Zn-10Al-2Sn alloy on graphite sheet at 1023K in argon
atmosphere as increasing contact time after dropping. Figure 3 shows the time dependence of contact
angle between Zn-10Al-2Sn alloy and graphite sheet at 1023K. The contact angle between graphite
sheet and Zn-10Al-2Sn alloy is 115-120 degree, which indicates this system has low wettability, whichshows the interface between the graphite and the alloy seems to have low interfacial strength. As
increasing time, the contact angle decreased for 1-4 degree, which means the chemical reaction
between alloy and graphite. Figure 3 shows the microstructure between alloy and graphite. Gap wasobserved at the interface between alloy and graphite. In equilibrium phase diagram between Zn-C,
Al-C and Sn-C, only Al reacts with C to form Al4C3. As Al4C3 is dissolved by water and humidity, it
seems the gap sign of the existence of Al4C3. On the other hand, the microstructure of alloy parts
shows two-phase, which is dendric Al base parts and Zn base parts. Sn was distributed in Zn base parts.
This structure is similar to monolithic Zn-10Al-2Sn alloy, and the effect of this reaction on properties
seems to be small.

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Figure 1 XRD analysis of surface of graphite sheet.

Figure 2 Configuration of droplet of molten alloy on graphite as a function of holding time after dropping.

Figure 3 Contact angle of droplet on graphite as afunction of holding time after dropping.

Figure 4 Microstructure around the interface between
droplet and graphite sheet by optical microstructure.
Holding time is 30 min.

References[1] K. Matsugi, Y. Iwashita, Y. B. Choi, G. Sasaki, K. Fujii, Long Time Stability of Pb-Free Sn-9Zn Elements
for AC-Low Voltage Fuse Performance, Materials Transactions 52,2, (2011), pp753-758[2] K. Matsugi, Y. Saki, Y. B. Choi, G. Sasaki, K. Suetsugu, K. Fujii, Control of Electrical and Thermal
Properties on Sn-50Zn Alloy by 8vol% Al2O3 Addition for Pb-Free AC Low Voltage Fuse Elements, Materials
Transactions 54, 2, (2013), pp. 231-237[3] K. Matsugi, K. Terada, S. Furukawa, Y.B. Choi, G. Sasaki, K. Suetsugu, Design of Zn-Al-Sn Alloys as
Lead-Free Solders for High Temperature Application and Their Characteristics, J. JFS, 86, 3, (2014), pp216-222[4] G. Sasaki, M.H, Niu, K. Sugio, Y.B. Choi, K. Matsugi, Development and Mechanical Properties of High
Temperature Lead Free Solder Composites, Abs of The 11th China-Japan Joint Conf. on Composite Materials(CJJCC-11) (2014)

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

THE ROLE OF HYDROGEN IN HIGH RATE PLASMA NITRIDING OF
MARTENSITE STAINLESS STEEL
D.J. Santjojo1*, T.Aizawa2, Istiroyah1
1

CRC-ASMAT Brawijaya University, 2Shibaura Institute of Technology*dsantjojo@ub.ac.id.

Keywords: High rate plasma nitriding, surface reaction, hydrogen plasma, metastable nitrogen
INTRODUCTION
High concentration of nitrogen with nearly constant profile toward the nitriding front end cannot be
explained with conventional nitriding mechanisms. Complex reactions on the surface of the substrateduring the nitriding process produce large amount of certain nitrogen species which can be diffuse easily
into the substrate. Analysis of optical emission spectroscopy data combined with Langmuir probe
measurement data showed that inclusion of hydrogen in the plasma leads to high density ion production.
The high density ion contains of metastable state nitrogen and NH species reacting to produce reactive
nitrogen ion and atoms. The hydrogen atoms played important role in the complex reactions with themetastable nitrogen and the NH.

EXPERIMENTAL PROCEDURE
Specimens for this study were produced by means of a RF-DC Plasma nitriding system(Fig.1a). The martensitic steel used was AISI 420 which was cut and shiny polished into 5 mm thicksilinders with a diameter of 10 mm. Investigation on the role of hydrogen in the high rate plasma
nitriding was carried out by a combination of optical emission spectroscopy and Langmuir probe
measurement (Fig.1b). Composition of elements of the nitride subsurface layer of the specimens was
obtained from auger electron spectroscopy (AES) measurements. The Composition of elements to thedepth was acquired by serial measurements of subsequent surface etching.

Fig. 1: a) Plasma nitriding system; b) A combination of OES and Langmuir probe measurement.

EXPERIMENTAL RESULTS and DISCUSSIONHigh rate plasma nitriding was estimated from depth profile auger electron spectroscopy. A highnitrogen concentration (40%) profile was found almost constant up to 3 µm. This was produced bymeans of a mixture of hydrogen and nitrogen plasma nitriding. The role of hydrogen was investigatedby comparing the result with the one produced using nitrogen plasma only (Fig.2) The active species
for steel nitriding are the metastable state nitrogen such as N2 (A3 ∑+u , ϑ) or vibrationally excited states
1+∑such asN2 (X g , ϑ). The active species cannot be observed directly from the OES spectra. However,Page 51

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015indirectobservationis possible by investigating transitions of the second positiveat 337, 358 and 370 nm which is related to the N 2* and the first negative2 ++N2+ (B 2 ∑+u → X ∑g ) at 391 and 427 nm which related to the N 2 . Other important state of the plasma
in the nitriding mechanism is the NH species which is found at 336 nm (Nagamatsu,2013).
45

N Concentration (%)

40

0%H

35

17%H

30

a)

25
20
15
10

c)

5
0
0

1

2

3

4

5

6

Depth (µm)
7dV/dI (k)

6

b)

5
4
3
2
1
0
0%

20%

40%

60%

80%

100%

Hydrogen Inclusion

Fig. 2: a) AES depth profile data; b) Plasma electrical resistance as a function of hydrogen inclusion; c)
OES spectra of nitrogen plasma and nitrogen-hydrogen plasma.

CONCLUSION
Our study concludes that the high rate nitriding is governed and controled by dynamic of nitrogen andhydrogen species in the plasma. The dynamic of the mixed gas plasma depends on the inclusion ofhydrogen in the nitrogen plasma. OES and Langmuir probe measurement suggest that low concentrationof hydrogen plasma (17%) contains a large density of positively charge species which are N2*, N2+
and NH complexes. Intermediate species exist in the plasma reaction and emit observable radiation. The
species react on the surface producing high concentration of high energy atomic nitrogen which diffuse
into the metal uniformly.
References[1] Alphonsa, A. Chaini, P.M Roale, B. Ganguli, P.I John, A study of martensitic stainless steel AISI 420 modified
using plasma nitriding, Surf. Coat. Technol., 150, 2002.
[2] Pinedo C. E., Waldemar A. Monteiro, On the kinetics of plasma nitriding a martensitic stainless steel type AISI
420, Surf. Coat. Technol, 179, 2004.
[3] Yun-tao Xi, Liu Dao-xin, han Dong, Improvement of errosion-corrosion resistance of AISI 420 stainles steel
by low temperature plasma nitriding, Appl. Surf. Sci., 258, 2008.
[4] Santjojo D.J., et.al, Formation of Expanded Martensite In Plasma Nitrided AISI420 Stainless Steel, Proceeding
SEATUC 8th Symposium, Malaysia, 2013.
[5] Nagamatsu H., et.al, Steel nitriding by atmospheric pressure plasma jet using N2/H2 mixture gas, Surface &Coatings Technology, 225,2013.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

DEVELOPMENT OF TITANIUM FIBER/ALUMINUM MULTIFUNCTIONALCOMPOSITES
1

1
2, S. Yamaguchi 2, T. Mochizuki 2, T. Yanaseko 2 and Suprianto3

Professor, Chiba University, 1-33 Yayoicho, Inage-ku, Chiba-shi, Chiba, 263-8522 Japan

Graduate student, Chiba University, 1-33 Yayoicho, Inage-ku,Chiba-shi, Chiba, 263-8522 Japan
3

Lecturer, University of Sumatera Utara, Medan 20155 Indonesia*Corresponding Author: asanuma@faculty.chiba-u.jp, Tel:+81-43-290-3201

Keywords: Multifunctional, Metal Matrix Composites, Innovative Processing, Interphase, Aluminum
Introduction

Asanuma et al. have developed a simple temperature and strain sensor by embedding surface
surface-oxidized nickel (NiO/Ni) fiber in an aluminum matrix in the previous work [1]. In thisstudy, two more useful functions, that is, heating and actuation were tried to be added in addition
to temperature sensing and strain sensing by using surface-oxidized titanium fiber by making
use of its higher electrical resistance and lower thermal expansion coefficient compared withthose of nickel fiber.
Experimental
Pure aluminum plates of 0.2 mm thickness, 8 mm width and 30 mm length were used as matrix. A 0.01
mm thick copper foil was used as insert. Titanium fiber of 0.15 mm diameter and 99.5% purity was used
as functional fiber. It was oxidized at 1073 K for 1.8 ks in the air in an electric furnace to obtain thesurface-oxidized (TiO2/Ti) fiber. These materials were consolidated by the IF/B method [1], which canbe briefly explained as following. The TiO2/Ti fiber was arranged in the U-grooves made on thealuminum plate laminated with the copper insert, and they were covered with another matrix aluminum
plate without U-grooves to make unsymmetrical configuration for actuation. These piled-up materials
were hot-pressed under the condition of 853 K, 1.36 MPa and 0.3 ks in a low vacuum.
Results and discussion
A SEM image of a cross section of the composite embedded with the TiO2/Ti fiber is given in Figure 1.
It shows that the TiO2/Ti fiber can be embedded in aluminum matrix without fracture of the surface
oxide layer when it is hot-pressed under the condition of 823 K, 1.36 MPa and 0.3 ks in a low vacuum.
The relation between the temperature obtained by the embedded TiO2/Ti fiber and the surface
temperature of the active composite measured with an external thermocouple is given in Figure 2, whichshows good coincidence of the temperatures even after deformation of the active composite up to higher
than 450K. According to the result, it is clear that the curvature of the active composite can be monitoredby the temperature measurement with the embedded fiber. The relation between the electrical resistanceof the embedded TiO2/Ti fiber and the surface temperature of the active composite is given in Figure 3.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

This shows an almost linear relationship between them, which means that the curvature of the active
composite can be monitored also by the resistance measurement though it includes the both effects ofthe temperature increase as well as the strain reduction. Figure 4 shows the curvature change of theactive composite during the fifth heating and cooling cycle between room temperature and 813K.
According to th result, the embedded fiber is working as a heater for actuation and its monitoring. In
addition, the shape of the composite at room temperature is reproducible even after the thermal cycle up
to 813 K. Figure 5 shows the curvature change of the active composite during the thermal cycle up to493 K. It shows that the curvature of the active composite during the thermal cycle up to about 490 K
can change without hysteresis.
Conclusions
The results obtained in this study are as follows:(1) The oxide layer formed on the titanium fiber at 1073 K for 1.8 ks in the air remains unbroken during
the composite fabrication by the IF/B method.
(2) The embedded titanium fiber works as a heater for actuation and a sensor for temperature anddeformation monitoring.
(3) The composite shape at room temperature is reproducible even after thermal cycle up to 813 K.
(4) The curvature of the composite during the thermal cycle up to 493 K changes without hysteresis.
References[1] H. Asanuma, The Development of Metal-Based Smart Composites, JOM, 52-10 (2000), pp. 21-25.

Fig. 1 SEM image ofa cross section of thecomposite.

Fig. 2 Relation between thetemperature measured by the embedded
fiber and that of specimen surface
measured with an external thermocouple.

Fig. 4 Curvature change of the active
composite during the fifth thermal cycle
between room temperature and 813K.

Fig. 3 Relation between theelectrical resistance change rateof the embedded fiber and thetemperature of specimen surface.

Fig. 5 Curvature change of the active composite
during the thermal cycle up to 493 K.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

PHASE-FILED SIMULATION OF
ELEMENTARY PROCESS IN POWDER BED FUSION
Kazunari Shinagawa
Faculty of Engineering, Kagawa University
shina@eng.kagawa-u.ac.jp

Keywords: powder, simulation, phase-field method, discrete-element method, selective beam melting

Introduction
Powder bed fusion, a 3D printing technique, is expected to be a new way of manufacturing by its
ability to produce high value-added products or supply wide variety of parts in small quantities.
Considering the benefits, optimizing the manufacturing conditions may become more important withincreasing the practical applications.

Manufacturing conditions, determined by empirical laws, havebeen mostly obtained through trial and error.
materials and machines will be desired.

A general law which can be applied to different

This study aims to establish a method for analyzing thefusion process of metal powder particles to optimize the conditions of the selective beam melting
process in a rational manner.

The author has developed a combined phase-field/discrete-element

method (PFM/DEM) to simulate both shrinkage and grain growth behavior in powder compacts
during sintering [1].

In this paper, PFM/DEM is applied to melted particles, and an elemental fusion

process is simulated as the first step of the study.
Numerical method
In PFM, the migration of interface between particles is computed with minimizing the interfacial
energy.

Combining with DEM, the force acting between particles (which stems from the surface

tension) is evaluated in the phase field by taking into consideration the surface curvature.
the force consists of two terms.

Note that

The first term is a compressive force directly acting on the contact

plane between particles, while the second one is an imaginary compressive force converted from theoutward force at the neck tip.

As for the simulation of solid state sintering, the position of the necktip can be easily detected in the phase field because the neck between particles is the grain boundary as
well, where the phase-field variable changes sharply.

For liquid phase, melted particles do not havethe grain boundary, but low grain boundary energy with different phase-filed variables is given toadjoining particles, for finding the positon of the neck by searching the change in phase-field variable.
The resultant forces are transferred into DEM to compute the motion of particles.
Calculation condition and result
As in the previous study [1], a two-dimensional model was used in the simulation.

The grid

points of 200x200 with an interval of x = y = 2 were generated for the finite difference method to
conduct the phase-field simulation.

A model of powder having size distribution with the mass

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

median diameter of 40 and the geometrical standard deviation of about 1.4 was prepared in the grid byusing random number generation.

In the central circular region with a diameter of 240 in the grid,

the particles were assumed as liquid phase, and a large value of interface mobility as well as a smallvalue of the grain boundary energy was set.

On the assumption that the outer region is affected byheat, the solid state sintering was simulated for the surrounding particles.
Figure 1 shows an example of the calculated results.

The particles in the central region fused

together with time t, while the sintering process with grain growth was observed in the other region.
In Fig. 1(c), rounded pores can be seen in a glob of liquid phase.

The liquid phase seemed to shrink alittle apart from the outer region.

(a)

t=0

(b) t = 200
Figure 1 Calculated result(c) t = 2000

Summary
By using a PFM/DEM combined method, the formation of a glob of liquid phase was calculated
in a two-dimensional simplified simulation, as the first step of the numerical study on powder bed
fusion.

Other factors of selective beam melting, including scanning operation, will be introduced

into the present method, and the whole process of powder bed fusion will be simulated in the next
study.
Acknowledgements
This work was supported by Grants-in-Aid for Scientific Research (B) (No. 26289273) from Japan
Society for the Promotion of Science.
References[1] K. Shinagawa: Simulation of grain growth and sintering process by combined phase-field/discrete-element
method, Acta Materialia, 66(2014), pp. 360-369.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

ESTIMATION OF FRICTIONAL PERFORMANCE OF LUBRICANTS FOR HOTFORGING OF STEEL BY CARIBRATION CURVES CONSIDERING HEATTRANSFER IN RING COMPRRESSION TEST
Kazuhito Asai1, Kazuhiko Kitamura1*
1

Dept. of Engineering physics, Electronics and Mechanics, Nagoya Institute of Technology,
Gokiso-cho, Showa-ku, Nagoya 466-8555, Japan,* Corresponding Author: kitamura.kazuhiko@nitech.ac.jp, Tel:+81-52-735-5351, Fax:+81-52-735-5442

Keywords: (Tribology, friction, hot forging, testing method, lubricant)
1. Introduction
Higher performance lubricants are required to improve die life in hot forging of steel. The frictional
performance of these lubricants has been often estimated by ring compression test. A low-speed hotring compression test1) was proposed to be helpful in classifying the tribological performances of
lubricants in hot forging of steel. By this test, high performances lubricants could be selected, whichshowed low-frictional coefficient and high heat-resistance performance. To use more precise values ofthese frictional coefficients, calibration curves considering heat transfer between a billet and dies are
determined by FEM analysis. In this report, frictional coefficients under poor-lubricated condition are
estimated by the low- and high-speed hot-ring compression test.
2. FEM analysis conditions of ring compression test
The conventional cold-ring compression test is widely used and roughly analyzed under anisothermal condition as shown in Table 1 (a). A ring billet is made of 0.45% carbon steel, and acommercial data base of this steel is used in FEM analysis. The temperature of this billet and dies is
20ºC. The billet has dimensions of 10 mm in height, 15 mm in internal diameter, and 30 mm in outer
diameter. A coulomb's frictional coefficient µ between dies and the billet is calculated when reductionin height Re is 50% and compression speed is 2.0 mm/s.
Table 1 (b) shows an analysis condition of hot-ring compression test considering heat transfer.
Billet temperature is 1100ºC before compression. Mechanical press takes 0.1 s and hydraulic press
takes 8.0 s for this test with 50% reduction in height. These testing times are input in FEM analysis. A
heat transfer coefficient2) α between the billet and the open air is 0.1 kW·m-2 ·K-1, and α between thedies and the air is also 0.1 kW·m-2 ·K-1. By considering the oxide thickness of the billet after heating, αbetween the dies and the billet is 12.5 kW·m-2 ·K-1. The emissivity2) of the billet or the die is 0.7.
An adiabatic condition of hot-ring compression test is also referred as shown in Table 1 (c), namely
the heat transfer coefficient between the billet and the dies is neglected under this condition.

Table 1Analysis conditions for hot ring compression test of steel

Upper die
BilletLower die(a) Cold-ring compression test (Isothermal condition)
Billet0.45% carbon steel, R.T
Original height: internal diameter: out side dia.
= 10 mm : 15 mm : 30 mm
Dies(Upper/Lower) Non-consider heat transfer, R.T
Frictional coefficient
Coulomb's friction µ
Reduction in height Re 50%
Compression speed
2.0 mm/s(b) Considering heat transfer for hot-ring compression test
Billet0.45% carbon steel, 1100ºC
Dies(Upper/Lower)
Consider heat transfer, R.T.
Compression time
8.0 s (Hydraulic press)(Re= 50%)
0.1 s (Mechanical press)
Heat transfer coefficienta) Billet - Air, Dies - Ai r 0.1 kW/m2Kb) Dies - Billet
12.5 kW/m2K
Emissivity (Dies and Billet) 0.7(c) Adiabatic condition for hot-ring compression test
Heat transfer coefficientb) Dies - Billet
0 kW/m2K

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Coulomb's frictional coefficient µ

3. Estimation of frictional coefficient in hot-ring compression test
Figure 1 shows a relationship between a decrease in an internal diameter and a coulomb’s frictionalcoefficient µ by hot-ring compression test. Under a poor-lubricated condition µ indicates 0.3 or more
for low-speed hot-ring compression tests when the hydraulic press is used. Calibration curves for
hot-ring compression the test leads the different µ from µ determined with the calibration curves forcold-ring compression test. It is necessary to calculate more appropriate calibration curves for hot-ring
compression test. The calculation considers temperature, compression speed, and heat transfer.
Re= 20%

Re= 40% Re= 50%

0.4
Cold-ring compression test1b) Considering heat transfer for hot-ring compression test(µ ≥ 0.3) Hot-ring compression test by mechanical press(µ < 0.3) Hot-ring compression test by mechanical press(µ ≥ 0.3) Hot-ring compression test by hydraulic press(µ < 0.3) Hot-ring compression test by hydraulic press

0.3
0.2
0.1
0
-10

Figure 1

0 10 20 30 40 50 60
Decrease in internal diameter /%Calibration curves of ring compression test with considering heat transfer

Coulomb's frictional coefficient µ

Figure 2 shows the relationship between a decrease in an internal diameter and a coulomb’s frictionalcoefficient µ without considering heat transfer in hot-ring compression test. These frictionalcoefficients µ are close to µ that is estimated with considering heat transfer because the test takes for
short time using mechanical press. The frictional coefficients calculated with considering heat transfer are
different from µ that is calculated without considering heat transfer because this test takes for long timeusing hydraulic press. The billet temperature can decrease when the test takes for long time inlow-speed hot-ring compression test. Particularly, the temperature of the billet near the dies can fall
down to be low. Material flow is locally difficult because the material is strengthened by cooling. It isalso important to consider the material strength in the analysis of hot-ring compression test.
Re= 20%

Re= 40% Re= 50%

0.4
0.3
0.2
0.1
0
-10

Figure 2

Cold-ring compression test1b) Considering heat transfer for hot-ring compression test(µ ≥ 0.3) Hot-ring compression test by mechanical press(µ < 0.3) Hot-ring compression test by mechanical press(µ ≥ 0.3) Hot-ring compression test by hydraulic press(µ < 0.3) Hot-ring compression test by hydraulic press1c) Adiabatic condition for hot-ring compression test
Hot-ring compression test by mechanical press
Hot-ring compression test by hydraulic press

0 10 20 30 40 50 60
Decrease in internal diameter /%Calibration curves with and without considering heat transfer

4. Conclusion and outlook
Calibration carves for hot-ring compression test were calculated with considering heat transfer undera poor-lubricated condition. The billet temperature, heat transfer coefficient, compression speed
directly influenced on increasing in material strength by cooling of the ring billet. To obtain the more
appropriate calibration curves, it is necessary to consider influence of lubricant on heat transfer undera well-lubricated condition in hot forging.
References[1] Kazuhito Asai, Kazuhiko Kitamura, Estimation of frictional property of lubricants for hot forging of steel
using low-speed ring compression test, 11th ICTP2014, Procedia Engineering 81 (2014), pp. 1970-1975.
[2] Nobuki Yukawa, Yoshihiro Nakashima, Takahiro Ishiguro, Eiji Abe, Modeling of heat transfer coefficient ofoxide scale in hot forging, 11th ICTP2014, Procedia Engineering 81 (2014), pp.492-497.

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DEVELOPMENT OF HIGH SPEED TORSIONAL FATIGUE TESTING MACHINEFOR MULTIPLE SPECIMENS
Katsunori TAKEZAWA1, Noriyasu OGUMA2,*, Mai FUKAMI3 and Masanori KONTANI3
1

Graduate Student, Graduate School of Science and Engineering, University of Toyama, Gofuku 3190,
Toyama-shi, Toyama 930-8555 Japan
2
Professor, Faculty of Engineering, University of Toyama, Gofuku 3190, Toyama-shi, Toyama 930-8555 Japan
3
Engineer, Tohshin Seiki Co., Ltd., Ha 18 Terai, Nomi-shi, Ishikawa 923-1121 Japanoguma@eng.u-toyama.ac.jp, Tel&Fax:+81-76-445-6776

Keywords: Torsional fatigue, Displacement controlled, Rotating-reciprocating-swinging conversion,
Sliding friction, Multiple specimens
Introduction
Torque transmission shafts, rolling bearings, coil springs and so on are important elements todynamic mechanical device. It is known that fatigue destruction of these mechanical elements occurs
mainly under shear stress. Therefore it is important to clarify fatigue strength characteristics of thematerial used for these mechanical elements under shear stress from design aspect. However, most of
tests to grasp fatigue characteristics have been performed under normal stress in rotating bending tests
or axial loading fatigue tests. This is because the torsion testing machine which is low-priced and canbe efficiently tested is not widely popularized. In these circumstances, we aimed for the development of
the compact torsion fatigue testing machine in which four specimens can be tested at high speed at thesame time to efficiently accumulate torsional fatigue data on materials for machine structures. And thevalidity of the torsion fatigue testing machine developed was estimated by the inspection of the torsional
angle of specimens and the fractographic observation of fracture surfaces of the failed specimens in
torsional fatigue testing.
Driving mechanism of torsional testing machine
The developed testing machine uses an electric motor as a drive source, and has a rotatingreciprocating conversion and a reciprocating-swinging conversion mechanism. A schematic diagram of
the former mechanism is shown in Fig. 1. Its mechanism consists of an eccentric shaft, a connecting rod
and a piston. Figure 2 is a schematic diagram of the latter one. Its mechanism consists of a piston, link
pins and four wheels. Specimens are fixed on the center of the wheels and twisted by swing of the wheels.
To set up several loading levels for the specimens, each wheel has plural pin holes where the distances
are different from center. However, since connecting the two wheels on a single link pin, the stress levelcan be tested at one time is two conditions. Needle roller bearings are press-fitted in the pin holes for
reduction of gaps and friction. Contact portions of the component parts with relative movement of thetesting machine are rolling contact and so as to suppress the heat generation due to friction.

Figure 1 Rotating-reciprocating conversion
mechanism

Figure 2 Reciprocating-swinging conversion
mechanism

Function confirmation of torsional testing machineFirst, the relationship between the operating frequency and the temperature
rotating-reciprocating conversion mechanism was examined to evaluate fever due to friction.
Temperature measured points were around the bearing, the bushing and housing. Temperature
measurement was carried out a step-up idling method using a thermocouple. As a result of the

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0.035

0.14
0.12

Torsion angle, radSwinging angle of wheel, radmeasurement, the temperature at any points became gradually higher with the increase of the operating
frequency, and the temperature in the vicinity of the bearing was the highest temperature. The
temperature was stabilized at 304 K with operating frequency 50 Hz. Thus we concluded that this testingmachine can operate by a frequency range to 50 Hz.
Next, the swinging angles of the wheels were confirmed to evaluate the test accuracy. The swinging
angle was calculated by measuring the displacement of the reflection plate which was attached to thewheel shaft by a laser displacement meter. Figure 3 shows the swinging angles of the wheels obtainedin each hole position by inserting the link pin and the design value. In the figure, two kinds of plots inthe same hole numbers are meant to be respective data of the two wheels. The measurement results
obtained at all of the hole positions fell within less than 5 % of errors from design values. Therefore we
judged that the intended dynamic functions of this testing machine could be reproduced.
Then, the torsion angle of specimen was confirmed to evaluate the test accuracy. The torsional angleof specimen was also measured using the laser displacement meter. Measurement points were threepositions of near the chuck portions and the center of the specimen. Figure 4 shows the measured
torsional angle of the specimen. Also the designed torsion angle is indicated by a solid line. The
measurement values fell within less than 1 % of error from design values.
Based upon the foregoing, we succeeded in development of the torsional fatigue testing machine
which can conduct fatigue tests for multiple specimens at high speed.

0.10
0.08
0.06

measurement value(a)measurement value(b)design value

0.04
0.02

1

2

3

4

5

6

7

8

9

measurement value

0.030

ー design value

0.025
0.020
0.015
0.010
0.005

10

Hole number

Figure 3 Swinging angle of wheels obtained ineach hole position and design value

0

5

10

15

20

25

Distance from chuck end, mm

Figure 4 Measured value and design value of
torsion angle of specimen

Torsional Fatigue Test
Torsion fatigue tests of JIS-S45C material were performed by means of the developed testingmachine in the room atmosphere without any control of temperature and humidity. Diameter of thechuck portion is 10 mm, and the parallel portion for dumbbell-shaped specimen is 3 mm. It is provided
with two fillets in order to reduce the stress concentration. The stress concentration factor of thisspecimen in torsion is given as 1.049 from FEM analysis. Example of photo image of a crack generated
in the specimen at stress level of τa = 705 MPa, N = 3.0×104 cycles is indicated in Fig. 5. It was
confirmed that fatigue crack was generated in the axial direction and perpendicular thereto acircumferential direction of the specimen due to shear stress. Exact fatigue life is currently underconsideration on the basis of the changes in the torsional angle data, because the rate of the changes of
torsional angle is found to vary depending on the relationship between the measurement position of
torsional angle and the crack generating position.

Figure 5 Example of photo image of a crack
generated in the specimen

Conclusions
Torsional fatigue testing machine in which four specimens can be tested at high speed at the sametime was developed. This testing machine can be operated up to a load frequency of 50 Hz. Result of
verification of the torsional angle of specimen, the accuracy was less than 1 % of the design value. In
addition, as results of the torsional fatigue tests, fatigue cracks were generated in the axial andcircumferential direction, it was confirmed that the shear stress was dominated fatigue fracture.

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GUIDED WAVE PROPAGATION IN WIRE CABLE UNDER TENSILE STRESS
Takuma MATSUO1* and Shunsuke OZONO1
1

School of Science and Technology, Meiji University*matsuo@meiji.ac.jp

Keywords: Wire cable, Guided wave, Nondestructive testing, Wavelet transform, Group velocity

Maximum amplitude, mV

Wire cables are widely used for mechanical and structural components. To evaluate the damage
and degradation of a wire cable, a nondestructive inspection technique utilizing guided waves isdesirable. However, calculating the propagation mode in wire cables is difficult because they consistof twisted threads. In our previous study, we examined the guided-wave propagation in a wire cablewith no tensile force [1]. We found that the group velocity and amplitude dispersion were not similar
to those of a rod. Moreover, the guided waves were mainly propagated in threads where the input
signal was excited. In the present study, we experimentally evaluated the propagation of guided waves
under tensile stress.
The propagation behavior of guided
waves under tensile stress was first evaluatedusing the experimental setup shown in Fig. 1.
Guided waves were produced and detectedby a pair of PZT transducers (PAC, R50).
The transducers were installed on thewire-cable surface, separated by 600 mm.
Step signals of 300 V was applied to one of
the transducers, and a guided wave was
excited. The guided wave was detected by theother transducer and fed to the oscilloscope
via a pre-amp and 100-kHz high-pass filter.
The wire cablea fiber core woundwith seven strands. The core and strands werecomposed of twined steel wires. The tested
Fig. 1 Experimental setup for detecting guided wave
wire cable was made of SUS304 and had an
propagated on wire cable with tensile stress.
outer diameter of 3.0 mm [2]. Under these
conditions, the amplitude of the signal had
dispersion because the contact force of the5threads in the wire cable was not constant.
On the other hand, the detected waveforms
were constant in the wire cable with tensile
4
force. According to the wavelet contour mapsof the detected waves, the propagation modeof the guided waves was similar to the3longitudinal and flexural modes of a rod.
These modes were denoted as the L-likemode and F-like mode. Figure 2 shows the2change in the maximum amplitude of the
L-like mode packet. These data were detected
0
100
200
using the positions of the transducers for
Tensile force, N
which the maximum amplitude signal was
detected with no tensile stress. The maximum
Fig. 2 Change in peak-to-peak amplitude of
amplitude decreased as the tensile force
longitudinal-like mode packet as a function ofincreased. This trend was caused by thecontraction of the cross section, which
tensile load.
caused the contact among the neighboring

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Rod

0

Rod

1.0

315

45

174 N

0

270

225

90

45
0.5

0

270

135

17.4 N

1.0

315

174 N

0.5

0

225

90

135

17.4 N 180

180

Fig. 3 Normalized maximum-amplitude profile of L-like (left)
wire and F-like (right) modes as a function of
Lmode normalize
Lmode normalize

cable(17.4N)Fmodeangle of input transducer.
wire

cable(174.0N)FmoderodFmode

Amplitude loss, %unit strands in the wire to be more greatly influenced by the increasing tensile-stress intensity.
The out-of-plane displacement of each mode was examined. Figure 3 shows normalized
maximum-amplitude profiles for the L-like (left) and F-like (right) modes. The out-of-plane
displacement for the L-like mode was almost constant, whereas the displacement for the F-like mode
had angular dependence and changed as a function of the axial distance from the excitation point. Thischange was according to the thread on which excitation transducer was mounted. These characteristics
were similar to those of an elastic wave in a rod [3].
Next, the relationship between the defect area and the amplitude of the guided wave was
evaluated. Three types of 10-mm wire cablewith different defect areas were compared,
100
and the amplitude of the transmitted wave wasmeasured. The experimental setup was thesame as that shown in Fig. 1; however, adefect was added in the center of the wire
cable. Figure 4 shows the relationship50
between the cross-sectional loss andamplitude loss for the L-like mode. Here, theblack circles indicate the average results andsolid bar indicated the dispersion of the data.
The amplitude loss of the L-like mode and0
defect size exhibited a relationship; however,
0
10
20
30
dispersion of the amplitude occurred. It was
Cross sectional loss, %determined that the contact between thesensor and thread was not constant.
Fig.4 Relationship between cross-sectional loss andReferences

amplitude loss for L-like mode.

[1] Matsuo, T. and Seshime, A., Elastic Wave
Propagation in Wire Cable for Corrosion
Monitoringand
Evaluation,
Corrosion
Engineering, 63-11(2014), pp.456-461.
[2] Ono, S., Kougaku-no-Suugaku (Mathematics for Engineers), AGNE Gijutsu Center, 2005, p.5 (in Japanese).
[3] Rose, J., Ultrasonic Waves in Solid Media, Cambridge University Press, 1999.

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BEHAVIOR ON GENERATION AND PROPAGATION OF FATIGUECRACK FOR WC-Co CEMENTED CARBIDES
Hiroko MIKADO1,2, Shingo KAWAMURA2 and Noriyasu OGUMA3,*
1

Graduate Student, Graduate School of Science and Engineering, University of Toyama, Gofuku 3190,
Toyama-shi, Toyama 930-8555 Japan
2
Engineer, Machinery and Engineering Group, YKK Corp., Yoshida 200, Kurobe-shi, Toyama 938-8601 Japan
3
Professor, Faculty of Engineering, University of Toyama, Gofuku 3190, Toyama-shi, Toyama 930-8555 Japan* oguma@eng.u-toyama.ac.jp, Tel&Fax:+81-76-445-6776

Keywords: WC-Co cemented carbides, Fatigue crack generation, Fatigue crack propergation, Grainsize, Crack growth pass
Introduction
It is well known that WC-Co cemented carbides have excellent wear resistance. Taking advantage
of its characteristics, the material is used for cutting tools and dies. However, the tools or dies in somecases leading to fracture failure in cracking or chipping during use. Therefore it is important to grasp
the fatigue crack growth resistance because of the brittle material. In this study, in order to examine theinfluence of WC grain size exerting the fatigue crack generation and its propagation, rotating bending
fatigue tests and short crack growth tests were performed using WC-Co cemented carbides specimens
with three grades WC grain-sizes. The plane specimens and the artificially defect specimens were used
for the tests respectively.
Specimen and Experimental Procedure
The materials used in this study were WC-13%Co cemented carbides with different WC grain sizes.
They are labeled as fine-grained WC (average grain size of 0.45µm), medium-grained WC (0.76µm)
and coarse-grained WC (1.08µm), respectively. Figure 1 shows the shape and dimensions of specimens.
Diameter of the critical portion for the hourglass type specimen is 3 mm, and the round notch radius is
10 mm. The stress concentration factor of this specimen in bending is given as 1.04. Figure 1(a) indicatesthe plane specimen, and (b) indicates the micro-defect specimen. For the latter, the artificial defect witha length of 80 µm, a width of 20 m and a depth of 40 µm was introduced into the center of the specimen
by laser beam machining.

(in mm)(a) Plane specimen(b) Micro-defect specimen
Figure 1 Shape and dimensions of specimens(in mm)Rotating bending fatigue tests were performed by means of a cantilever-type rotating bending fatigue
testing machine in the room atmosphere without any control of temperature and humidity. Frequency of
cyclic load applied 5-15 Hz. To evaluate the rate of fatigue crack growth of a micro surface crack, thefatigue tests were interrupted at constant intervals to obtain replicas of the specimen surface. The crack
lengths recorded on the replicas were measured using an optical microscope.
Experimental results
Figure 2 shows the experimental S-N characteristics for three grades of WC grain specimens. From
this figure, the fatigue life for the specimens of finer WC grain-size was longer. An example of SEM

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Figure 2 S-N characteristics for WC-Co
cemented carbides

Figure 3 SEM image of the fracture surface forcoarse-grained WC specimen atσa = 1100MPa

Figure 4 Crack growth behavior tested at σa =600MPaFigure 5 Crack propagation path appeared on
the surface for coarse-grained WC specimen

observation of the fracture surface for coarse-grained WC specimen is shown in Fig. 3. It is found thatthe crack initiation sites were relatively large WC grain in the surface. The crack origin of all the failed
specimens were such a large WC grain. Therefore it can be considered that the fatigue life depends onan existence probability of large WC grain.
Figure 4 indicates the increase of crack lengths of the specimen surfaces shown in Fig. 1(b) as afunction of number of cycles from the results of the micro surface crack growth tests. As can be seenfrom Fig. 4, the crack growth rate for the specimen of rougher WC grain-size was lower. The observation
photograph of crack propagation path of the surface for coarse-grained WC specimen is shown in Fig.
5. The fatigue crack grows mainly along the WC/WC boundary or WC/Co boundary with avoiding theWC grains.
Considering by comparing the results of the above two types of experiment, the crack propagation
life of fine-grained WC alloy is short, and its crack generation life is much longer. On the contrary, thecrack propagation life of coarse-grained WC alloy is long, and its crack generation life is not so long.
Thus the finer WC grain size to be suitable to improve the fatigue strength of WC-Co cemented carbides.
Conclusions(1) The results of rotating bending fatigue tests, the fatigue life of the finer grained WC-Co cemented
carbides was improved.
(2) According to the results of crack propagation tests, the crack propagation life of the coarse grained
WC-Co alloy was longer. Thus the fatigue life of WC-Co cemented carbides is estimated to be
governed to the crack generation life.
(3) From the two types of experiment, it is concluded that the finer WC grain size to be suitable toimprove the fatigue strength of WC-Co cemented carbides.

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A DEVICE OF GAS TRANSMISSION RATE FOR THIN FILMS
IN A HYDROGEN STORAGE SYSTEM
Hiroki Mano1,*, Mitsuo Notomi2
1

Graduate Student, Department of Mechanical Engineering, Meiji University, 1-1-1 Higashi-mita, Tama-ku,
Kawasaki, Kanagawa, 214-8571 JAPAN
2
Professor, ditto*ce42076@meiji.ac.jp

Keywords: Gas transmission rate, Hydrogen storage amount, Measuring device, Thin film, Carbon
based material
Introduction
Hydrogen storage materials, which are represented by hydrogen storage alloys and/or carbon
based materials, have attracted a great deal of attention. Although hydrogen storage alloys havehigh hydrogen storage density per unit volume, their hydrogen storage density per unit mass is low
due to their weight, and they have inferior cycling characteristics. Carbon based materials possess lightweight, high specific surface area and nanopores which are valuable in hydrogen adsorption and thematerials have excellent cycling characteristics due to the physical adsorption that depends on
temperature and/or pressure change monotonically. In addition the materials show high adsorption
performance at low temperature under high pressure and they alone can’t really adsorb at room underatmospheric one.
One of the solutions for increasing the adsorbing amount is that the materials are wrapped inthin films of which hydrogen permeation depends on temperature and/or pressure change
non-monotonically. If hydrogen permeation of the films has changed dramatically at a point of
temperature and/or pressure, the materials with the thin films could control the amount for thetemperature and/or the pressure. Hirotaki et al.[1] prepared the samples of Mg-Formate powder sealedwith the laminated film, originated Teflon tape with Silicon adhesive and sputtered Pt as a catalyst,
and they have reported that sealing with it improves hydrogen absorption amount of the Mg-Formate;while that of the commercial its powder is 0.297 wt.% at 298 K at 2.0 MPa, that of the sample is 0.424
wt.% at the same conditions. Since the Mg-Formate also adsorbs hydrogen with the physicaladsorption as with the carbon based materials, it can be expected that the hydrogen storage systemwith the materials shows higher hydrogen storage property through the improvement of the thin films.
We have considered using the laminated films which are composed of polymer and metal for the thinfilms, and we need to measure correctly hydrogen transmission rate and storage amount of the thinfilms alone at the same time. The purpose of this study was development of such a measuring device
for inventing the hydrogen storage system.
Design of the gas transmission rate measuring device
Gas transmission rates of films are determined according to JIS K 7126-1 (Plastics - Filmand sheeting - Determination of gas transmission rate - Part 1: Differential - pressure method)[2].
Figure 1 shows system diagram of the gas transmission rate measuring device which we designed on
the basis of the JIS without a valve, a tank and a pressure sensor written in red lines. Measurement
section is divided into high pressure side (HS) and low pressure side (LS) by a specimen, thelaminated film. We introduce test gas into HS and calculate the gas transmission rate with pressurechange in LS by the following equation.
Vc
GTR 
 dp
(1)
dt
R T  A P
u

GTR : Gas transmission rate[mol/(m2･s･Pa)], Vc : Volume of LS[l], T : Test temperature[K], Pu :Pressure of HS[Pa], A : Area of transmission[m2], dp/dt : Pressure change of LS per unit time[Pa/s],
R : 8.31×103[L・Pa/(K・mol)]Page 65

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Furthermore,
wedividedthemeasurement section in two parts, measurement
system (MS) and sample system (SS), by a valve
written in red line, and we added a pressure
sensor to MS for measuring the hydrogenstorage amount. It is based on JIS H 7201[3].
We can calculate the theoretical value ofpressure in MS plus SS using ideal gas law withvolume and former pressure in MS and volume
in SS. Hydrogen storage amount is obtained bycomparing pressure in MS plus SS when it isequilibrium state after proceeding transmission
and the theoretical value.
Figure 1 System diagram of the gas transmission
In addition, we added a tank to MS torate measuring device
avoid pressure decrease of HS due to gas
transmission from HS to LS. Entire assembly
drawing of the gas transmission rate measuring device based on the above is shown in Figure 2; left isa front view and right is a side view. In order to inhibit tubes deformation, each valve is fixed on apanel, the lower flange is fixed on a table and a flexible tube is used for a part of connecting pump.
Conclusion
We designed the gas transmission rate measuring device which can also measure hydrogenstorage amount for thin films, and this device was designed to combine the principle of JIS K 7126-1
and H 7201. Next, we will evaluate some polymer thin films with this device and eventually develop
the hydrogen storage system with the films and carbon based materials.
References[1] Shota Hirotaki, Mitsuo Notomi, Hydrogen absorption property of Magnesium Formate by spillover at theambient temperature, TMS2015, March 18/2015[2] JIS K 7126-1, Plastics-Film and sheeting-Determination of gas-transmission rate-Part 1: Differential-pressure
method, (2006), pp. 1-6[3] JIS H 7201, Method for measurement of pressure-composition-temperature (PCT) relations of hydrogen
absorbing alloys, (2007), pp,1-7

Figure 2

Entire assembly drawing of the gas transmission rate measuring device

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THE STUDY OF VARIANT SELECTION FOR SHAPE MEMORY EFFECTON FE-MN-SI BASED ALLOY WITH EBSD
Naohiro OKUBO1,*, Hiromu ISHII1, Mitsuo NOTOMI2
1 Graduate Student, Department of Mechanical Engineering, Meiji University
2 Professor, Department of Mechanical Engineering, Meiji University*

ce42017@meiji.ac.jp

Keywords: Fe-Mn-Si based alloy, shape memory effect, stress induced transformation, variant
selection, single variant
1. Introduction
Shape memory alloys can recover the original shape by heating and/or unloading after deformed.
Ni-Ti alloy has used practically for many applications due to the fine recovery effect. Contrast to poor
workability and expensive cost of Ni-Ti alloy, ferrous shape memory alloys have good workability andreasonable cost. Fe-Mn-Si based alloy has been utilized in pipe joints, although the amount of thisshape recovery is less than that of Ni-Ti alloy.
The shape memory effect (SME) in Fe-Mn-Si based alloy is associated with γ⇄ε martensitic
transformation which is caused by motion of the Shockley partial dislocation and martensite phase cansometimes made to form variants by stress induced transformation in several conditions. The fourorientations of the variants in Fe-Mn-Si based alloy could be formed according to Shouji-Nishiyama(S-N) orientation relationship between the austenite phase (A-phase) and the martensite phase(M-phase). If one direction of the variants only was formed in the transformation, what is called as the
single variant formation, this alloy would have a good SME since the back way would have been
limited in inverse transformation. Some authors [1][2] reported the single variant formation understress induced transformation and the effect of variant selection in SME.
We studied the crystal orientation of M-phase in Fe-Mn-Si based alloy with electron backscatter
diffraction (EBSD). The amount of shape recovery was quantified by means of bending test. The best
shape recovery ratio was about 86% for the specimen annealed 873[K]. From these results, the effectof variant selection to SME was discussed and also we studied the relationship between the variant
selection and the direction of the applied stress on the basis of crystallography.
2. Experimental methods
Fe-28Mn-6Si-5Cr alloy (mass%) was used as a wire, its diameter is φ1.2 [mm] and the wire wascut three pieces as the length about 5[mm]. The chemical composition is showed in Table 1. One ofthem is called As-received specimen. The others were annealed with a holding jig to memory theoriginal shape. They were heated at 873[K] and 1273[K] for one hour and then cooled in furnace andthese specimens are called each Ann.873 specimen and Ann.1273 specimen respectively.
To obtain the shape recovery, each specimen was bended at a constant radius of curvature using
the material testing machine with strain controlled and heated them to 923[K] for 15 minutes in the
purpose of shape recovery. The best shape recovery ratio was about 86% for Ann. 873 specimen as
shown in Table 2. The EBSD analysis was carried out to only Ann.873 specimen and to tensile area
after bending Ann.873 specimen. We cut them 5[mm] length and pile up specimens and carbon tape on
aluminum die, finally embedded together in cup of φ8[mm]. We polished the fine surface as the lessroughness to measure with EBSD.
Table 1

Chemical composition (mass%)

CSi
0.047 5.69
CrNi
5.02 0.033

Mn
PS
27.2 <0.003 <0.009
NO
0.012 0.002

Table 2

Shape recovery ratio in each specimen

Specimen nameShape recovery ratio
As-received specimen 11% (10[deg.]/90[deg.])
Ann.873 specimen
86% (62[deg.]/73[deg.])
Ann.1273 specimen 28% (28[deg.]/100[deg.])

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3. Result and discussion
Table 3 shows Inverse Pole Figure (IPF) maps and Pole Figures (PF) in Ann.873 specimen. In IPF
maps, each color level represents to the crystal orientation at the point for the normal direction to thepolished surface according to the stereo triangle indices shown in the top right of each figures. The PF
indicates the direction of crystal orientation at each position analyzed in IPF maps to a circle projected
sphere by means of stereographic projection and density of crystal orientation represents the color
contrast. In Table 3, the amount of M-phases (e) to (g) increases after bending instead of decreasing
the amount of A-phases (a) to (c). M-phases might be formed from A-phases due to stress-induced
transformation caused by bending deformation. In Table 3, tensile texture before bending, (b) isexisted to the longitudinal direction due to wire drawing and variants are formed to prior orientation
after bending, (h). Variants might be formed to prior orientation for variant selection changed bytensile texture and annealing (f) to (h). It found that formation of variants to prior orientation is related
to variant selection and that affects the improvement of SME.
We studied particularly the relationship between prior orientation of variants and tensile stress by
deformation. Fig. 1 shows PFs of each single grain enclosed color line in Table 3 (c). In Fig. 1, fourcolor circles of each single grain correspond to the crystal orientation of A-phase. Variants were only
formed to one direction of red circle after deformation and that means a single variant is formed bystress-induced transformation. Fig. 2 shows the angles of prior orientation of variants to tensile stress.
The red and blue arrows in Fig. 2 represent each direction of tensile stress and prior orientation. Highdegree of Schmid’s factor, 0.47 was calculated at the angle between red and blue arrows and shows
more activation of dislocation movement. If dislocation movement is more activated, variants might
be formed to prior orientation.
4. Conclusion
Variant selection might correspond to the activation of dislocation movement as one of the factor. The
grain of high Schmid’s factor produces concentrating a prior orientation in four directions and thenimproves SME.
References[1] O. Matsumura ,S. Furusako ,T. Furukawa and H. Otsuka, Formation of Surface Texture and Anisotropy of
Shape Memory Effect in an Fe-Mn-Si Alloy, ISIJ Int. 36, (1996), pp1103[2] B.H. Jiang , X. Qi , Y.L Ren and C.Q. Wang, Effect of Annealing Texture on Shape Memory Effect in an
Fe-Mn-Si-Cr Alloy Wire, Mater. Trans. JIM 41, (2000), pp1703
Table 3 Change of crystal orientation in Ann.873 specimen

Before bending
IPF map
Austenite (fcc)

PF

{111}
{111}

001

Martensite (hcp)

Tensile area after bending
IPF mapPF001

101 111

{111}
{111}

101 111

,

25μm

(a)

0001 10-10

1
2
3
4
5
6

1

(b)

2

{0001}{0001}(c)

25[μm]
000110-10

Y1

11-20

(d)
{0001}{0001}

Y1

11-20
X1

X1

25μm

(e)

3
6
9
12

(f)

(g)

25[μm]

(h)

1
2
3
4
5
6
7

After deformation
{111}
{111}

001
101

{111}
{111}
111

001

Tensile stress
101

111

<111>

54.74[deg.]

Sh =cos54.74cos(90-54.74)=0.47(a) Grain enclosed yellow line (b) Grain enclosed white line
Fig. 1 Crystal orientation of a single variant
Fig. 2

Variant selection and Schmid’s factor

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

CHARACTERISTICS OF METAL MATRIX COMPOSITES REINFORCED WITH
PALM OIL FLY ASH FABRICATED BY USING STIR CASTING AND
CENTRIFUGAL CASTING METHODSSuprianto1*, Tugiman1,Chandra Andrika2, Nursuci Ade2, Hiroshi Asanuma3
1

Lecturer, Mechanical Engineering Department, Faculty of Engineering, University of Sumatera Utara
Jalan Almamater Kampus USU, Medan 20155 INDONESIA
2
Student, Mechanical Engineering Department, Faculty of Engineering, University of Sumatera Utara
Jalan Almamater Kampus USU, Medan 20155 INDONESIA
3
Professor, Graduate School of Engineering, Chiba University
1-33, Yayoicho, Inage-ku, Chiba-shi, Chiba, 263-8522, Japan*Corresponding author:suprianto.t@gmail.com,Tel:(061)8211236, Fax: (061)8213250

Keywords: Metal Matrix Composite, Casting, Aluminum
1. Introduction
Metal Matrix Composites (MMCs) are type of materials that have been developed in recent years
to be used widely. The uses of MMCs have different goals such asreduction of mass, wear resistance,and so on.The properties of MMCsare determined by the nature of the basic constituent components[1]. Fly ash can improve the properties of MMCs such as hardness, compression strength and impact
strength due to an increase in mass%fly ash together with a decrease in the density [2,3]. MMCs havebeen fabricated by using casting methods, where ceramic particles mixed with molten metals as amatrix [4]. Stir casting is one of the methods which have succeeded in making MMCs [5] inconjunctionwith centrifugal casting to obtain higher hardness of castings [6]. Most of the manufacturesof MMCs for strengthening purposes use fly ash from coal,and rarely use that from palm oil ascomposite reinforcement. The objective of this research is to examine the opportunity of using palm
oil fly ash in the manufacture of MMCs using stir and centrifugal casting methods.
2. Experimental
Variation of composition was conducted from 2.5 to 12.5mass%palm oil fly ash, put into molten
aluminum in crucible at735oC pouring temperature and 450oC mold one. Locally produced fly ash
having the typical composition showed in Figure 1(b)was used in this research. Samples obtained fromthis processes were tested to obtain mechanical and physical characteristics.Figure 1(a) shows aschematic of centrifugal casting that was used to fabricate them.
3. Results and discussion
The densities of the MMCs fabricated by stir casting as a function of composition are shown in
Figure 2(a).Because of low density of the palm oil fly ash, density of the MMCs tends to decrease withincreasing its mass%. Figure 2(b) shows hardness of the MMCs, where the hardness increases withincreasing the percentage of fly ash. The results of impact testing are shown in Figure 2(c) andmicrostructures are shown in Figure 3(a), (b) and (c). Decrease in impact strength was found, which isthought to be caused by poor wettability and inhomogeneous distribution of fly ash. Microstructure
Ladel
o

735 C

CaO

MgO 0.44%

1.12%

14.35%Al melt

10.55%

Al2O3

450 C

Gas

3.18%

FeO

366 rpm

Permanentmold

Houshingobalance70.36%
SiO2

(b)
(a)
Figure 1(a) Schematic of centrifugal casting and (b) palm oil fly ash composition.

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consist of a matrix of aluminum MMCs with palm oil fly ash and Si distributed among the grainboundaries. Figure 4 shows the impact energy influenced by grain size and elongation decreased than
raw material, but stable as the increase in mass% fly ash.

2.60

BHN

Density, ρ/g・cm-3

90

(a)

2.65

2.55

50

2.50

30

30

(b)

70

2.45

(c)

Joule

2.70

25
20

----- Stir Casting-----Centrifugal Casting----- Stir Casting-----Centrifugal Casting

15

10

2.40

Raw
2.5

5

7.5

10

2.5

5

7.5

10

12.5

10

12.5

Mass%
0

2.5

5

7.5

10

12.5

Figure 2 Effect of mass% fly ash on (a) density, (b) hardness and (c) impact energy of MMCs.
(b)(a)
Palm oil flyash(c)More silicon
contents

Palm oilfly ash

Si

Si

Al

Figure 3 Microstructures of 5mass% palm oil fly ash MMC (a) optical micrograph (HF etched), (b) SEM
photograph, and (c) result of EDS analysis.
(c)

(a)

%Elongation

Joule

25
20
15

----- Stir Casting10
0.053 0.074 0.105 0.149

0.21

μm

6
5
4
3
2
1
0

(c)

(b)

Al
fly ash----- Stir Casting
Si

0

2.5

5

7.5

200µ

10 12.5 mass%

Figure 4 (a) Effect of particle size on impact energy, (b) effect of mass%fly ash on %elongation and (c)microstructure of 10 mass% fly ash.

4. Conclusions
MMCs reinforced with palm oil fly ash were tried to be fabricated in this study and the following
conclusions were obtained:
1) Stir casting and centrifugal casting have successfully produced MMCs reinforced with palm
oil fly ash, of which hardness increase with increasing mass% of palm oil fly ash in theexperimented range.
2) The distribution of the fly ash particles tends to localize along the grain boundaries and not
homogeneous.
References
[1]
[2]
[3]
[4]

[5]
[6]

K.U.Kainer, Metal Matrix Composites. Custom-made Materials for Automotive and Aerospace
Engineering, WILEY-VCH Verlag GmbH & Co. KGaA, Weinheim, ISBN: 3-527-31360-5 (2006).
K.V.Mahendradan and K. Radhakrisna, Fabrication of Al–4.5% Cu alloy with fly ash metal matrix
composites and its characterization, Materials Science-Poland, Vol. 25, No. 1 (2007), 57-68.
J. Babu Rao, D.Venkata Raodan and N.R.M.R.Bhargava, Development of light weight ALFA
composites, International Journal of Engineering, Science and Technology, Vol.2, No. 11 (2010), 50-59.
S.Sarangidan and D. Kumar, Fabrication and Characterization of Aluminium-Fly Ash Composite Using
Stir Casting Method, A Thesis, Department of Metallurgy & Materials Engineering, National Instituteof Technology,Rourkel, (2009), 16-17.
J. Hashim, The Production Of Cast Metal Matrix Composite By A Modified Stir Casting Method,
JurnalTeknologi, 35 (A), (2001), 9-20.
A.Jayakumardan and M.Rangaraj, Property Analysis of Aluminium (LM-25) Metal Matrix Composite,
International Journal of Emerging Technology and Advanced Engineering,Vol.4,Issue 2(2014),495-501.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

MECHANICAL PROPERTIES OF AMORPHOUS CARBON FILMSDEPOSITED FROM CARBON ISOTOPE MATERIALS
Yutaro Suzuki1, Junko Hieda1, Naoto Ohtake1, Hiroki Akasaka*1
1Tokyo Institute of Technology, 2-12-1 Ookayama, Meguro-ku Tokyo, 152-8552 Japan*akasaka@mech.titech.ac.jp

Keywords: Amorphous carbon film, Stable isotope, Hardness, True density, Isotopic effect
1. Introduction
An amorphous carbon (a-C) film is mainly consisted of sp2 and sp3-bonded carbons and showsinteresting mechanical properties owing to mixed structures of sp2 and sp3-bonded carbons. At isotopesof carbon atom, 13C is stable and usable nonradioactive isotope material. 13C has different characteristics
from 12C; thus, the incorporation of 13C into a-C films may lead to more advanced functions1). On studiesof electrical properties on the diamond, some studies reported about 13C isotopic effects2,3). On the other
hand, the deposition of amorphous isotopic carbon films are few. In this study, amorphous carbon films
which contained isotopes of carbon were deposited, and their structure and properties were evaluated toinvestigate isotopic effects.
2. Experimental
a-C and hydrogenated a-C(a-C:H) films which contained isotope of carbon were deposited on a ptype Si(100) substrate using a pulsed magnetron sputtering or a plasma chemical vapor deposition. On
a-C deposition, Carbon (99.98%) and 13C (99%) targets were used as source materials. On a-C:H
deposition, methane(CH4, 99.999%), deuterated methane(CD4, 99%) or 13C methane(13CH4, 99 %) gases
were used as source materials. These gases were introduced into the vacuum chamber and a 13.56 MHz
radio frequency power of 20 W was applied to a substrate at 20 Pa.
Film density was evaluated by X-ray reflectivity (XRR). The XRR pattern was obtained using a CuKα
X-ray diffractometer (X'Pert PRO MRD Phillips). The structure of the films was analyzed by Raman
spectroscopy (NRS-1000 JASCO) using a Nd:YVO4 laser (532 nm) with an aperture size of= 200 μm.
The hardness of the film was estimated by an indentation test using a Vickers diamond tip. To obtain thehardness of the film of nm-order thickness, a high-resolution indenter (HM500 Fischer Instruments)
with a pico-meter depth resolution was used. Indentation reached maximum loads of 100 N at 10 s. At
maximum loads, the load was kept 5 s to obtain the creep rate.
3. Results and discussion
In this abstract, results about a-C:H films were described. Although the deposition conditions of thesefilms were the same except for the raw material, film structure and properties indicated difference.
Figure 1 shows XRR patterns of the films. By simulation fitting, the XRR pattern densities of a-12C:H(1.70 g/cm3), a-12C:D (1.54 g/cm3) and a-13C:H (1.75 g/cm3) films were obtained. The density of the a12
C:D film was lower than that of the other films. One of the possible reasons for this difference is thatthe plasma condition at the deposition had large difference because their dissociation and ionizationenergies were affected by the isotopic effect.
Figure 2 shows the Raman spectra of the a-C:H films. In all the Raman spectra, two broad bands appear,
namely, the G (graphitic)-band at approximately 1500 cm-1 and the D (disorder)-band at approximately
1300 cm-1. G-band peaks originate from the E2g resonance oscillation in graphite structures and are dueto the residual sp2 bonds. Hence, the G-band represents the order of a graphite structure. On the other
hand, Tuinstra and Koenig assigned the D-band to first-order scattering from a zone boundary phonon
activated by disorder on the basis of the finite crystallite size4). Briefly, the D-band represents theboundary of the graphite structure. To analyze data on the order of graphite cluster size, the integrated
peak ratio, ID/IG, was calculated from the Raman spectra. The change in ID/IG suggests a change in thecontinuous order of graphite cluster size. These Raman spectra were fitted by Gaussian line shapes using
curve-fitting software. Values of ID/IG were 0.57, 0.70 and 0.37 for a-12C:H, a-13C:H and a-12C:D films,
respectively. The order of the graphite cluster size of a-13C:H is smaller than that of a-12C:H. The Gpeak position of the a-13C:H film indicated a clearly lower wave number than that of the a-12C:H and a12
C:D films. This shift was reported as an isotopic effect5). From these values of ID/IG, the graphite cluster

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Fig. 1. XRR patterns of the a-C:H films

Fig. 2. Raman spectra of the a-C:H films

size of a-12C:D is larger than that of other films. Hence, the reason of low density compared with other2 films is not graphite layer. This reason may be caused by containing of the polymer-like components
at a-C:D film. Raman spectra for a-C:D indicated the gradient base line which mean containing of theuniaxially oriented conjugating polymer.
Table Results of nano-indentation tests.
a-12C:H a-13C:H a-12C:D
Martens hardness [GPa]
4.2
4.1
3.0
Young's modulus [GPa]81
81
56
Creep rate [nm/s]
0.10
0.07
0.11

Figure 3 shows indentation depth-load profiles. Martens
hardness, Young’s modulus, and creep rates were calculated,shown in Table. The hardness and Young’s modulus of the a12
C:D film were smaller than those of other films. This result
were matched with suggestion of the containing of the polymerFig. 3. Indentation depth-load profiles.
like components from results of Raman measurements.
Although Martens hardness and Young’s modulus showed similar values on a-13C:H and a-12C:H films,
the creep rate indicated large difference. The creep rate suggests the deformation resistance, and isrelated to the vacancy size between the bonding networks. Domain size of sp2 bonded carbon network
is similar at these 2 films because ID/IG indicated little difference. One of the possible reasons for thisdifference is that the length of 13C-13C bond is shorter than that of 12C-12C bond.7). Hence, it is consideredthat the volume of vacancy between bonding networks of 13C-13C is smaller than that of 12C-12C.
4. Summary
To understand isotopic effects of amorphous carbon films, a-13C:H and a-12C:D films were deposited.
Results of structural estimation by XRR and Raman spectroscopy indicated structural difference.
Structure of a-12C:D showed large difference compared with other film because of the difference ofplasma condition. On mechanical properties, the creep rate of a-13C:H indicated a small rate comparedwith that of a-12C:H. This reasons for this difference is caused by the bonding length of 13C and 12C.
These results indicated that a-C:H films were affected by isotope effect by the change of 13C from 12C.
5. Acknowledgement
This work was supported by “Planting Seeds for Research” program in Tokyo Inst. of Tech. And someparts of this work supported by JSPS KAKENHI Grant Number 15K18038.
References[1] N. V. Novikov, A. P. Podoba, S. V. Shmegera, A. Witek, A. Zaitsev, A. B. Denisenko, W. Fahrner, M. Werner,
Influence of isotopic content on diamond thermal conductivity, Dia. Relat. Mater., 8 (1999), pp1602-1606.
[2] H. Watanabe, C. E. Nebel, S. Shikata, Isotopic Homojunction Band Engineering from Diamond, Science, 324(2009), pp1425-1428.
[3] A. T. Collins, S. C. Lawson, G. Davies, H. Kanda, Indirect energy gap of 13C diamond, Phys. Rev. Lett., 65(1990), pp891-894.
[4] F. Tuinstra, J. L. Koenig, Raman Spectrum of Graphite, J. Chem. Phys. 53 (1970), pp1126-1130.
[5] S. Karasawa, M. Mitsuhashi, S. Ohya, K. Kobayashi, T. Watanabe, K. Hirai, K. Horiguchi, F. Togashi, Crystal
growth of epitaxial CVD diamond using 13C isotope and characterization of dislocations by Raman spectroscopy,
J. Cryst. Grow. 128 (1993), pp403-406.
[6] H. Holloway, K. C. Hass, M. A. Tamor, T. R. Anthony and W. F. Banholzer, Isotopic dependence of the lattice
constant of diamond , Phys. Rev. B, 44 (1991), pp7123-7126.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

EVALUATION OF ADHESIVE STRENGTH BETWEEN Ti–29Nb–13Ta–4.6Zr ALLOYAND SOL-GEL FABRICATED HYDROXYAPATITE FILMS
Junko Hieda1*, Mitsuo Niinomi2, Masaaki Nakai2, Ken Cho2
1

Graduate School of Science and Engineering, Tokyo Institute of Technology
2Institute for Materials Research, Tohoku University* hieda@mech.titech.ac.jp

Keywords: -type titanium alloy, hydroxyapatite, adhesive strength, Ringer’s solution
1. Introduction
The -type Ti alloy, Ti–29Nb–13Ta–4.6Zr alloy (TNTZ), has been developed in order to achieve a
low Young’s modulus (~60 GPa) [1]. The Young’s modulus of TNTZ is lower than that of stainlesssteel (~180 GPa) and Ti-6Al-4V ELI (~110 GPa), which are conventional alloys for biomedical
applications. In addition, TNTZ consists of non-toxic and non-allergenic elements. Thus, TNTZ is avery suitable Ti alloy for implant devices. Although Ti and Ti alloys exhibit excellent hard-tissuecompatibility attributed to osseointegration, it takes a long time to achieve the remodeling of the boneon the surface of Ti and Ti alloys. Therefore, to improve the hard-tissue compatibility of Ti and Ti
alloys, various surface treatments have been developed for the formation of inorganic layers which
exhibit the desired high hard-tissue compatibility. The most common way for providing the highhard-tissue compatibility is coating of hydroxyapatite (HAp: Ca10(PO4)6(OH)2) on the surface of Ti
and Ti alloys, which is a component of human bone. The investigation of the effect of the bodyenvironment on the adhesive strength of HAp films fabricated on the surface of the metallic implant isvery important for biomedical applications. The adhesion between sol-gel fabricated HAp films and
TNTZ before and after immersion in Ringer’s solution was evaluated in this study.
2. Experimental procedures
TNTZ disks of 10 mm in diameter and 2 mm in thickness were cut from the TNTZ bar. The TNTZ
disks were wet-polished using waterproof abrasive papers of up to #4000 grit and then buff-polishedwith a colloidal silica dispersion. Some of the mirror-polished TNTZ disks were polished using awaterproof abrasive paper of #600. These TNTZ disks are denoted by mechanically polished TNTZ
disks hereinafter. The surfaces of mirror-polished and mechanically polished TNTZ disks wereobserved using an atomic force microscope. HAp films were fabricated on mirror-polished andmechanically polished TNTZ disks by a sol-gel process [2]. The phase constituents of the films wereidentified using an X-ray diffraction (XRD). Some of the samples were immersed in Ringer’s solution
at 310 K for 7 d in an incubator.
The adhesive strengths of HAp films fabricated on mirror-polished and on mechanically polished
TNTZ disks were evaluated by adhesion tests according to ASTM F 1147-05. The samples were fixed
between stainless steel rods (10 mm) at an applied load of 0.138 MPa using a
polymethylmethacrylate-based commercial dental adhesive. The samples and rods were set in theInstron-type testing machine with a maximum load capacity of 20 kN and pulled vertically with across-head speed of 4.17 × 10-5 m∙s-1. The tensile strength of the adhesive, which indicates the upper
limit of the present adhesion tests, was also measured by fixing the mechanically polished TNTZ disk
between the rods using this adhesive. The result is 29.7 ± 3.5 MPa (around 30 MPa). A statisticalanalysis of the adhesive strength was carried out by means of Student’s t-test. After the adhesion test,
the fracture surfaces of both sides of the samples were observed by scanning electron microscopy(SEM).
3. Results and discussion
The mirror-polished TNTZ disk exhibited a smooth surface and its average surface roughness wasapproximately 16 nm. The surface of the mechanically polished TNTZ disk had uniform asperities
with heights of less than 1 m. The average surface roughness of the mechanically polished TNTZ

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disk was approximately 187 nm.
In a typical XRD pattern of the HAp film fabricated on the mirror-polished TNTZ disk by thesol-gel process, the peaks attributed to HAp were present at around 26° and 32–36°. Thus, it was
confirmed that HAp was successfully formed on the TNTZ disk. Similar HAp was also obtained in thecase of fabrication on the mechanically polished TNTZ disks.
Figure 1 shows the adhesive strengths of HAp films fabricated on the mirror-polished and themechanically polished TNTZ disks before and after immersion in Ringer’s solution at 310 K for 7d.
As stated above, the average tensile strength of the adhesive was around 30 MPa, which represented
the limit value of the adhesive strength measurable by this test. The average adhesive strengths of the
HAp films fabricated on the mirror-polished TNTZ disk before and after immersion in Ringer’s
solution were about 16.0 MPa and 21.5 MPa, respectively. On the other hand, the average adhesivestrengths of the HAp films fabricated on the mechanically polished TNTZ disks before and after
immersion in Ringer’s solution were about 23.2 MPa and 21.2 MPa, respectively. Therefore, thesevalues showed no significant differences (p > 0.05). Thus, there was no difference in the adhesivestrength of HAp films with different surface morphology. Moreover, the adhesive strength exhibited
similar values before and after immersion in Ringer’s solution at 310 K for 7 d. In order to evaluate
the adhesion of HAp films deposited on the mirror-polished and the mechanically polished TNTZ
disks, the fracture areas at the interface between HAp and TNTZ were measured on the basis of SEM
images before and after immersion in Ringer’s solution as shown in Fig. 2. The average fracture areasat the interface between HAp films and mirror-polished TNTZ disks before and after immersion in
Ringer’s solution were about 3.9% and 23%, respectively, revealing a significant increase after
immersion in Ringer’s solution. No fracture at the interface between HAp film and the mechanicallypolished TNTZ disk was observed, both before and after immersion in Ringer’s solution. These results
indicated that the HAp film on the mechanically polished TNTZ disk exhibited greater adhesion incomparison with the HAp film on the mirror-polished TNTZ.

Figure 1 Adhesive strengths of HAp films
fabricated on mirror-polished and mechanicallypolished TNTZ disks.

Figure 2 Average fracture areas at interface
between HAp films and mirror-polished andmechanically polished TNTZ disks.

References[1] D. Kuroda, M. Niinomi, M. Morinaga, Y. Kato, T. Yashiro, Design and mechanical properties of new -type
titanium alloys for implant materials, Mater. Sci. Eng. A., 243(1998), pp. 244-249.
[2] L.H. Li, H.W. Kim, S.H. Lee, Y.M. Kong, H.E. Kim, Biocompatibility of titanium implants modified by
microarc oxidation and hydroxyapatite coating, J. Biomed. Mater. Res. A., 73(2005), pp. 48-54.

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ADHESION BEHAVIOR OF MECHANICALLY JOINED PARTS OF
ALUMINUM ALLOYS
Ruben Septianus1, Murugesan Jayaprakash1, Yukio Miyashita*1, Tetsuri Yamada2,
Kiyoshi Shirato2
1

Department of Mechanical Engineering, Nagaoka University of Technology,1603-1 Kamitomioka, Nagaoka, Niigata 940-2188, Japan
2
Isuzu Motors Limited*Corresponding Author: miyayuki@mech.nagaokaut.ac.jp, Tel: +81-258-47-9704

Keywords: Aluminum alloy, Mechanical joint, Adhesion, Contact damage
1. Introduction
Aluminum alloys have been widely used in automotive applications, especially in automotiveengine components due to their material characteristics, such as high deformability and castability,
appropriate strength, light weight and good corrosion resistance, etc.[1]. In automotive engine, many
parts are clamped together by using bolted joints.
Material transfer between two contact surfaces and formation of wear debris occur at the joining
part due to engine vibration under the service condition, which cause surface adhesion and adhesive
wear that reduces service life of the component. Therefore, it is important to understand themechanism of adhesion behavior at the contact part in aluminum alloy. In the present study, theadhesion behavior between cast aluminum alloys was investigated.
2. Experimental Procedures
Figure 1 shows the microstructures of ADC12 and AC4C. A plate specimen of AC4C was clampedwith a pad specimen of ADC12 as a bolt- nut assembly as shown in Figure 2. Both specimen and pad
surfaces were polished until mirror surface. A M10 hexagonal bolt was prepared with attaching strain
gage to measure the clamping force.
A servo hydraulic fatigue testing machine was used to apply cyclic loads.The experiments wereconducted by varying the contact pressures (50, 60, and 70 MPa) and cyclic stress amplitude (15, 20,
and 25 MPa). Cyclic loading was applied at frequency of 20Hz with 105 cycles for each test.
3. Results and Discussions
Figure 3 shows the adhesive conditions for ADC12 and AC4C mechanical joint. It was observed thatthe strong adhesion (when bolt is removed, the pad still firmly stick to the specimen) occurred at lower
stress amplitude of 15 and 20 MPa with relatively high contact pressure of 70 MPa. However, theweak adhesion (when bolt is removed, the pad still stick to the specimen, but after a few minutes it
separate by itself) occurred at stress amplitude of 25 MPa with contact pressure of 70 MPa. Weakadhesion also occurred at stress amplitude of 15 and 20 MPa with lower contact pressure of 60 MPa.
The adhesion did not occur at all below contact pressure of 60 MPa.
Specimen surface was observed by laser scanning microscope to study change in roughness before
and after the test. Figure 4 (a) and (b) show the surface damages after the test conducted at stressamplitude of 15 MPa with contact pressure of 70 MPa for ADC12 and AC4C, respectively. It can beobserved that AC4C specimen surface has more severe damage compared to ADC12 pad surface.
Cyclic load

(b)

(a)

AC4C
ADC12

Bolt

200mm

ADC12

Nut

200mm

Figure 1 Microstructures of (a) ADC12 contact pad and (b) AC4C specimen.

Figure 2 Mechanical joint
tested.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

30

(a)

(b)

25

20

15

Strong adhesion

10
40

Weak adhesion5060

No adhesion
70

80

cyclicloading

500 µm500 µmContact pressure , MPa

Figure 3 Adhesive conditions for ADC12 andAC4C mechanical joint.
4000nmFigure 4 (a) ADC12 and (b) AC4C surfaces after the test
carried out at stress amplitude of 15 MPa with contactpressure of 70 MPa.

before3000nm2000nm

100 µm1000nm

4000nmafter

3000nm2000nm1000nm

100 µmadhesive region

interface(a)ADC124000nmbefore3000nm2000nm1000nm

100 µm

4000nmafter
3000nm2000nm1000nm

100 µm(b)AC4C
Figure 5 Surface profile of (a) ADC12 and (b) AC4C
before and after the test carried out at stress amplitudeof 15 MPa with contact pressure of 70 MPa.

100 µmFigure 6 Cross section of ADC12 specimen
after the test carried out at stress amplitude of
15 MPa with contact pressure of 70 MPa.

Surface profiles measured are shown in Figure 5. The roughness increased significantly after the
test. Figure 6 shows cross section of adhesive region in ADC12 after the test. It can be observed thatthe metal transfer was adhered to the ADC12 contact pad surface.
The clean metallic surfaces have high surface energies and make strong metallic bonds when it
contact[2]. The initial contact by clamping force may lead the plastic deformation of the surfaces
asperities, which resulting adhesive bonds among the asperities[3]. When the contact pressure is highenough, the adhesion will occur.
Furthermore, when low cyclic stress amplitude is applied, the localized cold welding may occur
among the asperities and resulting a strong bonding. However, when higher cyclic stress amplitude isapplied, the relative slip movement between two surfaces become larger and adhesive wear will takeplace. In this case, the adhesion becomes weaker or may not occur.
4. Summary
Adhesion successfully occurred in the present experiment proposed. The strong adhesion occurredin the test conducted at relatively high contact pressure with lower cyclic stress amplitude. The
adhesive did not occur at all below a certain contact pressure.
References[1] Dwivedi DK, Adhesive wear behaviour of cast aluminium–silicon alloys: Overview, Materials and Design,Vol. 31 (2010), pp. 2517-2531.
[2] Bharat Bhushan, Modern Tribology Handbook, Vol. 1, CRC Press LLC, Florida (2001), pp. 195 – 213.
[3] J.Halling, Principles of Tribology, The Macmillan Press Ltd., London(1975).

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

DEVELOPMENT FOR HYDROGEN STORAGE Mg-Ti SOLID SOLUTION BY
MECHANICAL ALLOYING
Naoki MAEDA1*, Mitsuo NOTOMI1
1

Department of Mechanical Engineering, Meiji University*ce52061@meiji.ac.jp

Keywords: Hydrogen storage, Mg-Ti solid solution, Mechanical alloying, Stearic acid,Transformation hcp into fcc
1. Introduction
Although hydrogen, which has high energy density and environmentally-friendly properties, is oneof the alternative sources of energy to fossil fuel, its currently stored methods such as liquefaction andcompression have safety issues because these require extremely low temperature or high pressure to
store it. Metal hydrides have the advantage of moderate operating temperatures and pressures,compared to them, however it must need to be further improvement to put into practical use.
Mg is one of the most promising candidates for hydrogen storage material due to low cost and highhydrogen storage capacity by weight (7.6 wt.%, MgH2). The main defect of Mg is that its hydride
MgH2 is low (de)hydrogenation kinetics because MgH2 has high thermodynamic stability and there issome possibility to improving the defect by synthesizing with transition metals like Sc, Ti, V, Co. Ti islightweight (4.51 g/cm3) and low cost that is advantage for hydrogen storage and it is difficult tosynthesize the Mg-Ti solid solution by melting methods due to large difference of melting points
between Mg (923 K) and Ti (1943 K). The mechanical alloying (MA) which is a solid-state powder
processing method is capable of synthesizing such materials. Asano et al. [1] and Maweja et al. [2] had
successfully synthesized Mg-Ti solution by MA and it also had made under severe milling conditions
such as long milling time (more than 200 h) or high rotation speed (more than 400 rpm). The additionof process control agents (PCA) is one of the methods to enhance the MA process efficiency. It is tomodify the surface condition of deformed particles by impeding the clean metal to metal contact
necessary for cold welding, and promote powder fracturing [3].
The present study aims to synthesize Mg-Ti solution by MA with stearic acid as a PCA andinvestigate its hydrogen absorption/desorption properties.
2. Experimental
MA was performed with the planetary ball mill (LP-4, Ito Seisakusho) Powder and milling ballswith a diameter of 10 mm were set in milling pot whose internal volume was 500 cm3 and it was filledwith Ar gas. The milling balls and pot were made of chromium steel. Table 1 summarizes the milling
conditions of all the samples in this study. Sample III was conducted hydrogenation at 573 K and 2.5
MPa with hydrogenation equipment after heat treatment at 673 K to remove stearic acid.
The X-ray diffraction (XRD) was measured with the X-ray diffractometer (MiniFlex600, Rigaku
Corporation) with Cu Kα radiation. The lattice parameters of the samples were calculated from the
XRD peak positions. Desorption temperature was determined on detecting the exothermic reaction
with the differential scanning calorimetry (DSC-60, Shimadzu Corporation).

Sample No.
Composition
Amount of stearic acid
Ball filling rate
Rotation speed
Milling time
Powder-to-ball rate―
[wt.%]
[%]
[rpm]
[h]
―

Table 1 Milling conditions.
III
Mg-30at.%Ti
Mg-30at.%Ti
10
―
10
10
200
200
70
70
1:70
1:70

III
Mg-50at.%Ti
10
10
200
70
1:70

IV
Ti
10
10
200
40
1:70

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

3. Results and discussion
Figure 1(a)-(d) show the XRD patterns of milled sample I-IV respectively. The XRD peaks of pure
Mg or Ti are visible in Figure 1(a), (b) and (d), while these can’t be seen in Figure 1(c). The peaks
marked by red solid circle are at the same positions as the peaks of the Mg-Ti solution with fcc phasewhich were reported by Asano et al. [1].
3.1 Mechanism of synthesizing Mg-Ti solution
As shown Figure 1(d), Ti was transformed hcp into fcc phase by milling with stearic acid, and at thetime of the experiment, we observed some powder of sample IV was suddenly combusted and formed
TiO2 as soon as we opened the milling pot that had been finished milling, and hence the transformed
Ti seems metastable. These results suggested that the Mg-Ti solution was synthesized by dissolving
Mg in metastable Ti with fcc phase.
3.2 Hydrogen absorption/desorption properties of Mg-Ti solution
Figure 2(a) and (b) show the XRD pattern and the DSC results of sample III after hydrogenation
respectively. As shown Figure 2(a) and (b), a part of the Mg-Ti solution was formed MgH2 byhydrogenation and it desorbed hydrogen at approximately 637 K. The desorption temperature wasn’t
improved from pure MgH2 because the formed MgH2 wasn’t dissolved Ti, and consequently wasn’t
achieved destabilization.
4. Conclusion
The Mg-Ti solution was synthesized by milling Mg and Ti powder with stearic acid and dissolving
Mg in metastable Ti that was transformed hcp into fcc phase. It was formed MgH2 by hydrogenation
which wasn’t dissolved Ti and wasn’t achieved destabilization. Its desorption temperature was at
approximately 673 K which wasn’t improved from pure MgH2.
References[1] Kohta Asano, Hirotoshi Enoki, Etsuo Akiba, Synthesis of HCP, FCC and BCC structure alloys in the Mg-Ti
binary system by means of ball milling, Journal of Alloys and Compounds, 480(2009), pp. 558-563[2] Kasonde Maweja, Maje Phasha, Nic van der berg, Microstructure and crystal structure of an equimolar
Mg-Ti alloy processed by Simoloyer high-energy ball mill, Powder Technology, 199(2010), pp. 256-263[3] Li Lu, Y.F. Zhang, Influence of process control agent on interdiffusion between Al and Mg during mechanicalalloying, Journal of Alloys and Compounds, 290(1999), pp. 279-283(a) :Mg

Sample I : Mg-30at.%Ti

Intensity (a. u.):Ti(hcp)(b)

Sample II : Mg-30at.%Ti+stearic acid

:Mg-Ti(fcc)(c)

Sample III : Mg-50at.%Ti+stearic acid(d) :Ti(fcc)

Sample IV : Ti+stearic acid

:TiO2

30

(a)

20

40

50
60
Diffraction angle，2θ [deg]Figure 1 XRD patterns of sample I-IV.

:Mg-Ti (fcc)
:MgH2

30

40
50
60
Diffraction angle，2θ [deg]Figure 2

70

Heat flow (a. u.)

Intensity (a. u.)

20

70

80

(b)

80 400:Before hydrogenation:After hydrogenation

500

600
Temperature [K]700

XRD pattern and DSC results of sample III (after hydrogenation).

Page 78

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

DEVELOPMENT OF DAMAGE DETECTION TECHNIQUE FOR SHEET-TYPE
LITHIUM-ION BATTERIES USING ACOUSTIC EMISSIONS
Hiroaki ISHII1*, Takuma MATSUO2
1

Graduated Student, School of Science and Technology Meiji University
2
Lecturer, School of Science and Technology Meiji University* ce52010@meiji.ac.jp

Keywords: Lithium-Ion Battery, Degradation, Acoustic Emission, Health Monitoring, Optical Fiber
SensorLithium-ion batteries have the advantages of a high energy density and reduced memory
effect. They are widely used for systems such as portable devices. However, there have been someincidents due to safety problems. Therefore, a technique for detecting damage at an early stage isneeded. We applied the acoustic emission (AE) technique in order to detect the degradation anddamage in lithium-ion batteries. In this study, sheet-type lithium-ion batteries were monitored with an
optical fiber sensor because optical fiber sensors are explosion-proof and allow for free design of thesensor shape [1].
The AE generation in sheet-type batteries was first characterized with a conventional AE sensor.
Figure 1 shows experimental setup. The AE sensors were mounted onto the surface of the batteries; it
was constructed from an anode electrode, a cathode electrode, a separator, and a polypropylene plate
to prevent deformation of the batteries. The charging operation was carried out at a constant current of
4.16 mA, and the discharging operation was carried out at a constant current of -4.16 mA. AE signals
were monitored by two sensors (PAC, R50α) and mounted onto the surface of the batteries. The
detected AE signals were first amplified (40 dB) and filtered (10 kHz high-pass filter) using a
pre-amplifier and filtered (500 kHz low-pass filter) again and monitored with a digitizer (Alazartech,
ATS330). The potential and current were controlled by a potentiostat (Bio-Logic, SP-50). The results
obtained with a conventional AE sensor with normal charge/discharge cycles are presented below.
First, AE signals were mainly detected when the potential was over 80% full. Second, the peakfrequency of the AE signal was about 150 kHz. Finally, according to an analysis of the frequency
spectra from the AE waveforms, the detected AE signals were caused by gas evolution [2]. Then, we
repeated the charge/discharge cycles. After many cycles, the AE signals increased when the potential
did not increase. We confirmed that damaged batteries generated AE signals at random intervals. The
AE from the damaged batteries was 30–50 events, which was twice as many that from the normal
cycles. The peak frequency of the AE signal was about 50 kHz, which changed about 150 kHz during
the normal charge/discharge cycles.

Fig.1 Experimental setup for monitoring AE signals during charge-discharge cycles utilizing PZT sensors

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Then, we applied an optical fiber sensor. Experimental setup was the same shown in Fig.1,
however optical fiber sensor was used. Figure 2 shows the optical fiber AE monitoring system used inthis study. The optical fiber AE monitoring system was composed of an optical fiber sensor unit, a
feedback circuit, and a digitizer. The digitizer recorded the AE signals detected by the sensor unit, andthe feedback circuit improved the stability of the sensor. Optical fiber was wound over the aluminum
cylinder and the sensor was set on the sheet-type battery. The charging operation was carried out aconstant current of 0.8 mA, and the discharge operation was carried out a constant current of -0.8 mA.

Fig.2 Optical fiber AE monitoring system used in this study

Figure 3 shows the potential and cumulative AE events as a function of time with the optical fiber
sensor and the AE waveform the AE signals could be also detected with the optical fiber sensor when
the signal-to-noise ratio was 30–35, as in the results obtained with the conventional AE sensor.

Fig.3 Potential and cumulative AE events as a function of time (left) and the AE waveform (right)

Acknowledgment
This study was partially supported by the JGC-S Scholarship Foundation, grand for young researcher.
References[1] Matsuo, T., Sano, K., Sakakibara, Y., and Nakayama, G., Estimate of Stress Corrosion Cracking Initiation and
Propagation in High-Pressure, High-Temperature Water Environment Utilizing Acoustic Emission, Materials
Transactions of Metals and Matterials, Vol.53 No.3(2015), pp.327-332[2] Matsuo, T., Uchida, M., and Cho, H., Development of Acoustic Emission Clustering Method to Detect
Degradation of Lithium Ion Batteries, Journal of Solid Mechanics and Materials Engineering, 5-12(2011), pp.
678-689.

Page 80

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

IMPACT RESISTANCE OF THE GLASSES INSTALLED
NANO-LAMIRATED FILMS AND ARAMID FIBER CLOTH
Yudai YASUMOTO1*, Hikaru MASUI1, Mitsuo NOTOMI1 Katsumasa AOKI2
1

Department of Mechanical Engineering, Meiji University
2
SOU・SHOW CO.,Ltd.
*ce52067@meiji.ac.jp

Keywords: Glass，Nano-laminated film，Aramid fiber cloth，Impact fracture，Strain measurement
1. Introduction
Glass, which has excellent workability, production efficiency and the properties of airtightness,transparence, corrosion resistance and thermochemical stability, is used in various situation such as thewindows of buildings and cars. The strength of the glass depends on the number of Griffith flaws,
which is generated during the manufacturing process at the surface of the glasses and is less than its
ideal strength without any cracks[1]. In addition, if the stress over the limit of elasticity is subjected to
the glass, brittle fracture is occurred instantaneously without plastic deformation, and consequently
glass has a problem for impact resistance. There are many kinds of methods to reinforce glass andrecently, much attention has been paid to the method of film installation on the surface of the glass.
In this study, we considered the improvement of impact resistance of glass installed nano-laminatedfilm (Super layer URTRA600, 3M Japan Limited) and aramid fiber cloth (AW40/40, MAEDAKOSEN
CO., LTD.) over its surface.
2. Experimental procedure
In this research, five sheets of glass (300×300×3.2 mm) are used as specimens (Table 1). GLASS is
the specimen a glass without film and cloth. FILM-100 and ARAMID-100 are installed film or cloth(100×100) on impact surface. FILM-ALL-IMPACT and FILM-ALL-NON are installed film
(300×300) on impact or non-impact surface. It takes a month for curing period of glass installed film
or cloth. Their impact resistances were evaluated with the free-fall drop test mostly according to JIS R
3212. In this test, the steel ball (80 mm diameter, 2 kg weight) falls freely 2 meters high from thespecimen. Fracture behaviors of all specimens are taken by high-speed camera (EX-FH25, CASIOCOMPUTER CO., LTD.), whose flame rate is 240 flames per second, from the impact side of thespecimen. Specimen’s circumference strain is measured with strain gauge (FLA-1-11, TOKYO KEIKI
INC.). A center is set as the contact point between the specimen and the ball and the virtual circle linesof which each radius are 80 and 110 mm are drawn on the impact surface. All strain gauges areattached on parallel to the lines so that circumference strain is measured. The signals from the gages
translate to the bridge box (SB-121A, Tokyo Sokki Kenkyujo Co., Ltd.), dynamic strain meter(DC-97A, Tokyo Sokki Kenkyujo Co., Ltd.) and oscilloscope (DSO1024A, Agilent Technologies Inc.)in order. The range of oscilloscope is in 1 millisecond from the time the ball impacted to thespecimens.
Table 1 Specimen conditions.
Specimen name
GLASS
FILM-100
ARAMID-100
FILM-ALL-IMPACT
FILM-ALL-NON

Material

Installed material

Installed area

Installed side××

Glass(300×300 mm)

×Nano-laminated film
Aramid fiber sheet
Nano-laminated film

100×100
300×300

Impact side
Non-impact side

3. Results and Discussions
Figure 1 shows the specimens after fractured. Figure 1 (a) shows the specimen fractured perfectly
and that is penetrated by the ball. In this study this type of fracture is called All-Fractured (AF). Figure
1 (b) and (c) show the specimens fractured are also penetrated by the ball and the areas of film or cloth
installed however hold the plate shape even cracked. The type of fracture is called Partially-Fractured(PF). Figure 1 (d) and (e) show the specimens are not penetrated by the ball with a lot of cracks. The
type of fracture is called Non-penetration with cracks (NPC).

Page 81The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

The resistance of impact of FILM-ALL-IMPACT and FILM-ALL-NON is improved by the
installed films or cloth. This is because the impact energy diffusions caused by the film or cloth are
occurred in the installed areas. In addition, the glass fragments are not scattered and remained in thearea installed film or cloth and these fragments cause friction each other, and the areas remain the plate
shape even cracked. In the case of PF, the load by the impact ball is concentrated to the boundary
between the glass and the film or cloth, and consequently these specimens are penetrated by the ballwith the areas installed film or cloth held the plate shape.
Figure 2 shows the circumference strain profile in 1 millisecond from the time the ball impacted. In
this graph the vertical axis corresponds to the circumference strain of the specimen and the strain iscompressive when the vertical value is negative. Figure 2 (a)-(e) shows the strains of all specimens
repeat increase phase and decrease phase in compression. Table 2 shows the maximum strains of thespecimens defined minimum value of the strain profile in Figure 2. The maximum strains of
FILM-100, ARAMID-100, FILM-ALL-IMPACT and FILM-ALL-NON are lower than that of GLASS
in the 80 mm from the center of the specimen. In Table 2 the maximum strains in 80 mm are inclined
to decrease by installing film or
cloth but that of FILM-100,
ARAMID-100 and FILM-ALLIMPACT are almost the samevalue in spite of the differences of
their condition. It revealed thatthe maximum circumference strain
is irrelevant to the penetration(a) GLASS (AF)(b) FILM-100 (PF)property.

(c) ARAMID-100 (PF)(d) FILM-ALL-IMPACT (NPC)(e) FILM-ALL-NON (NPC)
Figure 1 Specimens at the time of fractured

0.001

0.001

0.001

0

0

-0.001-0.001Time [ms]

1

0.002

0.002

0.001

0.001

0

-0.001

1(b) FILM-100 (PF)

80mm
110mm
1

-0.0020

1Time [ms]Table 2 Experimental results-0.001Time [ms]-0.0020(c)ARAMID-100 (PF)

0

80mm
110mm
-0.0020Time [ms]

80mm
110mm

(d) FILM-ALL-IMPACT(NPC)

Strain [Strain [(a) GLASS (AF)-0.00200-0.001

80mm
110mm

80mm
110mm
-0.0020Strain [0.002

Strain [0.002

Strain [0.002Time [ms](e) FILM-ALL-NON (NPC)

Figure 2 Circumference train profile in 1 millisecond

1Specimen

Fracture
mode

G
F-100
A-100
F-A-I
F-A-N

AF
PF
PF
NPC
NPC

Maximumcompressive strain
×10-3[]
80mm
110mm
-1.42
-0.94
-1.08
-0.8
-1.04
-0.88
-1.08
-0.74
-1.22
-0.8

References[1] Hisashi HARAMORI, Relations between fracture strength of glass and Griffith flaws, The Japan Society
for Precision Engineering, Volume 31(1965), pp. 807-812.

Page 82

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

INFLUENCE OF ANNEALING ON SHAPE MEMORY PROPERTY IN
FE-MN-SI BASED ALLOY
Hiromu ISHII*, Naohiro OKUBO, Mitsuo NOTOMI
Department of Mechanical Engineering, Meiji University*ce52011@meiji.ac.jp

Keywords: Fe-Mn-Si based alloy, Shape memory effect, Stress-induced martensitic transformation,
Annealing, EBSD analysis
1. Introduction
Fe-Mn-Si based shape memory alloys (SMAs) have attracted much attention among the numerous
SMAs because these SMAs have many advantages such as high mechanical properties, fineworkability and cheap cost. Although they have been expected to be utilized as large-scale products,
their shape memory property is low compared with Ni-Ti alloys which have been generally used inmany engineering fields, so that it has been strongly required to be improved. It is known that shapememory effect (SME) in Fe-Mn-Si based alloys is associated with stress-induced γ ↔ ε martensitic
transformation, but detail processes for improvement have been widely studied and discussed. Li et al.
[1] reported that the property of the alloy was enhanced for proper annealing before deforming. The
objective of this study was to clarify influences of annealing on phase microstructure, crystal structureand shape memory behavior of Fe-28Mn-6Si-5Cr alloy.
2. Experiments
The chemical composition of the alloy used for this study is listed in Table 1. The alloy which is awire (φ1.2 mm) was manufactured by Awaji Materia Co., Ltd. and the process was cold-drawing
without subsequent solution treatment. In this study, heat treatment was conducted on the conditions
listed in Table 2, and the as-received alloy and the alloys annealed at 673 K, 773 K, 873 K, 973 K,
1073 K, 1173 K and 1273 K for 60 min were named specimen # 1, # 2, # 3, # 4, # 5, # 6, # 7 and # 8,
respectively. The outline of the experiments was shown in Figure 1. In order to evaluate shapememory behavior, all specimens were bent by material testing machine (Autograph AGS-H,
Shimadzu Corporation) with a bending punch and a jig on the condition listed in Table 3, then heated
at 923 K for 15 min as a recovery treatment. Electron backscatter diffraction (EBSD) and X-ray
diffraction (XRD) methods were employed to investigate the microstructural characterizations.
Samples for EBSD, parts of the cut specimens embedded into resin, were mechanically polished until
the specimens were halved, then the analyzed surfaces were prepared by colloidal silica nanoparticles.
EBSD analysis was conducted at 20 kV acceleration voltage with a step size of 0.955 μm on a fieldemission gun, and scanning electron microscopy (SEM JIB-4601F, JEOL Ltd.) equipped with
TSL-OIMTM EBSD set-up was used for the analysis. XRD analysis was conducted with XRD
instrument (MiniFlex600, Rigaku Corporation) using Cu-Kα radiation (λ = 0.154178 nm).
Table 1

Chemical composition of the alloy(mass %)

Mn
27.2
N
0.012

Si
5.69
S
< 0.009

Table 2Heat treatment condition in eachspecimen

Table 3Specimen
#1
# 2
# 3
# 4
# 5
# 6
# 7
# 8

Heat treatment condition
No treatment
673 K × 60 min, Furnace cooling
773 K × 60 min, Furnace cooling
873 K × 60 min, Furnace cooling
973 K × 60 min, Furnace cooling1073 K × 60 min, Furnace cooling
1173 K × 60 min, Furnace cooling
1273 K × 60 min, Furnace coolingSpecimen length
Test type
Maximum load
Minimum load
Test velocity
Span
Bending radius
Hold time
Temperature

Cr
5.02
P
< 0.003

C
0.047
O
0.002

Ni
0.033

Heat treatment(Annealing)

Bending test

Shape recovery
treatment

XRD analysis and
EBSD analysis*
EBSD analysis*(Compressive and tensile region)* Specimen # 1, # 2, # 4, # 6, # 8 were investigated.

Figure 1

Outline of the experimental methods
Bending test condition
50 mm
3-point bending
80 N
0N
15 mm/min
31 mm
30 mm
3.0 sec
Room temperature

Page 83

7KHWK$VLDQ6\PSRVLXPRQ0DWHULDO3URFHVVLQJ $603 /RPERN$XJXVW

3. Results and discussion
Table 4 shows recovery angles of the specimens after unloading and the recovery treatment, θSB andθSME, and also shows recovery ratios in SME, ηSME, which was calculated with a following equation.
ηSME =

θSME
�180 - θmax� - θSB× 100 [%]

(1)

Where θmax is 65 deg., a deformation angle in maximum loading. The results indicated that specimen #4 and # 5 had the most recovery in SME. Table 5 shows phase distribution maps obtained by EBSD
analysis, and it was found that allover austenite phases in specimen # 4 mostly transformed into εmartensite over the bending deformation. In contrast, specimen # 2 was formed few austenite phasesby the annealing. In addition, the high temperature caused grain coarsening in specimen # 8, so it
noticeably decreased the elastic recovery. Figure 2 shows quantitatively a relationship between ηSME
and the phase transformations. A Mq/Aq is a quantitative ratio of martensite to austenite in eachanalyzed area, and the high Mq/Aq means that there are many martensite phases in a specimen. Figure 2
suggested ε martensitite formed by bending led to fine SME. On the other hand, Figure 3 shows arelationship between ηSME and average crystallite size, D which was obtained by XRD analysis．D wasmost enlarged by annealing at 973 K, and that suggested the dislocation density was most decreased
with the temperature [2]. Besides, the variation trend of D was similar to that of ηSME and then therelationship implied the decrease of dislocation density encouraged γ ↔ ε martensitic transformation.
4. Conclusions
1. Shape memory property of the alloy was enhanced by annealing at 873-973 K, which formed
allover austenite phases without grain coarsening in the alloy.
2. The main factor in the fine SME was that allover austenite phases generated by the annealing
temperatures mostly transformed into ε martensite over the bending deformation.
3. The decrease of dislocation density was caused by relatively high annealing temperatures, thatwould be one of the factors in encouraging the γ ↔ ε martensitic transformation.
References

#1
28
14
16.1

#2
49
18
27.3

#3
51
35
54.7

#4
41
62
83.3

#5
42
61
83.6

#6
36
36
45.6

#7
15
35
35.0

#8
10
28
26.7

Table 5 Phase distribution maps before and after
bending deformation (blue: austenite，yellow:
martensite) obtained by EBSD analysis
Before bending
Optional region

After bending
Compressive region

Tensile region

20μm

20μm
[%]

100

#2

2.5
2.0

46

1.5
27

1.0
0.5

27

16

0.0
As-received

673 ℃
K
873 ℃
K
1073℃K
800
400
600
Annealing temperature

20μm

#4
20μm

20μm

20μm

#6

Recovery ratio in SME,

20μm

90

60

ηSME
D

50

70

40

60
50

30

40
20

30
20

10

10
0

20μm

20μm

20μm

90
80
70
60
50
40
30
20
10
0

1273 ℃
K
1000

80

20μm

[%]

Before bending
After bending
ηSME

84

3.0

Figure 2 Mq/Aq before and after bending
deformation and the variation of ηSME in theannealing temperatures

#1
20μm

3.5

Recovery ratio in SME,

SpecimenθSB [deg.]θSME[deg.]
ηSME [%]

Average crystallite size, D [nm]Table 4 Recovery angles by spring back and
SME and recovery ratios in SME

Quantitative ratio of martensite to austenite,
Mq /Aq[1] Li. HJ, Dunne. D, Kennon. N, Factors influencing shape memory effect and phase transformation behaviourof Fe-Mn-Si based shape memory alloys, Materials Science and Engineering A, Vol. 273-275, No. SI (1999),
pp. 517-523.
[2] Ting C. Huang, X-ray Polycrystalline Diffraction Analysis of Thin Films, Journal of the Crystallographic
Society of Japan, Vol. 34, No. 2(1992), pp. 110-117.

0

As-received 673 K

773 K

873 K

973 K

1073 K 1173 K 1273 K

Annealing temperature

#8
20μm

20μm

20μm

Figure 3 The Variations of ηSME and D in
the annealing temperatures

3DJH4

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

FRICTION STIR INCREMENTAL FORMING OF ALUMINUM ALLOY SHEETSWITH BACK SUPPORT DIE
Masaaki Otsu1,＊, Wei Jiang2, Takayuki Muranaka3, Ryo Matsumoto4, Masato Okada5
1
Professor, University of Fukui, 3-9-1 Bunkyo, Fukui, 910-8507, Japan
Graduate Student, University of Fukui, 3-9-1, Bunkyo, Fukui, 910-8507, Japan
3
Associate Professor, Fukui National College of Technology, Geshi, Sabae, 916-8507 Japan
4
Associate Professor, Osaka University, 2-1 Yamadaoka, Suita, 565-0871, Japan
5
Lecturer, University of Fukui, 3-9-1 Bunkyo, Fukui, 910-8507, Japan*Corresponding Author: otsu@u-fukui.ac.jp
2

Keywords: Friction Stir Incremental Forming, Sheet Metal Forming, Aluminum Alloy Sheet, Back
Support Die, Forming Accuracy
1. Introduction
Friction stir incremental forming (FSIF) has advantages of dieless forming and high formability for
non-ferrous metals. However, forming accuracy is not so high. In conventional incremental sheet
metal forming, back support die is used for improving the forming accuracy. In this study, a backsupport die is employed in friction stir incremental forming and formed shapes with and without backsupport die were compared.
2. Experimental method
The appearance of the working platform, die and tool were shown in Figure 1. The diameter of theforming tool was 6 mm and the top shape of the tool was hemispherical. AISI A5052 aluminum alloy
sheets were used for specimen. The size of the sheet was 100 mm x 100 mm and the thickness was 0.5
mm. During the forming, the sheet was clamped on the upside of the die using a screw. The tool
rotation rate ω feed rate v and wall angle θ were changed and formability was investigated. The tool
path was shown in Figure 2. The maximum height of pyramid, which is the objective shape, was 20
mm.

Figure 1 Appearance of working platform

Figure 2 Tool path for forming

3. Experimental result and discussion
3.1 Classification of formed parts
The formed parts can be classified into four types. Type A is forming succeeded without any defects.
Type B is forming was succeeded but burr was observed. Type C is forming was succeeded but theflange was curled up. Type D is forming failed with fracture. The formed specimens of each type wereshown in Figure 3.
3.2 Formable working conditions
The feed rate of the tool was fixed, and formable wall angle was studied with changing the tool
rotation rate. The experimental results when the tool feed rate was v = 1000 mm/min and v = 4000

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

mm/min were shown in Figure 4. In this figure, circle marks indicate that forming was completed until
forming height became to 20 mm. Cross marks mean the sheet was broken during forming. Triangle
marks mean that forming machine stopped by overloading.
As the tool rotation rate increased, formable wall angle became smaller and this means theformability was improved. However, as the tool feed rate increased, the minimum formable tool
rotation rate was increased and the minimum formable wall angle was increased and the formability
was decreased.
3.3 Cross-sectional shape
Cross-sectional shapes of ideal shape and formed parts with and without dies when the wall angle
was θ = 45 º and , tool feed rate was v = 1000 mm/min and tool rotation rate was ω = 10000 rpm, waswere plotted in Figure 5. Although there were no significant difference between the shapes at conicalsurface with and without die, those at the corner and flange were different and that with die was closer
to the ideal one.

(b) v = 4000mm/min

Figure 4 Relationship of wall angle and rotation rate

Ideal shape
FSIF with die
FSIF without die
10 mm

65
60
55
50
45
40
35
30
25
20
15
Tool feed rate10
v = 4000 mm/min
50
0 1 2 3 4 5 6 7 8 9 10111213
Tool rotation rate,  [1000 rpm]

Location

o

65
60
55
50
45
40
35
30
25
20
15
Tool feed rate10
v =1000 mm/min
50
0 1 2 3 4 5 6 7 8 9 10111213
Tool rotation rate,  [ 1000 rpm](a) v = 1000 mm/min(c) Curled(b) Burred
Figure 3 Appearances of formed sheets

Wall angle,  [ ]oWall angle,  [ ](a) Succeeded

10 mm
Location

Figure 5 Comparison of formed
shapes with and without die

4. Conclusions
Aluminum alloy sheets were formed by friction stir incremental forming with back support die. There
were some differences between with and without die as follows.
1. In some working conditions, burr or curling up were observed although those were not observed
without the back support die.
2. As the tool feed rate was increased, the minimum formable wall angle was increased with die.
3. Cross-sectional shapes at the conical surface were almost same, however, those at the corner andflange were different and that with die was closer to the ideal one.
Reference[1] Otsu, M., Ichikawa, T., Matsuda, M., Takashima, K., Friction Stir Incremental Forming of Aluminum Alloy
Sheets, Steel Research International, 81- 9(2010), pp. 942 - 945.
[2] Ambrogio, G., Galiardi, F., Bruschi, S., ,Filice, L., On the high-speed Single Point Incremental Forming oftitanium alloys, CIRP Annal-Manufacuring Technology 62(2013), pp. 243-246.

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DEVELOPMENT OF CU-SN ALLOY WITH MECHANICAL ALLOYING
Shunsuke WATANABE1*, Mituo NOTOMI2
Student, Department of Mechanical Engineering, Meiji University, 1-1-1 Higashimita Tama-ku, Kawasaki-shi,
Kanagawa-ken, 214-8571, Japan
2
Professor, Ditto*Corresponding Author: ce32068@meiji.ac.jp
Keywords: Mechanical alloying, Cu-Sn alloy, Powder molding, Ductility, XRD
1. Introduction
Cu-Sn alloys have been used through ages because of good workability, high corrosion
resistance and high electric conductivity in Cu-Sn solid solution. Recent research have revealed thesecondary transformation and new phase for the alloys, which made it as an interesting material tosubstitute the graphite anodes of Lithium ion batteries, lead-free solder as well as shape memory alloy(SMA)(1). The Cu-Sn alloys are potential to be used in various applications because it shows goodductility under each solid solution limit. Over each solid solution limit, brittleness appears due tointer-metallic compound form such as Cu3Sn, and Cu6Sn5 which made it difficult to process it using
plastic forming method that is commonly used as practical method. Although the traditional Cu-Sn
alloy is practically rare usage due to the occurrence of cleavage fracture under polycrystalline, theductile Cu-Sn alloy had a potential for SMA. SMAs have drawn significant attention due to theiruniqueness, i.e; they return original shape induced by heating or electromagnetic field. Cu-based SMA
is a promising substitution for the Ni-Ti alloy, which is the most common SMA due to superior shapememory effect and more expensive than Cu-Sn alloy.
In this research, Cu-Sn alloy in a mixing structure of Cu-FCC solid solution and Cu3Sn is
prepared to enhance ductility for the potential of shape memory effect. The atomic ratio is fixed inreference to the new phase diagram(1). Also, Cu-Sn (10at.%) alloy powder were synthesized bymechanical alloying (MA) from mixtures of pure crystalline Cu and Sn powders using planetary ballmill with the weight ratio of ball to powder. The phase structure of the alloy powders were evaluatedwith the X-ray diffraction (XRD) and the texture and hardness of specimens were investigated with thescanning electron microscope (SEM).
2. Experimental method
In this study, elemental powder of Cu with >99.8% of purity and < 100 μm of particle size (Niraco
Corp.) and Sn with > 99.8% of purity and < 45μm ofparticle size (Sigma Aldrich Corp.) were mixed inplanetary mills(LP-4, Ito Ltd.), to prepare Cu-10Sn(at. %). The ball-milling parameters are shown in Table 1.
Pre-alloyed powders were hot-pressed on a steel
die to rectangular shape of 12 × 12 × 5 mm withcompression testing machine (AYC-300P, Tokyo Testing
Machine Inc.) and done heat treatment under argon
atmosphere (FUA112DB, Advantec Toyo Kaisha Ltd.).
Heat treatment conditions are shown in Table 2. Phase
analyses were carried out on XRD (Rigaku Analix
A-26L-Cu). The microstructures of the specimen were
examined using a SEM (JEOL JSM-6700F) equippedwith energy-dispersive X-ray spectroscopy (EDS).

Table 1 The mechanical alloying parameter
Milling material
Ball-to-powder ratio
Filling rate [%]Rotation speed [rpm]Milling time [h]Milling atmosphere

Steel balls
30:1
10
200
40
Argon

Table 2 Heat treatment condition
Hot press
Pressure[MPa]

Principal
Powder80-

Heat treatment

Holding

Holding

Heat

Holding

temp
[℃]

temp
[sec]

treatment[temp.]-

time [hr]

-

-

(a)

200

5

(b)

200

-

500

-

5

(c)

-

-

(d)

200

-

80

200

300

specimen

500

(e)

5

500

Milledpowder

Cooling

method--200

Furnacecooling

(f)
(g)
(h)

5
500
80

500

300

(i)
-

200

(j)
(k)

5
500

(l)

3. Results and discussions
3.1 Crystal structure for Cu-Sn The structural evolution during MA, hot press and heat treatment
in order of Cu and Sn is shown in Fig.1. Cu and SnO appeared in the profile of the specimen made ofPage 87

The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

principal powder as shown Fig. 1(a). The figure indicated that Sn was promoting oxidation reaction
under high temperature. Also, Fig. 1(b) showed the peak of Cu, Cu-Sn solid solution and ε phase andthe phases were caused by the melting of Sn and react Cu by heat treatment. Okabe et al.(3) reported
that Cu-10mass%Sn compact forms ε phase during sintering at 500℃. The ε phase precipitated by theheat treatment at 500℃.
Fig. 1 (2) shows the profile of the pre-alloyed powder after MA treatment. The heights of all
diffraction peaks of Specimen (c) decrease while the width of the peaks increases. The lower
diffraction peaks are caused by the decrease of crystallite size and the increase of micro strain due tohigh stresses evolving during ball milling impact. The specimen (d) that has been recrystallized byheat treatment confirmed the peak of Cu-Sn solid solution and ε phase, and the position of Cu peakslightly shifted to smaller angles because of the diffusion of the Sn in the Cu matrix. Lattice parameterof FCC Cu-Sn solid solution consequently varied from 0.362 nm to 0.368 nm. We confirmed
Specimen (d) dissolved 7.1 at.% of Sn in Cu matrix from the results of N. Saunders et al.(4) reported
that the lattice parameter tends to increase by substituting or intruding Sn atoms in Cu matrix because
atomic size of Sn is bigger than Cu.
Fig. 1 (3) shows the specimen of the heat treatment at 200℃ with pre-alloyed powder by
MA treatment. Specimen (e) shows only the peaks of Cu-Sn solid solution. However, Specimen (f) has
the phase of Cu-Sn solid solution of which the peak is positioned bigger angle compared to Specimen(e) and ε phase. This is because the atom of Sn in Cu matrix is precipitated as ε phase and the lattice of
Cu-Sn solid solution is shrunk with the smaller atoms of Sn in Cu matrix.
Fig. 1 (4) shows the specimen at 500℃ with pre-alloyed powder by MA treatment. Both
Specimen (g) and (h) show the same diffraction patterns which appear only the peaks of Cu-Sn solid
solution.
3.2 Surface morphology for Cu-Sn Fig. 2 shows the images with SEM and EDS mapping ofspecimens. Specimen (a) and (e) can be seen porosity which leads to brittle fracture in Fig. 2 (1). On
the other hand, specimen (g) is confirmed to promote densification by applying high temperaturesduring sintering.
From EDS mapping images of Cu and Sn in Fig. 2, segregation of Sn can be seen on
specimen (a) while the specimen (e) and (g) shows that the distribution of the chemical composition ishomogeneous of Cu and Sn due to be pre-alloyed by MA.
Cu
Cu solid solution
Cu3 Sn (ε phase)
SnO(d) Heat treatment at 500℃

Intensity (a.u.)

Intensity (a.u.)(b) Heat treatment at 500℃(a) No heat treatment(a)(c) No heat treatment

(g)

(e)

100 μm

100 μm

100 μm(1) SEM micrograph

20

30

30

40

40

50

50

60

60

70

70

80

80

(1) Principal powder hot-pressed at 200℃

20
20

30
30

40
40

50
50

60
60

(2) Pre-alloyed powder(f) Heat treatment at 500℃

70
70

(g)

(e)

(a)
20

80
80

(h) Heat treatment at 500℃
100 μm

100 μm

100 μmIntensity (a.u.)

Intensity (a.u.)(2) Elemental mapping of Cu(e) No heat treatment

(e)

(a)

100 μm

20
20

30
30

40
40

50
50

60
60

70
70

80
80

(3) Pre-alloyed powder hot-pressed at 200℃

20
20

(g)

(g) No heat treatment

30
30

40
40

50
50

60
60

70
70

80
80

(4) Pre-alloyed powder hot-pressed at 500℃

100 μm

100 μm(3) Elemental mapping of Sn

Fig. Fig.
1 1XRD
specimens
Fig. Fig.
2 2Surface
morphologyof specimens
XRDpatterns
patterns of of
specimens
Surface morphologyof specimens
Reference[1] S.Furtauer, D.Li, D.Cupid, H.Flandorfer, The Cu-Sn phase diagram, Part 1: New experimental results,
Intermet al.lics, Voi.34, 2013, 142-147[2] T.Okabe, Y.Kado, Some Aspects on the Sintering of Cu-Sn Compact, J.Jpan.SOC.Powder Powder Met
al.lurgy, 14, 1967, 147-153[3] N.Saunders, AP Miodownik, The Cu-Sn (Copper-Tin) System, Bulletin of Alloy phase diagram, Vol.11,
Bo.3, 1990, 278-287

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

FEM SIMULATION FOR TWO-DIMENSION CUTTING OF TITANIUM-ALLOY
Makoto Nikawa1,*, Hiroki Mori2, Yuki Kitagawa2, Masato Okada3
1
2

Associate Professor, Gifu University, 1-1 Yanagido, Gifu, Japan, 501-1193

Graduate student, Graduate School of Engineering, Gifu University, 1-1 Yanagido, Gifu, Japan, 501-1193
3
Lecturer, University of Fukui, 3-9-1 Bunkyo, Fukui, Japan, 910-8507*Corresponding Author: mniikawa@gifu-u.ac.jp, Tel:+81-58-293-2517, Fax:+81-58-293-2517

Keywords: Cutting, FEM, Titanium alloy, Johnson-Cook model, Ductile fracture condition
Introduction
FEM simulation is a strong means to the creation of a new processing method and the optimization of
the existing processing method. There are various reports about the simulation of the cutting [1]. The
simulation was actively utilized in the field of plastic working [2] such as the forging and the sheet
metal forming, it had become indispensable tools for manufacturing. However, the simulation was not
utilized enough in the field of the cutting. The reason, it is difficult to actually measure the phenomenaduring cutting due to the severe environments of high temperature, high pressure, and high strain rate.
Therefore, the constitutive equation of materials was not determined.
FEM simulation for the two-dimension cutting of Ti-6Al-4V alloy was investigated in this study. The
cutting force, tool flank temperature, and dimension of the chip after cutting were measured by cutting
test and the parameters for the simulation were determined.
Testing methods
Cutting test
Titanium alloy (Ti-6Al-4V) of t3 x 100 mm was used as workpiece. Cutting tests of the workpiece
were carried out with an NC milling machine by dry cutting. Cutting was performed so that adown-cut cutting style was performed. The test conditions were shown in Table 1. The cutting force
was measured with a three-component dynamometer. The cutting force in this study was obtained the
maximum cutting pulse. The tool flank temperature during cutting was measured using a two-color
pyrometer with an optical fiber [3].
Table 1 Experimental conditions
Cutting speed V (m/min)25, 50, 100, 150, 200, 300
Rake angle α (deg.)-45, -25, -9, 5Radial depth of cut t1 (mm)0.1, 0.25, 0.3, 0.5

Computer simulation
Rigid-plastic analysis was carried out using DEFORM-2D Ver.9.0 which was commercial software,
and Johnson-Cook’s model was used for a flow stress equation of material. The equation is expressedby eq. 1.
1

ln

1

1

where A; initial yield strength, B; strain-hardening coefficient, C; strain-rate sensitivity, n;．
．
strain-hardening exponent, m; thermal-softening exponent, ε; equivalent strain rate, ε0; normalizedwith a reference strain rate, Tm; melting temperature of the material, Tr; room temperature.
The parameters in this model were used a reported value by Meyer-Kleponis [4]. The friction between
the chip and cutting tool was assumed the Coulomb friction. The friction coefficient that applied tosimulation was calculated by the results of the cutting test.
Failure accumulation in the Johnson-Cook model was considered by applying the normalized
Cockcroft & Latham law to the failure condition of materials. The equation is exposed by eq. 2.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

D

σ
σ

dε

2

where Df; fracture limit value, σmax; maximum principal stress, σ; equivalent stress, dε; equivalentstrain increment.
The validity of the calculation results was considered by comparing with the experimental results thatwere the cutting force, tool temperature, and chip shape.
Results
Figure 1 shows the effect of the ductile fracture condition. When a ductile fracture condition was notconsidered, the chip shape was calculated the flow type (Figure1 (a)), the characteristic saw-tooth typein titanium alloy was not calculated. However, when a ductile fracture condition was considered, thechip shape was calculated the saw-tooth type (Figure 1 (b)). The chip dimensions were varied
according to a limit value of the ductile fracture law.
(Limit value = 0.07)
Chip
Cutting tool

Workpiece(a) Non-consideration of the fracture condition(b) Consideration of the fracture condition

Figure 1 Effect of the ductile fracture condition to the calculation result of chip shape

Figure 2 shows the relation between cutting force, thrust force and the limit value of ductile fracture.
The cutting force was increased with increasing the limit value, the limited value was 0.1 or more, thecutting force was approximately constant value. The calculation result of cutting force wasapproximately 80 N lower than actual value. The thrust force was approximately accorded withexperiment.

Figure 2 Effect of the ductile fracture condition to the calculation result of chip

These results were suggested that the cutting simulation of the titanium alloy was possible by the by
the Johnson-Cook model that applied a ductile fracture condition.
References[1] P. Sartkulvanich, F. Koppka, T. Altan, Determination of flow stress for metal cutting simulation - a progress
report, J. Mater. Process. Technol., 146(2004), pp.61-71.
[2] R. Neugebauer, K.-D. Bouzakis, B. Denkena, F. Klocke, A. Sterzing, A.E. Tekkaya, R. Wertheim, Velocity
effects in metal forming and machining processes, Annals of the CIRP, 60(2011), pp.627-650.
[3] M. Okada, A. Hosokawa, N. Asakawa, T. Ueda, End milling od stainless steel and titanium alloy in an oil
mist environment, Int. J. Manuf. Technol, 74(2014), pp. 1255-1266.
[4] T. Özel and M. Sima, Finite element simulation of high speed machining Ti-6Al-4V alloy using modified
material models, Trans. of NAMRI/SME, 38(2010), pp.49-56.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

CORROSION RESISTANCE OF Si-DOPED AMORPHOUS CARBON FILMS
AGAINST NITRIC ACID
Jumpei Nishikawa1, *, Masayuki Nakano2, Junko Hieda1,Naoto Ohtake1, Hiroki Akasaka1
2

1
Department of Mechanical Sciences and Engineering Tokyo Institute of Technology,
Department of Chemical Science and Engineering Tokyo National College of Technology*Corresponding Author: nishikawa.j.ab@m.titech.ac.jp

Keyword: Amorphous carbon film, Si-doping, Nitric acid solution, Corrosion resistance, Surface
plasmon resonance
1. Introduction
Hydrogenated amorphous carbon films (a-C:H) are one of the functional films which consists of sp2 andsp3 hybridized carbon and hydrogen atoms[1]. Typical properties of the films are high hardness and low
friction coefficient[2], and it is also known that they have high corrosion resistance against acids andalkalis such as sulfuric acid, hydrofluoric acid and ammonium hydroxide[3]. On the other hand, thesefilms are corroded by nitric acid because amorphous carbon films contain sp2 bonding same asgraphite[4,5]. The film properties strongly depend on the film structure and doping of other elements into
the films can also change the film properties. In this study, the effect of silicon (Si) doping on the
corrosion resistance of a-C:H films was investigated by the precise corrosion rate measurement method.
2. Experimental method
Si-doped amorphous carbon films (a-C:H:Si) was depositedon Si (100) substrate by a plasma assisted chemical vapor
deposition (PECVD) method. Acetylene (C2H2) and tetramethylsilane (Si(CH3)4) were used as raw material gases. To
control Si concentration of a-C:H:Si films, the flow rate ofSi(CH3)4 was changed while the flow rate of C2H2 wasconstant. The prepared films were soaked for 0-18 h in nitricacid solution adjusted to 2 M and flushed with distilledwater. The surface of the sample was evaluated by an X-ray
photoelectron

spectroscopy(XPS).

Furthermore,

theFig. 1. Schematic illustration of SPR
measurement apparatus

corrosion rate in nitric acid solution was estimated by using the phenomenon of the surface plasmon
resonance (SPR). A schematic illustration of the measurement apparatus is shown in Fig. 1. A 40 nm
thick silver (Ag) layer was deposited on a glass substrate by magnetron sputtering method. a-C:H:Sifilms were deposited on Ag layer. The laser light incident from the back side was totally reflected on Ag
layer to generate evanescent wave and excite plasmon waves on the boundary between Ag layer and aC:H:Si film. Since these SPR conditions are determined by the dielectric constant of the space above
the Ag layer, it is able to estimate the reduction in the thickness of the film due to corrosion as a changeof the dielectric constant. The thickness was estimated in 0.3 M nitric acid solution by SPR measurement.
3. Results and Discussion
From the XPS spectrum, the proportion of sp2-bonding carbon in the film reduced by doping Si. The

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structure of each film did not change significantly by etching in nitric acid solution.
Figure 2 shows SPR profiles at each duration after the introduction of nitric acid solution. The intensityof the reflected light was markedly attenuated at an incident angle of about 57◦, which is the SPR angle.
Although the total reflection angle was not shifted in these profiles, the SPR angle was shifted by anamount that increased with the duration. From these SPR angles, amount of thickness change in film
thickness was calculated by fitting of SPR profiles. The results of the corrosion rate of the Si doped andnon-doped a-C:H films in the nitric acid solution estimated by SPR measurement is shown in Fig. 3.
Corrosion rate in nitric acid solution was reduced by the doping of Si into amorphous carbon films.

(a) Si-doped a-C:H(b) non-doped a-C:H

Fig. 3. SPR profiles in nitric acid solution.

Fig. 2. Change in thickness by corrosion.

4. Conclusions
High-resistance portion and low-resistance portion against nitric acid exist in the a-C:H films and no
significant change of the structure of a-C:H by oxidation occurred during etching in nitric acid solution.
By doping Si into a-C:H films, sp2-bonds ratio decreased and corrosion rate was reduced. These results
indicate that doping Si into a-C:H film increases its corrosion resistance against nitric acid.
5. Acknowledgement
This work was supported by Mitutoyo Association for Science and Technology.
References[1] J. Robertson, Comparison of diamond-like carbon to diamond for applications, Phys. Status Solidi., 205, 9(2008) pp233-2244.
[2] Y. Liu,A. Erdermir, E. I. Meletis, A study of the wear mechanism of diamond-like carbon films, Surf. Coat.
Tech., 82 (1996) pp48-56.
[3] R. D. Mansano, M. Massi, A. P. Mousinho, L. S. Zambom, L.G. Neto, Protective carbon layer for chemicalcorrosion of stainless steel, Dia. Relat. Mater., 12 (2003) pp749-752.
[4] K. Kamegawa, K. Nishikubo, H. Yoshida, Oxidative degradation of carbon blacks with nitric acid (I)—
Changes in pore and crystallographic structures, Carbon 36, 4 (1998) pp433-441.
[5] A. Takarada, T. Suzuki, K. Kanda, M. Niibe, M. Nakano, N. Ohtake, H. Akasaka, Structural dependence ofcorrosion resistance of amorphous carbon films against nitric acid, Dia. Relat. Mater., 51 (2015) pp49-54.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

TENSILE AND FATIGUE BEHAVIORS FOR LIP–SEAL WELDS OF LIQUID
LITHIUM PIPES AT ELEVATED TEMPERATURE
Kazuyuki Furuya1*, Fumio Takeo1, Hiroshi Yokoyama1, Tomoharu Akagaki1
1

Department of mechanical engineering, Hachinohe National College of Technology*Corresponding.kazuyuki–m@hachinohe–ct.ac.jp

Keywords: Li–Flow, Lip–Seal welding, Metallurgical property, Tensile property, Fatigue property
Introduction
An intense neutron irradiation system like the International Fusion Materials Irradiation
Facility is an accelerator–based D–Li neutron source designed to produce an intense high–energy
neutron flux (2 MW/m2) for testing candidate materials and components to be used in the International
Thermonuclear Experimental Reactor and the fusion DEMO reactor. Following a three–year key
element technology phase up to the end of 2002 [1], a transition phase is being performed prior toengineering validation and engineering design activities [2]. To realize such a condition, two 40–MeV
deuteron beams with a total current of 250 mA are injected into a liquid Li stream at 523 K, flowing at
20 m/s. Intense neutrons are emitted inside the Li flowing on a thin back–plate attached to the target
assembly. The target assembly operates under a severe neutron irradiation condition (50 dpa/year forthe back–plate). Therefore, the target assembly must be designed as a removable component to be
replaced from a pipe in a loop system in which the Li flows, using a remote handling system. In thecurrent design, the target assembly and the Li pipe are designed to have lip parts, each of which are
made of an austenitic stainless steel type–316L (hereinafter known as 316L steel). Several remote cutsand re–welding between the lip parts are anticipated to be done using a fiber laser which is the mostpromising tool in the present design. In case of welding, the welding method is called for a lip–seal
welding. The lip–seal weld operate under a corrosion/erosion condition due to the high speed Li–flowat the elevated temperature described above although the neutron irradiation damage at the weld does
not exceed about 1 dpa. For that reason, it is important to elucidate metallurgical conditions andmechanical properties in the weld to select optimum welding conditions. This extended abstract
presents microstructure, hardness and high temperature tensile and fatigue properties of the weld before
neutron irradiation.
Experimental Methodology
Specimens of two types were prepared using a 316L steel. One specimen, which is for the
metallurgical and hardness tests, simulated for lip–seal welding (hereinafter known as LSW). The
other butt–joint, which is for the mechanical tests, simulated for thickness of the lip–seal part (2–mm–
thick). The LSW and butt–joint were fabricated using a fiber laser tool, in the welding speed of 4000
mm / minutes and the supplied power of 1000 ~ 1100 W. Argon gas was utilized as covering gas
because there is the possibility that nitrogen gas which is used normally for covering gas, reacts with Li
inventory in Li pipe. As for the LSW, two 316L boards (2–mm thick each) overlapped were fixed byclamping tool, followed by lip–seal welding. Test pieces sampled from the weld were surface–finished.
Some of them were followed by electrolytic etching for observation of the microstructure using an
optical microscopy. Hardness distribution at the welds was measured using a Vickers hardness tester,with the testing load of 300 g and dwell time of 10 seconds. As for the butt–joint, which is specimens
for tensile and fatigue tests, is 24–mm long, 16–mm wide, and 2–mm thick: the cross–section is 32 mm2.
A weld is located at the center of the specimen. The specimens were surface–finished for all thesurfaces. Tensile and fatigue specimens of the same type were also sampled from the as–received 316L
steel for reference, prepared in the same way. Tensile and fatigue tests were carried out at elevatedtemperatures according to relevant ISO specifications.
Result and Descussion
Figure 1 shows macroscopic and microscopic views of the LSW. As shown in [I], soundwelding is visible in the weld with no harmful defects such as overlap, undercut, or macro–cracks

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100μm
B
316L
board

316L
board

0.3
0.5
0.7
190-200 0.9
180-190 1.1
170-1801.3
160-1701.5
150-160
1.7
1.9-1.6 -1.2 -0.8 -0.4 0 0.4 0.8 1.2 1.6
Distance from Center of Weld (mm)Figure 2 Micro–Vickers Hardness
Distribution on Weld.

Gap
Expulsion

Figure 1 Macrostructure and Microstructureof LSW. [I]: Cross–sectional macrostructureof the weld. A and B: Magnified views of bondregion and fusion metal

References[1] IFMIF International Team (Ed.), “IFMIF–KEP
International Fusion Materials Irradiation Facility
Key Element Technology Phase Report”, JAERI–
Tech 2003–005 (2003) pp. 5–16.
[2] H. Takatsu, M. Sugimoto, S. Jitsukawa, H. Matsui,
Proceedings of the 12th International Conference
on Fusion Reactor Materials, J. Nucl. Mater. 329–
333 (2007) pp. 178–182.

1000

Strength (MPa)

100μm

YS (BM)
UTS (BM)
UE (BM)TE (BM)

800

100

YS (WJ)UTS (WJ)
UE (WJ)TE (WJ)

80
Elongation (%)

Porosities

A

Approx. 3.0 mm[I]

Distance from Surface of Weld (mm)although there were a couple of porosities in these welded metals. Typical examples of thesemicroscopic views are depicted in panels A and B. As revealed in those views, no micro–crack isapparent at around the bond. As for heat affected zone (HAZ), no marked change occurs in the grain
size. As can be seen in the B of the figure, the LSW has small gap about 15–μm wide, and it can beseen a material like an inclusion between the gap. In result of EDX analysis for the material, BM andfusion metal, it was revealed that the material was expulsion came from fusion metal. Micro–Vickers
hardness distribution of the LSW is presented in Figure 2. The minimum and maximum hardness are
150 HV in the base metal region and 200 HV in the fusion metal region, and the increment of hardness
is 50 HV. Result of high temperature tensile test are portrayed in Figure 3 for butt–joint and referencematerial. Almost all the fracture occurred not at the welded region but at a base metal region. As canbe seen in the figure, it can be said that property change of the butt–joint are almost the same trend at
all temperature range. The 0.2 % off–set yield stress was decreased up to 140 MPa in accordance withtemperature increases. The other properties indicated the same trend, i.e., the ultimate tensile stress,
uniform elongation and total elongation were also decreased up to about 400 MPa, 7.5 % and 44 %,respectively. In result of fatigue test at RT for a reference material, the fatigue limit was 370 MPa.
On the other hand for the butt–joint tested at 623 K, the fracture occurred not at the welded region but
at a base metal region, and the fatigue limit was 280 MPa. Therefore it can be said that the value (280MPa) is comparable to a fatigue limit of the reference material at 623 K. The fatigue limit of the butt–
joint tested at 623 K decreased by 24 % compared to the value of the reference materials tested at RT(370 MPa). Consequently, it can be said that strength of the butt–joint are nearly equal to the referencematerials, and it is thought to be that the welding condition are relevant in the range of the mechanicaltest condition.

600

60

400

40

200

20
0

0
0

100 200 300 400
Testing Temperature (ºC)

500

Figure 3 Comparison of Tensile Propertiy ofReference Material and Butt–Joint.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

EFFECT OF LOCAL PLASTIC PROCESSING ON FATIGUE STRENGTHOF STEEL PLATE WITH A HOLE
Fumio Takeo1*, Kazuyuki Furuya1, Michitoshi Toyota1, Hiroki Naomachi1,
Daiki Yamada1, Syogo Sasaki1, Tomoharu Akagaki1, Hitoshi Soyama2
1

Department of Mechanical Engineering, Hachinohe National College of Technology,2
Department of Nanomechanics, Tohoku University* takeo-m@hachinohe-ct.ac.jp

Keywords: Fatigue Strength, Plastic Processing, Steel Plate, Circular Hole, S-N Curves
INTRODUCTION
Fatigue fracture of metallic materials has caused many serious accidents.
Fatigue is easy to occur at large stress. Therefore, it is effective to improve mechanical properties instress concentrated area of mechanical parts for enhancing the fatigue strength of them. In the previous
studies, for example, the effects of local plastic deformation [1] and cavitation peening [2] on the bending
fatigue strength of the metallic plate with a hole have been reported. In this study, the effect of the local
plastic processing on tensile fatigue strength of steel plate with a circular hole which is often used forfastener was examined. Tensile fatigue tests were carried out with two types of loading methods. Basedon the S-N curves, the effects of two types of local plastic processing on fatigue strength were comparedand discussed.

50

50

EXPERIMENTAL APPARATUS AND PROCEDURES
The fatigue strength of steel plates
having a circular hole was examined. As shown in Figure 1(a) and (b), the cyclic tensile load was applied
to the specimen by two methods; we call them the uniform loading type and the pin loading typerespectively. In the uniform loading type, tensile load was applied by gripping both ends of the specimen.
In the pin loading type, tensile load was applied to the pin which was inserted into the specimen’s hole,while the other end of the specimen was gripped. Thispin was made of quenched alloy steel and was 19.5mm
in diameter. The material used for the specimen was
cold-reduced carbon steel sheets of JIS G 3141 SPCCSD. Figure 2 shows the shape and size of the specimen
for two loading types. The specimens were 300 mm or
200 mm in length, 50 mm in width and 3.2 mm inthickness with a hole of 20 mm in diameter.
(a) Uniform loading(b) Pin loading
The plastic processing was introduced by pressing
Figure1Typesof loading
the edge of the hole using two types of punches asshown in Figure 3. The flat type punch was 30 mm indiameter, so the region within 5 mm from the hole-edge
20
Gripped area
Gripped area
was deformed. Compressive load of 180 kN was
applied with the punch for 30 seconds, then the
thickness of processed area was reduced by about
t 3.2
5.4 %. On the other hand, the load of 60 kN was applied
45
45
300
with the cone type punch, and the region within about(a) For uniform loading
1 mm from the hole-edge was deformed like thechamfered shape as shown in Figure 3(b).
20Gripped area
Tensile fatigue tests were carried out for all
combinations of plastic processing types and loading
types. The stress ratio of the minimum tensile stresst 3.2
min, to the maximum tensile stress max, was 0.1. In this45
50
Unit: mm200
paper, max and min were calculated from the applied
load and the cross sectional area perpendicular to the(b) For pin loading
loading direction through the shorter dimension of the
Figure 2 Geometry of specimen
specimen that passes through the center of the hole.

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The 4th Asian Symposium on Material Processing (ASMP), Lombok 10-13 August 2015

Punch

Load P

Load P

Punch

30

45°
3.2

3.2

20

20

(a) Flat type(b) Cone type

Unit: mmFigure 3 Schematic diagram of plastic processing

Frequency of stress cycle was 30 Hz. When the specimen was not fractured until 107 cycles, the test was
discontinued.

Maximum stress σmax [MPa]Figure 4
EXPERIMENTAL RESULTS
300
shows S-N curves obtained by fatigue tests.
Uni., Unprocessed
Blank and solid plots show the results forUni., Cone type280
uniform loading type and pin loading type,
Uni., Flat typePin, Unproccessed260
respectively. Arrows in Figure 4 mean that
Pin, Cone type7
the specimen does not fracture at 10 cycles.
Pin, Flat type
240
Compared with two loading types about
unprocessed specimen, the fatigue limit for220
pin loading type was about 6 % lower than200
that for uniform loading type. As to theeffects of local plastic processing on the180
enhancing of fatigue limit, it can be found160
that fatigue limit of processed specimen for
the uniform loading type was improved by140
about 20 % than that of unprocessed120
specimen. Considering with the magnitude ofrequired load for plastic processing, the cone
100
type processing seems to be more effective
1.E+04
1.E+05
1.E+06
1.E+07
1.E+08
104
105
106
107
108
than the flat type. On the other hand, theNumber of cycles N
effect of plastic processing for the pin loading
Figure 4 S-N curves obtained from fatigue tests
type was lower than that for the uniform
loading type. Improvement of fatigue limit
was about 10 % for the flat type processing. Especially in the case of cone type processing, fatigue limit
was slightly lower than unprocessed specimen. As one of the reason for this result, the friction between
the pin and the inner wall of the hole could be considered. In the case of the pin loading type, fretting
might contribute to the fatigue strength reduction. For the cone type processed specimen, fatigue strength
would be more affected by fretting because of its small contacting area.
CONCLUSIONS The effect of local plastic processing on fatigue strength of a steel strip with a holewas examined. It was found that fatigue limit could be improved by about 20 % by plastic processing
around the hole-edge for the uniform loading type. In contrast, the improvement effect for the pinloading type was low. In the case of pin loading type, fatigue strength might be affected by friction
between the pin and the inner wall of the hole.
References[1] Sheng-wu Wang, Shinich Nishida, Nobusuke Hattori and Norikazu Nakamura, Improvement in Fatigue
Strength of Notched Plate-Specimen by Local Plastic Deformation, Trans. JSME, 67(A)-664(2001), pp. 19962001.
[2] Hitoshi Soyama, The use of cavitation peening to increase the fatigue strength of duralumin plates containing
fastener holes, Mater. Sci. Appli., 5-6(2014), pp. 430-440.

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