Journal of Plasma and Fusion Research SERIES, Volume 10
Proceedings of the 4th Japan-China Workshop
on Fusion-Related Tritium Science and Technology
May 9–11, 2012, Toyama, Japan

Yukino Otani in
Tateyama Kurobe
Alpine Route

Tulip in Tonami

The Japan Society of Plasma Science and Nuclear Fusion ResearchContents
Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble
by Catalytic Metals...........................Chengjian XIAO, Chunmei KANG, Xingbi REN, Xiaoling GAO, Xiaojun CHEN,
Heyi WANG and Xiaolin WANG

1

Effects of Li4TiO4 Structure on Tritium Release Kinetics from Lithium-Enriched Li2+xTiO3......................................... Makoto KOBAYASHI, Kiyotaka KAWASAKI, Katsuyoshi TATENUMA,
Masanori HARA, Masao MATSUYAMA, Toshiyuki FUJII, Hajimu YAMANA,
Yasuhisa OYA and Kenji OKUNO

7

Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet............................................................................... Deqiong ZHU, Takuji ODA and Satoru TANAKA

12

Mass Loss of Li2TiO3 Pebbles in Atmosphere Containing Hydrogen....................................... Hideaki KASHIMURA, Masabumi NISHIKAWA, Kazunari KATAYAMA,
Shohei MATSUDA, Satoshi FUKADA, Tsuyoshi HOSHINO

18

Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University...................................................................................................................................Satoshi FUKADA

22

Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide...............................................................................Wei MAO, Takumi CHIKADA, Akihiro SUZUKI,
Takayuki TERAI and Kenji YAMAGUCHI

27

Solubility of Hydrogen Isotopes in Zirconia Ceramics........................... Kenichi HASHIZUME, Kosuke OGATA, Satoshi AKAMARU and Yuji HATANO

33

Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent...........................................................Kenzo MUNAKATA, Keisuke HARA, Takahiko SUGIYAMA,
Kenji KOTOH, Masahiro TANAKA and Tatsuhiko UDA

36

Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for
Methane Decomposition....................................................... Heyi WANG, Wenhua WANG, Yong YANG and Shuming PENG

42

Measurement of Tritium Distribution in Nickel and Vanadium by Means of a Combined
Technique of an Imaging Plate and Thin Absorbers.......................................................................................Hiroko YOSHIDA-OHUCHI, Yuji HATANO,
Akram MOHAMMADI and Takao KAWANO

49

III

The R&D Status of ITER SDS........................................... Sei-Hun YUN, Chang Shuk KIM, Min Ho CHANG, Hyun-Goo KANG,
Seungyon CHO, Hyun Gon LEE, Ki Jung JUNG, Hongsuk CHUNG,
Dae Seo KOO, Kyu-Min SONG, Soon-Hwan SHON and Duk Jin KIM

54

Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD............................... Masao MATSUYAMA, Shinsuke ABE, Kiyohiko NISHIMURA, Yasushi ONO,
Yasuhisa OYA, Kenji OKUNO, Tomoaki HINO and Akio SAGARA

64

Simulating Tritium Retention in Tungsten with a Multiple Trap Model in the TMAP Code.........................................Brad J. MERRILL, Masashi SHIMADA and Paul W. HUMRICKHOUSE

71

Implantation Energy Dependence on Deuterium Retention Behaviors for the
Carbon Implanted Tungsten................................................................. Yasuhisa OYA, Makoto KOBAYASHI, Naoaki YOSHIDA,
Naoko ASHIKAWA, Akio SAGARA, Yuji HATANO and Kenji OKUNO

76

Tritium Distribution on the Tungsten Surface Exposed to Deuterium Plasma and then to
Tritium Gas................................ Kanetsugu ISOBE, Vladimir Kh. ALIMOV, Akira TAGUCHI, Makiko SAITO,
Yuji TORIKAI, Yuji HATANO and Toshihiko YAMANISHI

81

Study on Plasma-Tungsten Surface Interactions Using the New Experimental Device EXPRESS...................................Ryo MIURA, Tetsuo FUJISHIMA, Yuto MIYAHARA, Tomohisa TAGUCHI,
Masahiro TANAKA, Akio SAGARA,Yasuhisa OYA and Kenji OKUNO

85

Hydrogen Incorporation into Tungsten Deposits Growing under Hydrogen and
Argon Mixed Plasma........................................................................................ Kazunari KATAYAMA, Yasuhito OHNISHI,
Satoshi FUKADA and Masabumi NISHIKAWA

89

Hydrogen Storage Properties of Nanocrystalline Mg2Ni Based Alloys Prepared by Ball-Milling.......................................................... Yifu XIONG, Jingwen BA, Wuwen QING and Wenyong JING

94

Hydrogen Generation Property on the Surface of Plasma Sintered Be12Ti..................................... Kohei WADA, Kenzo MUNAKATA, Yusuke AKIMOTO, Haruki TAKEDA,
Jae-Hwan KIM, Daisuke WAKAI, Kazuo YONEHARA and Masaru NAKAMICHI98

IV

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Preface
It is our great pleasure that the publication of the contributed papers to the 4th Japan-China Workshop(JCW) on Fusion-Related Tritium Science and Technology held in Toyama, Japan, during May 9-11, 2012,
has been completed here as the Journal of Plasma and Fusion Research Series published by the Japan Societyof Plasma Science and Nuclear Fusion Research.
This meeting is the 4th tritium workshop, formerly held at Mianyang in 2006, 2nd at Sendai in 2008,
and 3rd at Kunming in 2010. The Workshop on fusion-related tritium science and technology provides
the opportunity for discussion about tritium science and technology including tritium handling, tritium
confinement, tritium breeding for establishment of a thermonuclear fusion reactor and environment effects,
and other related areas for fission as well as fusion.
The present workshop was focused on the following four topics: (1) Fusion-related tritium science andtechnology, (2) Breeding blanket technology, (3) Environment effects, and (4) Other related tritium scienceand technology. The total number of the participants was 53 (30 from Japan, 19 from China, 2 from US
and 2 form Korea), and 35 invited and contributed papers were presented. Among these papers, 19 excellent
papers are finally published in this proceeding after reviewing by the publication committee.
I would be grateful to all the members of the Japanese and Chinese Organizing Committee and Local
Committee. Finally, I would express my sincere appreciation to the members of the Local Organizing Committee and the voluntarily supporting staff of the Hydrogen Isotope Research Center (HRC) of Universityof Toyama, who have let to the success of this workshop.

Organizing Committee
for the 4th JCW, Chair
Kenji Okuno

I

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Guest Editorial Board
Chief Guest Editor
Yuji Hatano (University of Toyama, Japan)

Editors
Zhu An (Sichuan University, China)
Masanori Hara (University of Toyama, Japan)
Kenichi Hashizume (Kyushu University, Japan)
Yukiharu Ohsawa (Nagoya University, Japan)
Yasuhisa Oya (Shizuoka University, Japan)
Yuji Torikai (University of Toyama, Japan)
Heyi Wang (China Academy of Engineering Physics, China)
Yuji Yamauchi (Hokkaido University, Japan)
Kozo Yamazaki (Nagoya University, Japan)
Hongguang Yang (China Institute of Atomic Energy, China)

IIJ. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Improvement
Improvement ofof Hydrogen
Hydrogen Isotope
Isotope Release
Release on
on Lithium
Lithium Orthosilicate
Orthosilicate Ceramic
Ceramic
PebblePebble by
by Catalytic
Catalytic Metals
MetalsabaaaChengjianRENa, Xiaoling GAO a, Xiaojun CHEN a,
Chengjian XIAOXIAO a,, Chunmei
Chunmei KANG
KANG b,, Xingbi
XingbiREN,
Xiaoling
GAO,
Xiaojun
CHEN
,
a
c, *
Heyi
WANG
a, Xiaolin WANG c, *Heyi WANG , Xiaolin WANG
aaInstitute of Nuclear Physics and Chemistry, CAEP, Box 919-214, Mianyang, Sichuan 621900, P. R. China

Institute of Nuclear Physics and Chemistry, CAEP, Box 919-214, Mianyang, Sichuan 621900, P. R. ChinabbScience and Technology Information Center, CAEP, Box 919-803, Mianyang, Sichuan 621900, P. R. China

Science and Technology Information Center, CAEP, Box 919-803, Mianyang, Sichuan 621900, P. R. ChinaccChina Academy of Engineering Physics, Box 919-214, Mianyang, Sichuan 621900, P. R. China
China Academy of Engineering Physics, Box 919-214, Mianyang, Sichuan 621900, P. R. China(Received: 9 May 2012 / Accepted: 26 October 2012)
Lithium Orthosilicate (Li4SiO4) ceramic pebble is considered as a candidate tritium breeding material of Chinese helium cooled
Lithium Orthosilicate (Li4SiO4) ceramic pebble is considered as a candidate tritium breeding material of Chinese helium cooledsolid breeder test blanket module㸦CH HCSB TBM㸧for the International Thermonuclear Experimental Reactor (ITER). In this paper,
solid breeder test blanket module㸦CH HCSB TBM㸧for the International Thermonuclear Experimental Reactor (ITER). In this paper,
Li4SiO4 ceramic pebbles deposited with catalytic metals, including Pt, Pd, Ru and Ir, were prepared by wet impregnation method. The
Li4SiO4 ceramic pebbles deposited with catalytic metals, including Pt, Pd, Ru and Ir, were prepared by wet impregnation method. The
metal particles on Li4SiO4 pebble exhibit a good promotion of hydrogen isotope exchange reactions in H2-D2O gas system, withmetal particles on Li4SiO4 pebble exhibit a good promotion of hydrogen isotope exchange reactions in H2-D2O gas system, withconversion equilibrium temperature reduction of 200í300 qC. The out-of-pile tritium release experiments were performed using 1.0
conversion equilibrium temperature reduction of 200í300 qC. The out-of-pile tritium release experiments were performed using 1.0
wt% Pt/Li4SiO4 and Li4SiO4 pebbles irradiated in a thermal neutron reactor. The results showed that Pt was effective to increase the
wt% Pt/Li4SiO4 and Li4SiO4 pebbles irradiated in a thermal neutron reactor. The results showed that Pt was effective to increase thetritium release rate at lower temperatures, and the ratio of tritium molecule (HT) to tritiated water (HTO) of 1.0 wt% Pt/Li4SiO4 wastritium release rate at lower temperatures, and the ratio of tritium molecule (HT) to tritiated water (HTO) of 1.0 wt% Pt/Li4SiO4 wasmuch more than that of Li4SiO4, which released mainly with HTO. Thus, catalytic metals deposited on Li4SiO4 pebble may help tomuch more than that of Li4SiO4, which released mainly with HTO. Thus, catalytic metals deposited on Li4SiO4 pebble may help toaccelerate the recovery of bred tritium particularly in low temperature region, and increase the tritium molecule form released from theaccelerate the recovery of bred tritium particularly in low temperature region, and increase the tritium molecule form released from thetritium breeding materials.
tritium breeding materials.

Keywords: Lithium Orthosilicate (Li4SiO4); Hydrogen Isotope Exchange Reaction; Catalytic Metal; Tritium
Keywords: Lithium Orthosilicate (Li4SiO4); Hydrogen Isotope Exchange Reaction; Catalytic Metal; Tritium

rate, it is difficult to reflect tritium diffusion in the bulk ofrate, it is difficult to reflect tritium diffusion in the bulk ofgrain. Then, platinum and palladium were tried to deposit
grain. Then, platinum and palladium were tried to depositon lithium-containing ceramics, and exhibited aon lithium-containing ceramics, and exhibited a
promotion of hydrogen isotope exchange reactions at low
promotion of hydrogen isotope exchange reactions at low
temperatures [6-9]. Considering the practical application
temperatures [6-9]. Considering the practical applicationof neutron activity materials in a fusion reactor, manyof neutron activity materials in a fusion reactor, many
problems need systematic study, such as comparison ofproblems need systematic study, such as comparison ofdifferent metals, the optimization of catalytic metal
different metals, the optimization of catalytic metal
content and catalyst poisoning.
content and catalyst poisoning.
Now, Li4SiO4 ceramic pebble is considered as aNow, Li4SiO4 ceramic pebble is considered as a
candidate tritium breeder of Chinese helium cooled solid
candidate tritium breeder of Chinese helium cooled solid
breeder test blanket module for the international
breeder test blanket module for the international
thermonuclear experimental reactor㸦ITER㸧[4]. In orderthermonuclear experimental reactor㸦ITER㸧[4]. In order
to investigate the energetic tritium moving mechanism in
to investigate the energetic tritium moving mechanism inthe crystal grain of Li4SiO4 ceramic pebble, study thethe crystal grain of Li4SiO4 ceramic pebble, study theisotope exchange reactions on the gas/solid interface of
isotope exchange reactions on the gas/solid interface ofthe pebble, and search some way to enhance the tritium
the pebble, and search some way to enhance the tritium
release at low temperaure, it is necessary to investigate
release at low temperaure, it is necessary to investigate
the behavior of hydrogen isotope exchange influenced by
the behavior of hydrogen isotope exchange influenced bycatalytic active metal. In this work, Li4SiO4 ceramic
catalytic active metal. In this work, Li4SiO4 ceramic
pebbles deposited with four kinds of catalytic metals, Pt,
pebbles deposited with four kinds of catalytic metals, Pt,
Pd, Ru and Ir, were prepared. The effects of different Pt
Pd, Ru and Ir, were prepared. The effects of different Ptcontent and different metal role on the hydrogen isotope
content and different metal role on the hydrogen isotope
exchange reactions were examined in H2/D2O gas system.
exchange reactions were examined in H2/D2O gas system.
The tritium release behavior of the neutron-irradiated 1.0
The tritium release behavior of the neutron-irradiated 1.0
wt% Pt/Li4SiO4 and Li4SiO4 was compared through the
wt% Pt/Li4SiO4 and Li4SiO4 was compared through theout-of-pile tritium release experiments.
out-of-pile tritium release experiments.

1. Introduction
1. Introduction

Lithium-containing ceramics㸪such as Li2O, LiAlO2,
Lithium-containing ceramics㸪such as Li2O, LiAlO2,
Li2ZrO3, Li2TiO3 and Li4SiO4㸪are considered as tritium
Li2ZrO3, Li2TiO3 and Li4SiO4㸪are considered as tritium
breeding materials for securing an adequate tritium
breeding materials for securing an adequate tritium
supply to sustain the D-T fusion reaction in a fusion
supply to sustain the D-T fusion reaction
in a fusion
reactor [1-4]. Tritium is produced via 66Li(n,D)T in thereactor [1-4]. Tritium is produced via Li(n,D)T in thelithium ceramic pebbles, and extracted by a helium sweep
lithium ceramic pebbles, and extracted by a helium sweepgas with a fraction of hydrogen or deuterium from thegas with a fraction of hydrogen or deuterium from thesurface of the pebbles. Considering the reduction for thesurface of the pebbles. Considering the reduction for the
residence time of tritium in the breeder blankets, the bred
residence time of tritium in the breeder blankets, the bred
tritium must be carried out with a fast recovery rate overtritium must be carried out with a fast recovery rate overa broad range of temperatures, so some means shall bea broad range of temperatures, so some means shall be
developed to accelerate the tritium release rate.
developed to accelerate the tritium release rate.
The mass transfer mechanism of tritium released from
The mass transfer mechanism of tritium released fromlithium-containing ceramics is complicated, but it islithium-containing ceramics is complicated, but it isdeveloped and assumed commonly that tritium is bred indeveloped and assumed commonly that tritium is bred ingrain interior, diffuses to the surface layer, and thengrain interior, diffuses to the surface layer, and thentransfer to the surface water layer which consists oftransfer to the surface water layer which consists ofphysical, chemical and structural adsorbed water. Finally,physical, chemical and structural adsorbed water. Finally,
tritium at the surface water discharges to the purge gas
tritium at the surface water discharges to the purge gas
through such surface reactions as desorption, isotope
through such surface reactions as desorption, isotope
exchange reactions between hydrogen and water vapor in
exchange reactions between hydrogen and water vapor inpurge gas [5]. However, the isotope exchange reactions atpurge gas [5]. However, the isotope exchange reactions at
the interface between solid surface and sweep gas are
the interface between solid surface and sweep gas arevery slow at low temperature, only fast at relatively
very slow at low temperature, only fast at relatively
elevated temperatures (>700 qC) [6, 7]. When isotope
elevated temperatures (>700 qC) [6, 7]. When isotope
exchange reactions contribute strongly to tritium releaseexchange reactions contribute strongly to tritium releaseauthor’s e-mail: xiaocj@hotmail.com
author’s e-mail: xiaocj@hotmail.com

1©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

C. Xiao et al., Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble by Catalytic Metals

2. Experimental

the experiments were conducted under steady state
conditions at ambient atmosphere pressure. The
experimental conditions and the expeimental flow chart
for hydrogen isotope exchange reactions are shown in
Table1 and Fig.1 respectively.

2.1 Preparation of Sample
Li4SiO4 ceramic pebbles were fabricated byfreeze-sintering process[10]. The preparation processincluded mixing Li4SiO4 powder, polyvinylalcohol
solution and antifoam agent, dropping the preparedmixture into liquid nitrogen through a nozzle to generate
spheres, freeze-drying the spheres in a freeze dry system,
and calcinating the spheres programedly to get ceramic
pebbles.
Li4SiO4 ceramic pebbles (av. 0.9 mm in diameter)
were selected to be deposited with noble metals by wet
impregnation method. The pebbles were dried in a
vacuum drier at 150qC for 3h. Then the metal-containing
solution, H2PtCl6, H2IrCl6, PdCl2 and RuCl3, was drippedon the pebbles respectively, and was adsorbed swiftly by
the pebbles. When the pebbles were saturated with thesolution, they were dried for 1-2 h by infrared ray. After
repeating the above steps until the pebbles had adsorbed apredetermined content, they were transferred to deoxidize
in a 5 vol% H2/N2 mixture gas at 400-450 oC for 3h.
Finally Li4SiO4 ceramic pebbles with catalyst wereobtained.

Fig.1 Experimental flow chart for hydrogen isotope
exchange reactions
2.4 Tritium Thermal Desorption
Li4SiO4 and 1.0 wt% Pt/Li4SiO4 ceramic pebbles,
which had been dried at 150 oC for 2 h in vacuum, weresealed in quartz capsules with He gas less than 10 Pa.
Then the encapsulated samples were irradiated in areactor with a neutron flux of 7.1×1012 cm-2s-1 for 100min
at ambient temperature (29-36 oC). Tritium was producedby thermal neutron capture of lithium via the reaction of
6
Li(n, Į)T. After irradiation, the samples were placed atroom temperature for two years. The quartz capsules
were broken in a glove bag filled with He gas. Somepebbles about 0.01-0.02 g were transferred into a reactor
tube made of quartz for tritium thermal desorption, andthe others were removed into a small test tube which wasstored in a desiccator for stand-by.

2.2 Physical Characterization
Powder X-ray diffraction (XRD) patterns of 1.0 wt%
Pt/Li4SiO4 were determined on a DX-2600SSC
diffractometer (Cu KD radiation at 40 kV and 30 mA) at ascanning speed of 1.2o miní1 from 15 o to 95 o. Pt particle
size was analyzed by Philips- CM30 transmission
electron microscope at an acceleration voltage of 200 kV.
Pt particles more than 100 were chosen randomly to test
the metal particle size by Photoshop CS2 software.
Table1 Experimental conditions

2.3 Hydrogen Isotope Exchange Reactions
The samples were dried at 150 oC for 2h in a vacuum
drier, and then installed in a settled reactor (I = 8 mm)
made of stainless steel. A mixture gas of 1.08 % H2/He
and 1.12 vol% D2O/He was introduced into the reactor.
The concentration of D2O was controlled by D2/He gas
passing through a CuO bed of 350 oC. The deuterium
concentration of HD in the outlet stream was measured
by gas chromatogram (Agilent 6890N). 601 column wasused as the column material. The HD peak could beobtained depending on the thermal conductivity
coefficient when hydrogen was used as a carrier gas. All

Fig.2 Experimental flow chart for tritium release
Tritium release experiments were performed byapplying the out-of-pile temperature programmed
desorption technique. The flow chart of the experimental
system is shown in Fig.2.The reactor tube was heated in a
tube furnace at a heating rate of 5 oC min-1 from room
temperature to 800 oC, while a purge gas of He containing
1vol% H2 carried the released tritium (HT and HTO)
through the experimental system at a flow rate of 50 ml
22

C. Xiao et al., Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble by Catalytic Metals

min-1. The first ionization chamber measured total tritium
released from the ceramic pebbles, the second ionizationchamber measured tritium released as HT. The amount oftritium in the bubblers, which were used to adsorb HTO(T2O) with glycol solution as adsorbent, was measured byliquid scintillation counting. Details of the experiments
were cited in the previous literature[11,12].

3.Results and Discussions
3.1 Preparation of Sample
As the data shown in Table 2, Li4SiO4 ceramic
pebble fabricated by freeze-sintering process has good
sphericity, high density and material phase purity .

Fig.5 TEM micrograph of 1.0 wt% Pt/Li4SiO4 pebble

Table 2 Characteristics of Li4SiO4 ceramic pebble
Distribution( %)

40
30
20
10
0

0

2

4

6
8 10 12
Particle size( nm)14

16

Fig.6 Distribution of Pt particles on the surface of 1.0 wt%
Pt/Li4SiO4 pebble

impregnation method. When the catalytic metals were
deposited on Li4SiO4, the surface color changed fromwhite to greyδ0.1 wt% Ptε, brown (0.5 wt% Pt) or blackδ1.0 wt% Ptε as shown in Fig.3. It is determined by
X-ray diffraction that Pt particles on the pebble surface
belong to face centered cubic structure with the
crystalline planes of Pt (1 1 1) and Pt (2 0 0) in Fig.4.
TEM analysis in Fig.5 shows that the agglomeration of Pt
particles is obvious, but the maximal size of Pt particles isless than 20 nm as shown in Fig.6. The mean size of Pt
particles is 6.7 nm.

Fig.3 Appearance of 1.0 wt% Pt/Li4SiO4 pebble

3.2 Isotope Exchange Reactions on pebble
The mixture gas of H2, D2O and He was feeded into
the reaction bed, D2O was adsorbed on the grain surfaceof the sample, and then the isotope exchange reactions
took place over the grain surface between the adsorbed
D2O and H2 in the gas. The exchange reactions on thebasis of the simplified model were expressed as
D2O(s) + H2(g) я HD (g) + HDO(s)HDO(s) + H2(g) я HD (g) + H2O(s)Fig.4 XRD patterns of Li4SiO4 and 1.0 wt% Pt/Li4SiO4

0.1í1.0 wt% Pt/Li4SiO4 and 0.1 wt% M/Li4SiO4 (M=Pd, Ru, Ir) ceramic pebbles were obtained by the wet

was
3
3

(1)
(2)

The deuterium concentration in the outlet stream
measured with gas chromatogram, and the

C. Xiao et al., Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble by Catalytic Metals

influence the enhancement of the isotope exchange
reactions. The conversion ratio of 0.1 wt% Pt/Li4SiO4reached equilibrium after 350 oC, and that of 0.5-1.0 wt%
Pt/Li4SiO4 reached around 300 oC. It indicates that more
content of Pt on the pebble surface would provide more
active adsorption localities, and the lower equilibrium
temperature for the isotope exchange reaction would beobtained. However, the efficiency of noble metals of thesurface should depend on the surface coverage for the
samples having different specific surface area. Y. Narisato
and K. Mochizuki reported that the better content rangeof Pd on lithium ceramic pebbles was 0.15-1.6 wt% or
0.02-0.2 wt% [5, 13].
Pd, Ru, and Ir were also deposited on Li4SiO4respectively, and exhibited to enhance the isotope
exchange reactions as shown in Fig.9. It is suggested thatmany catalytic metals could be used to promote hydrogenisotope exchange reactions on Li4SiO4 ceramic pebble.

conversion of D-atoms for the isotope exchange reaction
was defined as the formula (3) , where CD,out was theconcentration of HD in the outlet stream, and CD,in was
the concentration of D-atoms in the inlet stream.

RCDˈoutCDˈin(3)

Fig.7. Conversion ratio of hydrogen isotope exchange reactions

With respect to the virgin Li4SiO4, the rate of theisotope exchange reaction taking place on the surface wasquite slow, and the conversion ratio reached equilibrium
above 650 oC. However, the conversion ratio of 1.0 wt%
Pt/Li4SiO4 ceramic pebble increased rapidly between 200o
C and 300 oC, and reached equilibrium above 300 oC.
The experimental curves in Fig.7 indicate that theconversion equilibrium temperature decreased about 300o
C if 1.0 wt% Pt was deposited on the surface.
Catalytic metal deposited on Li4SiO4 ceramic pebbles
would provide active adsorption sites for H2 in the gas,and had a tendency to decompose H2(ads) into H(ads).
The H(ads) atom had higher reactivity and quicker
transfer capability, so the activation energy for theexchange reaction was reduced, and it’s easy to occur
hydrogen isotope exchange between the H(ads) atom andD2O molecule adsorbed on the surface of ceramic pebbles[6-9]. The hydrogen isotope exchange reaction process
may be expressed as follows.

Fig.8. Conversion ratio of isotope exchange reactions on
Pt/Li4SiO4 pebble

According to the thermodynamic functions,hydrogen isn’t adsorbed on catalytic active metal surface
in intact molecule, and it combines with metal atoms insingle atom [13, 14]. Water vapor molecule exists inintact molecule on Pt surface, but has a tendency to bedissociated to OH(ads) and H(ads) on Ir or Ru surface. So
there were different hydrogen isotope exchanging
processes on the surface of M/Li4SiO4 pebbles. It
indicates that if Pt or Pd with rational intermingles of Ru,
Ir or Fe, deposited on Li4SiO4 pebbles together, would
incease the catalytic activity. In that case, the amount of
the noble metal could be reduced in the tritium breeder
materials.

H2 (g) + 2M (s) я 2M-H (s)(5)
M-H(s) + D2O(ads) я M-H(D2O) я M-D(HDO) я
M-D(s) + HDO(ads)(6)
M-H(s) + M-D(s) я 2M(s) + HD(g)(7), where M(s) is the catalytic metal deposited on solid
surface.
To investigate the catalytic effect depended on
different metal content, the Li4SiO4 ceramic pebbles were
deposited with 0.1 wt%, 0.5 wt% and 1.0 wt% content of
Pt respectively. According to the experimental results
shown in Fig.8, the content of Pt on the surface could

3.3 Tritium Thermal Desorption
Fig.10 and 11 present the tritium release behavior ofneutron-irradiated Li4SiO4 and 1.0 wt% Pt/Li4SiO4pebbles under the 1.0 vol% H2/He purge gas flow. After
44

C. Xiao et al., Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble by Catalytic Metals

Fig.9. Fig.9.
Conversion
Conversion
ratio ofratio
isotopeof isotope
exchangeexchange
reactions onreactions
on4SiO
M/Li
M/Li
4SiO4 pebble
4 pebbleFig.10 Tritium release of neutron-irradiated Li4SiO4 pebble

the specimens were swept by pure He gas to wipe off thetritium adsorbed physically on the pebble, the purge gas
was changed to 1.0 vol% H2/He, and then the reactor tube
was heated at a rate of 5 oC /min. The blue curve
represents the total tritium concentration denoted as
HT&HTO released from the pebbles, and the red curve
represents the molecular form of tritium denoted as HT.
Fig. 10 shows the tritium release curves of virgin
Li4SiO4. There are two large peaks of total tritium from
R.T to 800 oC. One peak, appeared at around 259 oC, andthe second one, relatively larger than the former,
appeared at 418 oC. It’s evidently that tritiated water,
denoted as HTO, is much more than HT in the annealingcourse, especially at low temperature less than 400 oC. As
shown in Table 3, the fraction of HTO was 96.2 % of thetotal tritium, but HT was 3.8 %. It indicated that tritium
was released mainly as the chemical form of tritiatedwater in the experimental conditions of this paper.
Fig.11 shows the tritium release curves of 1.0 wt%
Pt/Li4SiO4. There are two peaks for both the total tritium
and the tritium molecular form during the annealingcourse. With regard to the total tritium curve, the larger
peak was observed at 218 oC, and the smaller peak
appeared at 468 oC. It means that most of tritium wasreleased at lower temperature less than 400 oC. The
tritium release curve of HT had a similar shape as that of
the total tritium. Table 3 shows that the fraction of HTO
was 36.9 %, but HT was increased to 63.1%. The fractionof chemical forms was quite different from that of thevirgin Li4SiO4. It indicates that the tritium molecularform (HT) became the main component when noble
metal deposited on the surface of Li4SiO4.
To compare the tritium release amount of Li4SiO4
and 1.0 wt% Pt/Li4SiO4, 37.4% of the total amount oftritium was released at 350 °C for Li4SiO4, but 66.9% for
1.0 wt% Pt/Li4SiO4. The largest peak of total tritium was
at 418 °C for Li4SiO4, but 218 °C for 1.0 wt% Pt/Li4SiO4.
So the catalytic effect of Pt was evidently. Even at
temperature less than 200 °C, the tritium concentration in
the outlet stream of the reactor tube increased rapidly

Fig.11 Tritium release of neutron-irradiated 1.0 wt% Pt/Li4SiO4pebbleTable 3 Fraction of tritium released from samples

from 1.0 wt% Pt/Li4SiO4. It means that when catalytic
metal was deposited on the surface of Li4SiO4, thehydrogen isotope exchange reactions were enhanced
between H2 in the purge gas and tritium adsorbed on thesolid surface.
With regard to the tritium chemical form, 3.8 % of
the total tritium was released as HT from room
temperature to 800 °C for Li4SiO4, and a small peak ofHT appeared at 463 °C. However, the fraction of HT was
63.1 % for 1.0 wt% Pt/Li4SiO4 with only one large peakat 236°C. The HT curve strongly influenced the shape ofthe total tritium curve for 1.0 wt% Pt/Li4SiO4. It indicatedwhen 1.0 wt% Pt was deposited on Li4SiO4, the hydrogenisotope exchange reactions, expressed as (5)-(7),wereeasy to proceed at the temperatures less than 460 °C.
The isotope exchange reactions on 1.0 wt%

5C. Xiao et al., Improvement of Hydrogen Isotope Release on Lithium Orthosilicate Ceramic Pebble by Catalytic Metals

Acknowledgements

Pt/Li4SiO4 proceeded much faster than that on virgin
Li4SiO4 , and the chemical forms of the released tritiumchanged. It was attributed to the catalytic effect of noble
metal on the solid surface. Without the effect of catalyst,
the exchange reaction rates of (1) and (2) were far smaller
than the desorption rate or the rate of isotope exchange
reaction (8), then the conversion of tritium (HTO intoHT) was quite slow.

H2O(g) + HTO(s) я HTO(g) + H2O(s)This work is financially supported by theDevelopment Fund of China Academy of Engineering
Physics (2010B0301035) and the National Magnetic
Confinement Fusion Science Program (2010GB112004).

References[1][2]

N. Rouxa, et al, Fusion Eng. Des. 41 (1998) 31-38.
C. E. Johnson, et al, J. Nucl. Mater. 258-263 (1998)
140-148.
[3] C. Xiao, et al, Progress in Chem. 23 (2011) 1906-1914.
[4] X. Chen,et al, Fusion and Plasma Phys. 26(2006)
210-216.
[5] T. Kinjyo, et al, Fusion Eng. Des. 81 (2006) 573–577[6] Y. Narisato, et al, J. Nucl. Mater. 329-333(2004)
1370-1373.
[7] K. Munakata, et al, J. Nucl. Mater. 307-311(2002)
1451-1455[8] K. Munakata, et al, J. Nucl. Mater. 386-388 (2009)1091-1094.
[9] K. Munakata, et al, Fusion Eng. Des. 58-59 (2001)
683-687.
[10] X Gao, et al, Atomic Energy Sci. and Tech. 44(2010)1099-1104[11] C. Kang, et al, J. Nucl. Mater. 412 (2011) 62–65.
[12] C. Xiao, et al. Fusion and Plasma Phys. 31(2011)
224-227.
[13] K. Mochizuki, et al, Fusion Eng. Des. 85(2010)
1185-1189.
[14] S. Hu, et al, International Jornal of Hydrogen Energy.
19(2010) 10118-10126

㸦8㸧

With the help of catalyst, it’s easy to occur hydrogenisotopes exchanging between H(ads) atoms and T
adsorbed on the surface of ceramic pebbles at lower
temperature as the formula(6). As soon as T exchanged
with H(ads) and migrated to the active adsorption locality,
it would discharge quickly combined with anotherH(ads).
Considering the high neutron activity in practicalapplication, the content of noble metals should be
reduced on lithium ceramic pebble, and some new metal
with low neutron activity need to discover.

Conclusions

Wet impregnation method has been used to fabricate
Li4SiO4 ceramic pebbles deposited with catalytic active
metals including Pt, Pd, Ru and Ir.
Lithium orthosilicate pebble deposited with noble
metals exhibite a promotion of hydrogen isotope
exchange reactions in H2/D2O gas system. With respect tovirgin Li4SiO4, the rate of the isotope exchange reaction
taking place on the surface was quite slow, and theconversion ratio reached equilibrium after 650 oC.
However, the conversion ratio of M/Li4SiO4 wasincreased obviously from 200 oC and 300 oC. Morecontent of the catalytic metal provided more active
adsorption localities on the pebble, so the conversion
ratio of 0.1 wt% M/Li4SiO4 reached equilibrium after 350
oC, and that of 0.5-1 wt% M/Li4SiO4 reached equilibrium
after 300 oC.
Most tritium was released at lower temperaturesfrom neutron-irradiated 0.1 wt% Pt/Li4SiO4, and thechemical form of HT was increased remarkably
compared with that of Li4SiO4. The result indicates thatcatalytic metal deposited on Li4SiO4 ceramic pebble
could reduce activation energy, facilitate H-T(D) isotope
exchange reaction, and influence the tritium releasemechanism. Consequently, the deposition of catalytic
metal on Li4SiO4 is an effective way to promote thetritium release from the breeder materials, and is infavor of the design of tritium cycle system about
extracting, separating and purifying in a D-T fusion
reactor.

6
6

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Effects of Li4TiO4 Structure on Tritium Release Kinetics from
Lithium-Enriched Li2+xTiO3
Makoto KOBAYASHI1, Kiyotaka KAWASAKI1, Katsuyoshi TATENUMA2, Masanori HARA3,
Masao MATSUYAMA3, Toshiyuki FUJII4,
Hajimu YAMANA4, Yasuhisa OYA1 and Kenji OKUNO1
1

Radioscience Research Laboratory, Faculty of Science, Shizuoka University, Japan
2
Kaken Co. Ltd., 1044, Hori, Mito-city, Ibaraki, 310-0903, Japan
3
Hydrogen Isotope Research Center, University of Toyama, Toyama, Japan
4
Research Reactor Institute: Kyoto University, Kumatori, Japan(Received: 10 May 2012 / Accepted: 3 December 2012)

Keywords : Tritium, Li2TiO3, Li4TiO4, diffusion, de-trapping, TDS
The release kinetics of tritium for Li2+xTiO3 (x = 0, 0.2, 0.4) and Li4TiO4 were studied by means of Thermal Desorption
Spectroscopy (TDS). Tritium-TDS spectrum at the heating rate of 0.5 K / min for Li2+xTiO3 irradiated with thermal
neutron consisted of two release stages at 480 K and 580 K, namely Peaks 1 and 2, respectively. The activation energy ofPeak 1 was estimated to be around 0.37 eV, while that of Peak 2, about 0.63 eV. The latter corresponded to the activationenergy of tritium diffusion in Li2TiO3. For Li4TiO4 exposed to tritium-deuterium mixture gas, two tritium release stages
named as Peaks A and B were also observed at 450 K and 600 K, respectively, in tritium-TDS spectrum at the heating rateof 0.5 K / min. As the release temperature regions of Peak 1 and Peak A were almost the same, tritium releases of thesepeaks were considered to be originated from a same process. O-T bonds formed on the surface of Li4TiO4 were
decomposed as Peak B in TDS spectra.

1. Introduction
For the development of D-T fusion reactors, it isimportant to establish an effective fuel recycling system
and a comprehensive model of tritium migration processes
in solid tritium breeding materials. In the test blanket
module for ITER, lithium titanate (Li2TiO3), one ofternary lithium oxides, is a candidate of tritium breeding
materials because of good tritium release property,
chemical stability, and so on [1]. During the operation of
D-T fusion reactor, energetic tritium is generated by thereactions of 6Li(n, D)T and 7Li(n, nD)T in Li2TiO3 andlithium will be consumed by these reactions, resulting inlithium-depletions in Li2TiO3. As lithium-depletion in
Li2TiO3 induces lower tritium generation rate, thelithium-depleted tritium breeding materials will be
replaced after certain periods of operations.
The use of Li2+xTiO3 (x=0.2, 0.4), which is thelithium-enriched Li2TiO3, is proposed recently to improve
the life cycle of tritium breeding materials with respect totritium generation property [2-4]. Li2+xTiO3 consists of
Li2TiO3 structure and excess lithium: i.e. Hara et al. [5],reported that both of Li2.2TiO3 and Li2.4TiO3 were themixture of Li2TiO3 and Li4TiO4 structures, indicating that
excess lithium induced the formation of Li4TiO4 structure
in Li2TiO3 structure during fabrication.
In our previous study, tritium release behaviors and theirkinetics for Li2+xTiO3 irradiated with thermal neutron were
investigated by means of TDS [6]. Tritium-TDS spectrum
for Li2.0TiO3 at the heating rate of 0.5 K/min consisted ofsingle release stage at around 580 K. An additional releasestage was observed at around 480 K for the tritium-TDS
spectra for Li2.2TiO3 and Li2.4TiO3, although the major
release stage was located at 580 K, which was assigned tothe tritium release controlled by tritium diffusion in
Li2TiO3 structure from kinetic analyses. In addition, theamount of tritium released at 480 K was increased withincreasing lithium concentration. It was considered that
this additional tritium release stage would be associatedwith excess lithium, that is, Li4TiO4 structure. However,
the effects of Li4TiO4 structure on tritium release forLi2+xTiO3 are still unclear because tritium releasemechanisms for Li4TiO4 have not been studied in detail.
Therefore, in this study, tritium release behavior for
Li4TiO4 was studied by means of TDS. Tritium was
generated in Li4TiO4 by thermal neutron irradiation or
introduced by thermal tritium gas absorption. The tritium
release kinetics for Li2+xTiO3 irradiated with thermal
neutron was also evaluated for comparison.

2. Experimental
Powders of Li4TiO4 and Li2+xTiO3 purchased from Kaken
Co. were used as samples. The average grain diameters of
Li4TiO4, Li2.0TiO3, Li2.2TiO3 and Li2.4TiO3 measured by
SEM (Scanning Electron Microscope) were 12, 3.0, 1.0
and 1.0 μm, respectively. For thermal neutron irradiation
experiments, 0.3 g of these samples were introduced into

f5144008@ipc.shizuoka.ac.jp

7©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

M. Kobayashi et al., Effects of Li4TiO4 Structure on Tritium Release Kinetics from Lithium-Enriched Li2+xTiO3

He

M.F.

CH4

M.F.

MS-5A

Furnace

Cu

SampleProportional
Counter 1

MS-5A

Proportional
Counter 2

CuO

Stack

Water bubbler 䐠

Water bubbler 䐟

Fig. 1 Schematic drawing of TDS systemof oxidized T and reduced T trapped in WB1 and WB2,
respectively, are quantified by a Liquid Scintillation
Counter (ALOKA LSC-5100). The CH4 gas at the flowrate of 25 sccm was used as a counting gas for the
proportional counters. The CH4 gas was humidified bypassing through another water bubbler for reduction ofmemory effect in proportional counters. The humidified
CH4 gas was converged to He gas flow just before PC1 sothat humidified CH4 is not purged directly to Mo cruciblewith sample during heating. The quartz tube, stainlesstubing of this system and the proportional counters werealso heated at around 373 K to reduce the memory effect.
In the tritium-TDS experiments, the samples were heated
up to 1173 K. Isothermal heating experiments were also
conducted around 600-700 K.
FT-IR (Fourier Transform Infrared Spectroscopy)
measurements were also carried out for Li4TiO4 exposed
to pure D2 gas at Shizuoka University to reveal chemical
states of deuterium in Li4TiO4. Changes of crystal
structure of Li4TiO4 by D2 gas exposure were also
investigated by X-ray diffraction (XRD) technique atUniversity of Toyama.

quartz tubes individually and annealed at 1073 K for
Li4TiO4 and 1173 K for Li2+xTiO3 under He gas with thepressure of less than a few Pa for 3 h. After heating, thesetubes were sealed under vacuum. Thermal neutron
irradiations were conducted at Pneumatic tube 2 (Pn-2) of
Research Reactor Institute, Kyoto University (KURRI).
The thermal neutron flux was 5.5 × 1012 n cm-2 s-1. The
thermal neutron irradiation was carried out for 10 min,corresponding to the thermal neutron fluence of 3.3 × 1015
n cm-2. The temperatures of the samples during irradiation
were estimated to be lower than 350 K.
For hydrogen isotope gas exposure, Li4TiO4 sample washeated at 1073 K for 3 h in vacuum under the pressurearound 10-5 Pa, then sample temperature was kept at 673
K. Thereafter, the tritium-deuterium mixture gas
consisting of D2 and DT (T/D = 0.005) or the pure D2(isotopic purity: 99.4%) gas was introduced into thesample with the total pressure of 200 Pa for 1 day. Afterthe hydrogen isotope gas exposure, the sample wasquickly cooled down by water and liquid nitrogen.
Out-of-pile tritium release experiments were performedin tritium-TDS system at Shizuoka University as shown in
Fig. 1. The sample put in Mo crucible was set in quartz
tube surrounded by a ceramic furnace. The He gas at theflow rate of 15 sccm was introduced into MS-5A and aheated Cu at 623 K to exclude impurities such as waterand oxygen in He gas, and then purged into the system
during experiments. The released tritium is recovered bypurge gas and is carried to the first proportional counter(PC1). The volume of PC is 30 cm3, and the voltage of 2.4
kV was applied to PC during measurement. After PC1,
purge gas with tritium passes through first water bubbler(WB1) to trap the water-form tritium such as HTO and
T2O, named as “oxidized T”. The second proportionalcounter (PC2) set after WB1 measures gas-form tritium
such as HT and T2, named as “reduced T”. Thereafter, thereduced T are oxidized by heated CuO at 623 K after PC2
and trapped by second water bubbler (WB2). The amounts

3. Results and discussion
Fig. 2 shows the tritium-TDS spectra for Li2+xTiO3 at theheating rate of 0.5 K / min measured by PC1, representing
the overall release of tritium. These spectra can beseparated into two release stages, namely Peak 1 at 480 K
and Peak 2 at 580 K. The TDS spectra for Li2.0TiO3
consisted of one release stage as Peak 2, althoughadditional release stage of Peak 1 was also found forLi2.2TiO3 and Li2.4TiO3. The amount of tritium for Peak 1
was increased with increasing lithium concentration. To
understand the tritium release kinetics for these peaks, theheating rate dependence experiments were carried out. In
thisanalysis,
Kissinger-Akahira-Sunose(KAS)
model-free-kinetics method [6] which has been widely
applied to analyze TDS spectra was adopted in the present

8

㻌

20

Li2.2TiO3
Li2.4TiO3

10

2

Peak 250300

Peak 1
Peak 2-15.0

ln (E / Tp )

15

-14.5

Li2.0TiO3

Peak 1㻌-1

Tritium release rate / Bq s g-1

M. Kobayashi et al., Effects of Li4TiO4 Structure on Tritium Release Kinetics from Lithium-Enriched Li2+xTiO3

-15.5
-16.0
-16.5
-17.0
-17.5

400

500 600 700 800
Temperature / K
Fig. 2 Tritium-TDS spectra of Li2+xTiO3
at the heating rate of 0.5 K / min

900

1.4

-1
-1

Tritium release rate / Bq s g

Table 1 Peak temperatures of each peak for Li2.4TiO3
at various heating rates
TP for Peak 2 (K)

0.5

453

591

1

515

603

5

580

705

10

601

718

2.2

1.2 eV [11], indicating that Peak 1 would not be
associated with Li2TiO3.
Fig. 4 shows the tritium-TDS spectrum for Li4TiO4irradiated with thermal neutron measured at the heating
rate of 0.5 K / min. This spectrum was measured by PC1
and no tritium release was found by PC2, indicating that
almost all of tritium was released as water form. The
tritium-TDS spectrum of Li4TiO4 can be separated into
two release stages at 450 K and 600 K, and named as Peak
A and Peak B, respectively. Especially, FWHM of Peak B
was quite narrow compared to the other spectra as shownin Fig. 2. The peak temperature of Peak A was almost
consistent with that of Peak 1.
Isothermal heating experiments were done for Li4TiO4exposed to tritium-deuterium mixture gas to determine thetritium release kinetics for Li4TiO4. Because of the
exposure temperature of 673 K, which is higher than therelease temperature of Peak A, the release kinetics of Peak
B was examined in these experiments especially. Two
probable processes which control the tritium release in
Li4TiO4 are considered as diffusion and de-trapping
processes. In the diffusion limited release process, tritium
would diffuse slowly toward the surface and release on thesurface quickly. In the de-trapping limited one, tritium
would be initially trapped strongly by the trapping sites.
Tritium release would be initiated by de-trapping from the§ E ·§ RA · E a¸(1)
ln¨ 2 ¸ ln¨¨¸¨T ¸
© E a g (D ) ¹ RTP© p ¹
where R is gas constant, A is pre-exponential factor, Ea is
activation energy of reaction. The term of g(D) is amodel-function determined by rate limiting process ofreaction. The linear relation between ln(E/TP2) and 1/TP
was deduced as Eq. 1, and the slope of which isrepresenting -Ea/R. The peak separation analysis was done
for TDS spectra of Li2.2TiO3 and Li2.4TiO3 under various
heating rates, then peak temperatures for Peaks 1 and 2 ineach spectrum were estimated. Table 1 summarizes thepeak temperatures of Peaks 1 and 2 for Li 2.4TiO3 atvarious heating rates. The peak temperatures were shifted
toward higher temperature region with increasing heatingrate. Fig. 3 shows the ln(E/TP2) vs 1/TP plot for Peaks 1
and 2 for Li2.4TiO3. Liner relations between ln(E/TP2) and1/TP were found for each peak. The obtained activationenergies of these release peaks were evaluated to be 0.37
eV and 0.63 eV for Peaks 1 and 2, respectively. The
similar activation energies were also obtained forLi2.2TiO3. Especially, the activation energy of Peak 2 wasconsistent with that of tritium diffusion for Li2TiO3obtained in our previous study and other works [7, 8]. The
activation energy around 0.4 eV has not been reported for
Li2TiO3 with neutron irradiation [8-10]. In addition, theactivation energy for desorption of hydrogen isotopes
adsorbed on the surface of Li 2TiO3 was reported as about

TP for Peak 1 (K)1.8 -1 2.0
1000/Tp / K

Fig. 3 Analytical results of release kinetics
for Peaks 1 and 2 in Li2.4TiO3study. In this method, ln(E/TP2) are plotted against 1/TP for
the determination of activation energy, where E is theheating rate and TP is the peak temperature, based onfollowing equation,

Heating rate (K / min)

1.6

30

Peak A

20

Peak B

10
0
300

400

500 600 700
Temperature / K

800

Fig. 4 Tritium-TDS spectrum of Li4TiO4with thermal neutron irradiation

9900

M. Kobayashi et al., Effects of Li4TiO4 Structure on Tritium Release Kinetics from Lithium-Enriched Li2+xTiO3

trapping sites. The other processes are quite faster than thede-trapping process, so these processes are not apparent on
tritium release kinetics. The tritium release controlled bydiffusion process is expressed by Fick’s diffusion equation[12]. In particular for tritium release via diffusion in
spherical media, tritium release with respect to elapsed
time, t can be expressed in the following equation,
f§  Dn 2S 2 t ·61

¸
¨
D
1 2
exp
2
2
¸
¨
S
n
a

¹
©
n 1
where D is fraction of released tritium to total amount, a is
the radius of media, namely grain radius and D isdiffusivity. On the other hand, the rate of de-trapping
limited release should be controlled only by theconcentration of trapped tritium in the material, indicatingthat de-trapping limited release is the first-order reaction,
which is expressed as follows,k (1  D )

D

(3)

㻌

D

0.2
0

5
3
Time / 10 s

5

103Time / 10 s

15

decomposition energy of hydrogen from O-H bonds [14].
A FT-IR spectrum for Li4TiO4 exposed to pure D2 gas
showed the IR absorption in the wavenumber range of
2450-2750 cm-1 (not shown here). It was reported thatO-D bonds in lithium ceramics showed IR absorption inthis wavenumber region[15], indicating that O-D bonds
were formed in Li4TiO4 by D2 gas exposure. Therefore, it
was considered that tritium trapped as O-T bonds werede-trapped by the decomposition of O-T bonds andreleased as Peak B for Li4TiO4.
However, the activation energy of Peak 2 in Li2+xTiO3
was different from that of the decomposition of O-T bonds
in Li4TiO4. In addition, the tritium retentions as Peak 2 inLi2+xTiO3 were almost similar to that of Li2TiO3, althoughmore than half of tritium was retained as O-T bonds in
Li4TiO4 with neutron irradiation. These facts would
indicate that the tritium releases derived from O-T bonds
should be small in Li2+xTiO3.
The crystal structure of Li4TiO4 before and after D2 gas
exposure were compared by XRD method as shown in Fig.
7. In the measurements, silicon was used for reference.
The diffraction patterns of Li4TiO4 exposed to D2 gas werequite similar to that without D2 gas exposure, indicatingthat the crystal structure was not changed by D2 gas
exposure. These results indicated that LiOD structure(crystal) was not formed in Li4TiO4. It was reported that
Li4TiO4 is reactive with water vapor and forms O-H bondson the surface [4], considering that water vapor as
impurity was adsorbed on the surface of Li4TiO4 during
air exposure or TDS measurements, which would induce
the formation of O-H bonds. Similar results have beenobtained in Li2O [13]. The rate limiting process of tritium
release for Li2O was changed by heating temperature: i.e.
in a lower temperature region, tritium release wascontrolled by decomposition of surface O-T bonds. In ahigher temperature region, tritium release was governedby diffusion process. The rate limiting process waschanged at 570 K, and this temperature was closed to Peak
B appearing in this study, indicating that O-T bonds was
mainly formed on the surface of Li4TiO4 and decomposed

0.4

0.0

0

Fig. 6 Isothermal tritium release for Li4TiO4exposed to tritium-deuterium mixture gas

673 K
623 K
detrapping model
diffusion model

0.6

0.4

0.0

㻌

0.8

0.6

0.2

where k is the rate constant of reaction [13]. Comparing
the theoretical curve derived from Eqs. 1 and 2 withtritium release by isothermal heating, the rate limiting
process on tritium release was determined. Fig. 5 shows
the isothermal tritium release behavior at various
temperatures for Li4TiO4 exposed to tritium-deuterium
mixture gas. The theoretical curves by de-trapping limited
release model and diffusion limited release model werealso shown in this figure. As shown in Fig. 5, theoreticalcurves based on diffusion limited release model could
hardly reproduce isothermal tritium release for Li4TiO4.
The theoretical curves by de-trapping limited releasemodel could reproduce tritium release curve in thetemperature region around 600-700 K, as shown in Fig. 6,
indicating that tritium release for Li4TiO4 was controlled
by de-trapping process.
The activation energy of tritium de-trapping was derived
to be about 1.2 eV from the de-trapping rate constants at
various temperatures, which was close to the

1.0

673 K
648 K
623 K
598 K
detrapping model

0.8

¦

dD
dt

㻌

1.0

10

Fig. 5 Comparison of isothermal tritium release totheoretical curves derived from de-trappingand diffusion processes

10M. Kobayashi et al., Effects of Li4TiO4 Structure on Tritium Release Kinetics from Lithium-Enriched Li2+xTiO3Intensity / arb. unit

200

Before D2 exposure

‫ە‬

300

After D2 exposure

‫ە‬

‫ی‬

‫ی‬
‫ی‬

‫ی‬

‫ە‬

‫ە‬

‫ی‬

100
0
20

30

40

50
2 theta / deg.

60

70

80

Fig.7 X-ray diffraction patterns for Li4TiO4 before and after D2 gas exposure‫ ی‬indicates Li4TiO4 structure and ‫ ە‬indicates reference silicon [5]release stages named as Peaks A and B were also observed
at 450 K and 600 K, respectively, in tritium-TDS spectrumat the heating rate of 0.5 K / min. As the releasetemperature regions of Peak 1 and Peak A were almost thesame, tritium releases of these peaks were considered to
be originated from the same process. The O-T bonds were
easily formed on the surface of Li4TiO4, appearing as Peak
B in TDS spectrum.

as observed as Peak B. The small tritium retention as O-T
bonds in Li2+xTiO3 would imply that the amount of wateron the surface of Li4TiO4 in Li2+xTiO3 was small. Because
the Li2+xTiO3 was the mixtures of Li2TiO3 and Li4TiO4,
the effective surface area of Li4TiO4 in Li2+xTiO3 should
be small, resulting in the small retention of tritium as O-T
bonds.
Peak B appearing in the TDS spectrum of Li 4TiO4 could
be assigned to the release of tritium as O-T bonds in thisstudy. However, the assignment for Peak A was notcompleted although this peak also appeared in TDS
spectra of Li2+xTiO3. The water vapor on the surface of
Li4TiO4 may provide trapping sites for tritium as O-T
bonds. A part of tritium migrated from bulk region of
Li4TiO4 at lower temperatures, corresponding to tritium
release as Peak A, would be trapped on the surface atlower temperature region. Li4TiO4 would easily react withwater vapor in air, therefore, the experiments for Li4TiO4without any air exposure are required to elucidate thetritium retention and release as Peak A, and consequently
the relations of tritium retention as Peak 1 and the amountof excess lithium in Li2+xTiO3. The experiments in vacuum
system would be a promising way to exclude the effects ofsurface O-T bonds. Therefore, it was considered thathydrogen isotope gas exposure and subsequent TDS
experiments in the same vacuum system would reveal thetritium release mechanisms from Li4TiO4 in detail.

Acknowledgements
This work was supported by the Grant-in-Aid for
Scientific Research (A), KAKENHI 20246131.

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[5] M. Hara et al., J. Nucl. Mater. 404, 217–221 (2010).
[6] T. Ozawa, Netsu Sokutei 31(3), 133 (2004).
[7] M. Kobayashi et al., Fusion Eng. Des., 87, 471 (2012).
[8] T. Kulsartov et al., Fusion Sci. Technol., 60 1139 (2011).
[9] C. Alvani et al., Fusion Eng. Des., 69, 275 (2003)[10] S. Casadio et al., J. Nucl. Mater., 329, 1252 (2004).
[11] K. Hashimoto et al., Fusion Eng. Des., 61 375 (2002).
[12] K. Okuno and H. Kudo, Fusion Eng. Des. 8, 355 (1989).
[13] H. Kudo and K. Okuno, J. Nucl. Mater., 101 38 (1981).
[14] K. Okuno and H. Kudo, J. Nucl. Mater., 133, 192 (1985).
[15] T. Oda et al., J. Nucl. Mater. 346, 306 (2005).

4. Conclusions
The release kinetics of tritium for Li2+xTiO3 andLi4TiO4 were studied by means of TDS. Tritium-TDS
spectrum with the heating rate of 0.5 K / min for Li2+xTiO3
irradiated with thermal neutron consisted of two releasestages at 480 K and 580 K, namely Peaks 1 and 2,
respectively. The activation energy of Peak 1 wasestimated to be around 0.37 eV, while that of Peak 2 wasabout 0.63 eV. The latter one corresponded to theactivation energy of tritium diffusion in Li 2TiO3. For
Li4TiO4 irradiated with thermal neutron, two tritium

11

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Release behavior of hydrogen isotopes in Li2TiO3 pellet
Deqiong Zhu, Takuji Oda and Satoru Tanaka
Department of Nuclear Engineering and Management, The University of Tokyo(Received: 25 May 2012 / Accepted: 20 August 2012)

Release behavior of hydrogen isotopes thermally sorbed in Li2TiO3 pellet was studied by thermal desorption
spectroscopy (TDS). As microstructure (grain size, surface condition, open and closed pores) of pellet samples werechanged by high-temperature sintering, the release behavior of hydrogen isotopes was changed. The amount ofhydroxyl groups chemically adsorbed on the surface was found to increase with the specific surface area. The gas
migration in open pores was delayed to high temperature region (450~650 K), for which Knudsen diffusion andinteraction between gas molecules and walls of narrow open pore channels is considered to be the main mechanism.
The release of hydroxyl groups absorbed in the bulk was delayed due to the existence of closed pores, and thesehydroxyl groups have to undergo bulk diffusion and trapping in closed pores several times before desorption.
Keywords: tritium breeder, hydrogen isotopes, Li2TiO3, pellet, microstructure

Instrument Corporation) was adopted to measure the open
porosity, open pore size distribution. The specific surface
area was measured by BET (Kr adsorption at 77 K) on an
Autosorb-1-C (Quantachrome.Co). The influence of
microstructure (grain size, specific surface area, open andclosed pores) on hydrogen isotopes release behavior was
discussed.

1. Introduction
The tritium release in ceramic breeder materials is acomplex process, in which the contributions of bulk
diffusion, surface processes and interaction with radiation
defects were included as investigated in previous studies[1-4]. Just a few literatures were reported on the influenceof porosity on tritium release behavior. Tanifuji et al.
investigated the porosity dependence of tritium release inneutron-irradiated Li2O sintered pellets. It was shown thatthe tritium release behavior is significantly affected evenby a slight porosity change in pellet sample with densities
above 87% T.D., and that tritium release can be delayed
due to trapping in closed pores [5]. Peeters et al. compared
the tritium residence time of different Li2TiO3 sample
materials in the experiments EXOTIC-8 and
EXOTIC-9/1, and suggested that tritium release improves
with increasing open porosity and the closed porosity has
no effect on tritium release behavior [6]. Tam et al.
developed a random-lattice approach to study the tritium
percolation through porous ceramic breeders by modelingthe short range transport in pores via a
convection-diffusion-reaction approach and long range
transport via a matrix technique [7].
In the present work, Li2TiO3 pellet samples withdifferent microstructure (grain size, specific surface area,
open and closed pores) were prepared. The release
behavior of hydrogen isotopes thermally sorbed in pellet
samples was studied by thermal desorption spectroscopy(TDS).
Scanning
electronmicroscope(SEM,
JSM-6510LA, JEOL, Ltd.) was used to observe the poremorphology and grain size. Mercury intrusion
porosimetry (autopore(R) IV 9520, Micromeritics

2. Experiment
Li2TiO3 pellet samples were prepared by pressing thepowder into pellet and sintering. The starting materials of
Li2TiO3 poly crystal powder, PVA (5 wt%) and stearic acid(5 wt%) were ball-milled for 6 h and mixed
homogeneously. The mixture of powders was pressed into
pellets stepwise at 20 kN (2 minutes), 40 kN (2 minutes)
and 60 kN (2 minutes) per 1.3 cm2. By sintering theas-prepared pellets in air at different temperatures for 12 h,
five pellet samples of various apparent densities (% oftheoretical density: 3.43 g.cm-3 for Li2TiO3) were obtained:(a) 76% T.D., 1173 K, (b) 83.5% T.D., 1223 K, (c) 86.6%
T.D., 1273 K, (d) 88.6% T.D., 1373 K and (e) 88.6% T.D.,
1473 K. Note that these apparent densities were evaluatedby measuring the mass and size of pellets. The color of all
the pellet samples is white after sintering in air at hightemperatures.
Deuterium was incorporated into pellet samples byheating in 100 Pa D2O vapor at 1073 K for 1 h inabsorption system. After cooling down and evacuation, thesamples were transferred into a high-vacuum TDS system,and then were heated to 973 K at 5 K/min for TDS
experiment. TDS spectra of released gas versus
temperature or time were obtained during the heatingprocess. Figure 1 shows the photographs of five pellet

zhudeqiong@flanker.n.t.u-tokyo.ac.jp

12©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

D. Zhu et al., Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet㻝㻚㻞㼤㻝㻜㻙㻤
㻼㼍㼞㼠㼕㼍㼘㻌㼜㼞㼑㼟㼟㼡㼞㼑㻛㻼㼍

samples after TDS experiments. The gray color in samples(a) and (b) and darker color in samples (c) to (e) wereobserved. It was reported that Li containing species (Li andLiOD) were observed in mass spectra for Li2TiO3 byheating above 973 K in D2O atmosphere. And the color of
Li2TiO3 changed from white to light blue after heated in
D2O atmosphere, which was attributed to be due to theoxygen defects induced by LiOD vaporization [8]. In thepresent work, the color of Li2TiO3 pellet samples changed
to gray or black after heated in 100 Pa D2O vapor, which issimilarly considered to be induced by reduction.

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㼀㼑㼙㼜㼑㼞㼍㼠㼡㼞㼑㻛㻷㻝㻜㻜㻜

Fig. 2 TDS spectra for Li2TiO3 pellet samples: (a) 76% T.D., (b)83.5% T.D., (c) 86.6% T.D., (d) 88.6% T.D. (1373 K), (e) 88.6%
T.D. (1473 K)

㻝㻜㻜㻜

㼎

Figures 2 show the TDS spectra for Li2TiO3 pellet
samples. Deuterium thermally sorbed in pellet samples was
released mainly in HDO and D2O. One peak was observedat around 360 K for HDO, D2O and H2O, named peak (1).
For H2O, one more peak was observed at around 625 K,
named peak (2). Except sample (e), no obvious HDO

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3.1. Release behavior of hydrogen isotopes㻝㻜㻜㻜㻟㻚㻜㼤㻝㻜㻙㻤

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3. Results and discussion㻢㻜㻜㻤㻜㻜㼀㼑㼙㼜㼑㼞㼍㼠㼡㼞㼑㻛㻷㻌㻴㻞㻻㻝㻚㻞㼤㻝㻜㻙㻤

Fig. 1 Photographs of pellet samples after TDS experiment

㼏

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13

D. Zhu et al., Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet

the present experiment, obvious H2O peak with higher
level than HDO and D2O was observed at around 625 K.
Three possibilities were considered for this H2O peak: (i)–OH formed due to H2O residue in D2O tank and thendesorbed as H2O after recombination, (ii) –OH formed byisotope exchange reaction between –OD on the sample
surface and H2O residue in TDS system and thendesorbed as H2O after recombination, and (iii) isotope
exchange reaction between released molecules (HDO,
D2O) and H2O residue in TDS system.
The quantitative analysis for peak (1), peak (2) andpeak (3), and assignment of peak (3) and peak (4) will be
discussed in detail in section 3.4.

release corresponding to peak (2) was observed. A peak for
HDO and D2O was observed at 650~900 K in samples (b)to (e), named peak (3). In sample (d), a broad peak
between peak (1) and peak (3) was observed at 450~650 K
for HDO and D2O, named peak (4). For peak (3) and peak(4), no obvious H2O peak was observed.
Based on the previous studies [9, 10], peak (1) is
assigned to molecules physically adsorbed on the surface,which is released through desorption of physicallyadsorbed water molecules and migration along open porechannels, and peak (2) to hydroxyl groups chemicallyadsorbed on the surface, which is released byrecombination desorption or isotope exchange reactions. In
order to assign peak (3) and peak (4), the D/H ratios forpeaks (1) to (4) were calculated and compared in Table 1.
For peaks (3) and (4), no H2O peak was obviously
observed, which could be submerged in the high-levelbackground. In order to give more precise evaluation of
D/H ratios for peaks (3) and (4), the H2O peak was
estimated from HDO and D2O peaks by assuming theequilibrium for the following reaction:
H2O + D2O = 2HDO
(1)The D/H ratios for peaks (4) and (5) evaluated
without (shown in brackets) and with considering theestimated H2O peaks were compared.

3.2. SEM observation

Table 1 D/H ratios for peak (1) to peak (4)Sampleab

peak (1)~ 360 K
0.96
0.91

c

0.82

d

0.89

e

0.86

D/H ratio
peak (2)peak (3)~ 625 K 650~900 K
1.8
(2.8)
1.9
(2.9)
1.9
(2.9)
0.01
1.9
(2.9)

peak (4)450~650 K
1.9
(2.9)
-

The D/H ratio reflects the probability of isotope
exchange reaction between deuterium species (D2O, HDO,
-OD) and H2O residue in TDS system. For peak (1),isotope exchange reaction between released molecules(HDO, D2O) and H2O/-OH residue in TDS system partiallyoccurred. For peak (2) with the lowest D/H ratio of around
0.01, it is considered that isotope exchange reaction
between –OD on the sample surface and H2O residue in
TDS system largely happened. The D/H ratios for peak (3)and peak (4) are larger than that for peak (1) and peak (2),and similar to that of D2O tank (D/H ratio=1.8). It indicatesthat the probability of isotope exchange reaction is muchsmaller in peak (3) and peak (4), which suggests that theorigins of peak (3) and peak (4) are located at some places
where residual H2O gas in TDS system hardly reaches. In

Fig. 3 SEM observation for Li2TiO3 pellet samples: (a) 76% T.D.,(b) 83.5% T.D., (c) 86.6% T.D., (d) 88.6% T.D. (1373 K), (e)88.6% T.D. (1473 K); 1- surface, 2- cross section

SEM observation for Li2TiO3 pellet samples is shownin Fig. 3. With the density increasing, grain sizes increased,
and the porosity decreased. A broad grain size distribution
was observed in sample (c) and especially obvious insample (d) (Fig. 3.c-1 and d-1). Small pores with dead-end
appeared in samples (b) to (e), and especially obvious insamples (d) and (e) (Fig. 3.b-2, c-2, d-2 and e-2). In
samples (d) and (e), the grains in the bulk melted to form

14

D. Zhu et al., Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet

amount of peak (2) (hydroxyl groups chemically adsorbedon the surface) was evaluated from the area of H2O peak ataround 625 K (HDO peak was added in sample (e)). The
specific surface area of pellet samples was measured by
BET.

big grain (Fig. 3.d-2 and e-2).

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3.3. Mercury intrusion porosimetry
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㻜㻚㻜㻜㻤

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㻌㼎㻚㻌㻤㻟㻚㻡㻑㻌㼏㻚㻌㻤㻢㻚㻢㻑㻌㼐㻚㻌㻤㻤㻚㻢㻑㻌㼑㻚㻌㻤㻤㻚㻢㻑

Table 2 Correlation between specific surface area and amount ofpeak (1) and peak (2)

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㻝㻜㻜

㻝㻜㻜㻜
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㻼㼛㼞㼑㻌㼐㼕㼍㼙㼑㼠㼑㼞㻛㼚㼙

Fig. 4 Open pore size distribution for Li2TiO3 pellet samples: (a)76% T.D., (b) 83.5% T.D., (c) 86.6% T.D., (d) 88.6% T.D. (1373
K), (e) 88.6% T.D. (1473 K)

The open porosity and open pore size distribution were
measured by mercury intrusion porosimetry under thepressure of 6.8 MPa, for which pores smaller than 183 nm
could not be measured. As a result, the open porosity andopen pore size distribution could be underestimated tosome extent. The open porosity was calculated from theopen pore volume (volume of mercury intruded into
sample) and external sample volume. The open porosity
for pellet samples is: (a) 12.7%, (b) 11.9%, (c) 10.9%, (d)9.8% and (e) 10.6%, which decreased from sample (a) tosample (d), and slightly increased in sample (e). Figure 4
shows the open pore size distribution, which was
determined from the mercury pressure by the Washburnequation assuming that the geometry of pore channels canbe approximated by cylindrical capillaries [11]. It seems
that there are some pores smaller than 168 nm existing inpellet samples especially in sample (a). With the density
increasing, the open pore size decreased in samples (b) and(c), and then increased in samples (d) and (e). With thesintering temperature increasing, the pores shrank first in
samples (b) and (c), then expanded in samples (d) and (e)due to over-sintering at high temperature (1373 K and1473 K). A small fraction of pores with larger size aroundseveral microns were also observed in sample (e). The
closed porosity was evaluated from the measured open
porosity and apparent density for pellet samples as [closed
porosity] = (1- [apparent density/theoretical density]) [open porosity]: (a) 11.3%, (b) 4.6%, (c) 2.5%, (d) 2.4%and (e) 0.8%. Due to the underestimation of open porosity
by mercury intrusion porosimetry, the closed porosity
might be overestimated especially in sample (a).

SampleSpecific surface
area of pellet /m2·g-1

Peak (1)molecules/
1017 g-1

Peak (2)hydroxylgroups/ 1016 g-1

a
b

0.762
0.120

22.5
2.2

11.2
6.4

c
d

0.035
0.058

0.9
1.1

1.3
1.5

e

0.039

0.4

19.8

The amount of molecules physically adsorbed on thesurface (peak (1)) is proportional to the specific surface.
Except sample (e), the amount of hydroxyl groupschemically adsorbed on the surface (peak (2)) is roughly
proportional to the specific surface area in samples (a) to(d). It is considered to be due to that the concentration of
surface defects induced by vaporization of Li-containing
species is different in sample (e) sintered at the highest
temperature (1473 K).
The vaporization characteristics of Li2TiO3 have been
investigated by the mass spectrometric Knudsen effusion
method, and it was shown that the vaporization oflithium-containing species increased with the temperature
increasing [8, 12]. Hoshino et al. investigated thevaporization of lithium containing species under different
atmospheres (D2, D2O, vacuum and O2) from 973 to 1473
K [8], and the sum of partial pressures of Li-containing
species (Li, LiO, Li2O and LiOD) was found to be: D2 =D2O > vacuum > O2. In the present study, pellet samples
were prepared by sintering in air, where Li 2TiO3 should
show an intermediate vaporization behavior between O2
atmosphere and D2O atmosphere due to the existence ofmoisture (H2O, similar to D2O) and O2. In another work[13], the non-stoichiometry of Li2TIiO3 heated in
reduction atmosphere was quantitatively analysed bythermogravimetry. The mass of Li2TiO3 was found to
decrease with time and the weight loss due to Li 2O and O
vaporization reached the maximum after heating at 1273
K for around 16 h in H2 atmosphere. And the maximum
weight loss due to Li2O vaporization was found to be
0.982 mg of 288.5 mg, which results in Li 2O deficiency
with molar fraction of 0.01. The amount of defects
induced by sintering in air (without O vaporization) is
considered to be smaller than this value in samples (a) to

3.4. Discussion
3.4.1. Quantitative analysis for peak (1) and (2)Table 2 shows the correlation between the specificsurface area of pellet and the amount of peak (1) and peak(2). The amount of peak (1) (molecules physicallyadsorbed on the surface) was evaluated from the area of
D2O, HDO and H2O peaks at around 360 K, and the15

D. Zhu et al., Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet(d), for which the influence from the defects on therelease behaviour of hydrogen isotopes is considered tobe negligible. However, the molar fraction of Li 2O might
be close to or larger than 0.01 in sample (e) which was
sintered at higher temperature (1473 K), and the influenceof surface defects on the amount of surface hydroxyl
groups is expected in sample (e).
Oda et al. [14] studied the influence of radiation
defects on hydrogen isotope behavior in Li2O by in situ IR
absorption analysis, and suggested that hydrogen isotopes
were stabilized in Li2O through an interaction with the Li
vacancy or mutual aggregation. Bikondoa et al. [15] used
STM to image the reaction of water molecules withbridging-oxygen vacancies on rutile TiO2 (110) and oxygen
vacancies were visually observed being transformed into
OH species as a water molecule dissociates in the vacancy.
It indicates that oxygen vacancies act as active sites forhydroxyl groups. In the present work, defects induced by
Li species vaporization might facilitate relatively larger
amount of hydroxyl groups adsorbed on the surface,especially in sample (e).

diffusion length needed to reach the surface (or open
pores). The “effective grain size” is calculated fromtemperature of peak (3) by using a simple diffusion
model, the detail of which was reported in the previous
study [9]. The real grain size of pellet samples was decided
by SEM observation. Table 3 shows the correlation
between the real grain size and “effective grain size”. The“effective grain size” is several times larger than the real
grain size. It indicates that hydroxyl groups mightexperience bulk diffusion in multiple grains due to theexistence of closed pores. It is generally known that the
migration along the grain boundaries is fast. The surface
processes are considered to proceed fast at thetemperature region of peak (3) (650~900 K) (the peak
was observed at 625 K in the present work). Besides,
peak (3) was not observed in 5 μm and 50 μm single
crystal powder samples [10]. Accordingly, the release ofhydrogen isotopes is considered to be jointly controlled
by diffusion in the bulk, trapping in the closed pores anddissolution in the grains at the temperature region of peak(3). When the hydroxyl groups come to the surface of
closed pores, molecules desorbed from the grain surface
can be trapped in the closed pores, dissolve in grain anddiffuse in the grain bulk again. After several-times bulk
diffusion and trapping in closed pores, the hydroxyl
groups can arrive at the surface of pellet or open poresand finally be released to the TDS system. No obvious
dependence of hydrogen isotopes release temperature on
the grain sizes decided by SEM was observed in thepresent work.

3.4.2. Assignment of peak (3)Peak (3) was observed at higher temperature region(650~900 K) than the peak due to hydroxyl groupschemically adsorbed on the surface in high-density pellet
samples (d) to (e) and might located at some places whereresidual H2O gas in TDS system hardly reaches. Here twokinds of possible origins for peak (3) are considered: (i)hydroxyl groups absorbed in the bulk, (ii) molecules
trapped in closed pores. In order to clarify the origin ofpeak (3), the amount of molecules/hydroxyl groups forpeak (3) was estimated from the area of D2O and HDO
peaks in TDS spectra, which is around 1015.
In the case of (ii) molecules trapped in closed pores,
the pressure needed to trap 1015 molecules was roughly
estimated. By assuming the closed porosity in high-density
pellet samples (d) to (e) to be 5%, the volume of closed
pores is estimated to be around 0.005 cm3 (the volume ofpellet is around 0.1 cm3). Based on the ideal gas law
equation, around 800 Pa of D2O vapor is needed to trap
1015 molecules at 300 K. However, 100 Pa D2O vapor wasapplied in the present experiment. Here we can infer thatmolecules trapped in closed pores are not a main source ofpeak (3).
In the previous study, 7.8×1016 cm-3 hydroxyl groups
were thermally absorbed in the bulk of Li2TiO3 single
crystal plate after heating in 100 Pa D2O vapor at 1073 K
for 1 h [10]. The amount of hydroxyl groups for peak (3) isat the same level, which convinced us that peak (3) can be
assigned mainly due to (i) hydroxyl groups absorbed in thebulk.
In order to make clear the release path for peak (3),the “effective grain size” is defined to be an effective

Table 3 Correlation between the real grain size and “effective
grain size” for peak (3)SampleReal
grainsize/ μm

Peaktemperature/K“Effective
grainsize”/ μm

a

~1

-

-

“Effective
grain size”/
real grain
size
-

b
c

1~3
3~10

700
775

12
50

4~12
5~14

d
e

>5
>10

750
780

32
55

<6
1~5

3.4.3. Assignment of peak (4)For peak (4) detected in sample (d), a similar peakalso appeared between peak (1) and peak (2) in 5 μm (av.)
Li2TiO3 single crystal powder sample, in which sever
agglomeration was observed [10]. The possible origin ofwater molecules physically adsorbed on the surface of
inner open pores is considered, which is released throughmigration along complicated-structure pore channels. In
the simulation by using the random-lattice approach

16

D. Zhu et al., Release Behavior of Hydrogen Isotopes in Li2TiO3 Pellet

developed by Tam, in which convection-diffusion-reaction
approach was considered for short range transport in open
pores, a sharp pulse of tritium-bearing species was
injected at the center of interconnected pores. And therelease of injected tritium-bearing species exhibited abroad distribution with the long tail, which indicatesmany different pathways through which the gas moleculecan reach the outlet from the source [7]. In the present
work, two kinds of open pore channels are assumed: (I)pore size much larger than molecule mean free path, (II)pore size smaller than molecule mean free path. In the case(I), interaction between gas molecules and walls of porechannels is negligible, and the gas diffusion is fast, whichis released in peak (1). In the case (II), interaction between
gas molecules and walls of pore channels has to be taken
into account. The broad peak (4) is considered to bemainly due to Knudsen diffusion and interaction between
gas molecules and walls of different-shape narrow open
pore channels [16, 17]. It was reported that Knudsen
diffusion is dominant in pores with size of several (tens)
nm [18, 19].In the present study, no obvious peak (4) was
observed except sample (d). Hence, we may think thatnarrow pore channels were evolved in the sintering
condition of sample (d). However, small pore channels of
several (tens) nm could not be measured by mercury
intrusion porosimetry. The detailed mechanism for peak (4)needs further study.

ABDOU, Modeling of tritium transport in ceramic breeder
single crystal, J. Nucl. Mater. 210 (1994) 143-160.
[3] M. Oyaidzu, H. Kimura, A. Yoshikawa, et al., Correlation
between annihilation of irradiation defects and tritium
release in neutron-irradiated lithium zirconate, Fusion Eng.
Des. 81 (2006) 583–588.
[4] Makoto Oyaidzu, Yusuke Nishikawa, Taichi Suda, et al.,Detrapping behavior of tritium trapped via hot atomchemical process in neutron-irradiated ternary lithiumoxides, J. Nucl. Mater. 375 (2008) 1-7.
[5] Takaaki Tanifuji, Daiju Yamaki, Tadashi Takahashi, Akira
Iwamoto, Tritium release from neutron-irradiated Li2O
sintered pellets: porosity dependence, J. Nucl. Mater.
283-287 (2000) 1419-1423.
[6] M.M.W. Peeters, A.J. Magielsen, M.P. Stijkel, J.G. van derLaan, In-pile tritium release behaviour of lithium
metatitanate produced by extrusion–spheroidisation–sintering process in EXOTIC-9/1 in the high flux reactor,
Petten, Fusion Eng. Des. 82 (2007) 2318–2325.
[7] S.W. Tam, V. Ambrose, Tritium percolation through porous
ceramic breeders – A random-lattice approach, Fusion Eng.
Des.17 (1991) 43-48.
[8] T. Hoshino, et al., Vapor species evolved from Li2TiO3heated at high temperature under various conditions,
Fusion Eng. Des. 81 (2006) 555–559.
[9] Deqiong Zhu, Takuji Oda and Satoru Tanaka, Effect ofgrain size on hydrogen isotope behavior in LiNbO3,
Fusion Sci. Technol. 60 (2011) 1147-1150.
[10] Deqiong Zhu, Takuji Oda, Yohei Shono, Satoru Tanaka,
Study of hydrogen isotopes inventory in Li2TiO3, J. Nucl.
Mater. Submitted.
[11] J. Van Brake1, S. Modrý, M. Svatá, Mercury Porosimetry:
State of the Art, Powder Technol. 29 (1981) 1-12.
[12] Hamazo Nakagawa, Mitsuru Asano, Kenji Kubo, Mass
Spectrometric Investigation of the Vaporization of Li2TiO3,
J. Nucl. Mater. 110 (1982) 158-163.
[13] T. Hoshino, M. Dokiya, T. Terai, et al., Non-stoichiometry
and its effect on thermal properties of Li2TiO3, Fusion Eng.
Des. 61-62 (2002) 353-360.
[14] Takuji Oda, Yasuhisa Oya, Satoru Tanaka, Dynamic
observation of the behavior of 3 keV D2+ irradiated into
Li2O using IR absorption spectroscopy, J. Nucl. Mater.
346 (2005) 306-311.
[15] Oier Bikondoa, Chi L. Pang, Roslinda Ithnin, et al., Direct
visualization of defect-mediated dissociation of water on
TiO2 (110), Nat. Mater. 5 (2006) 189-192.
[16] Jaime Benítez, Principles and Modern Applications ofMass Transfer Operations (A JOHN WILEY&SONS,INC.,
New Jersey, 2009), p.59.
[17] Douglas M. Ruthven, Principles of Adsorption andAdsorption processes (A Wiley & Sons, New York, 1984),
p 136.
[18] Kourosh Malek, Marc-Olivier Coppens, Knudsen selfand Fickian diffusion in rough nanoporous media, J.
Chem. Phys. 119 (2003) 2801-2811.
[19] Aaron W. Thornton, Tamsyn Hilder, Anita J. Hill, James
M. Hill, Predicting gas diffusion regime within pores ofdifferent size, shape and composition, J. Membrane. Sci.
336 (2009) 101-108.

4. Summary and Future work
The release behavior of hydrogen isotopes
thermally sorbed in Li2TiO3 pellet was studied by
TDS. The release behavior of hydrogen isotope was
obviously changed due to the change in the
microstructure of materials by high-temperature
sintering. The amount of hydroxyl groups chemicallyadsorbed on the surface increases with the specificsurface area. The release of hydroxyl groups in thebulk was delayed due to the existence of closed pores
through undergoing bulk diffusion in multi-grains andtrapping in closed pores several times. The peakrelated to gas migration in open pores was observed atrelatively high temperature region (450~650 K),
which is considered to be due to Knudsen diffusion
and interaction between gas molecules and walls ofnarrow open pore channels, which needs further
study.

References
[1]

[2]

Peter C. BERTONE, The kinetics that govern the releaseof tritium from neutron-irradiated lithium oxide, J. Nucl.
Mater. 151 (1988) 281-292.
A. RENÉ RAFFRAY, Seungyon CHO, Mohamed A.

17

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Mass loss of Li2TiO3 pebbles in atmosphere containing hydrogen
Hideaki Kashimuraa,* , Masabumi Nishikawaa, Kazunari Katayamaa,
Shohei Matsudaa, Satoshi Fukadaa, Tsuyoshi HoshinobaInterdisciplinary Graduate School of Engineering Science, Kyushu University, 6-10-1, Hakozaki, Higashi-ku, Fukuoka
812-8581, Japanb

Blanket Irradiation and Analysis Group, Fusion Research and Development Directorate, Japan Atomic Energy Agency,
2-166, Oaza-Obuchi-Aza-Omotedate, Rokkasho-mura, Kamikita-gun, Aomori 039-3212, Japan.
(Received: 8 May 2012 / Accepted: 24 September 2012)

It has been found that water vapor is released from the ceramic b reeder materials into the purge gas due todesorption of adsorbed water and the water formation reaction, and affects the tritiu m release behavior. Additionally,
the release of water vapor possibly promotes the mass loss of Li fro m the ceramic breeder materials. In this study,
we investigated the changes in the mass of Li2 TiO3 (NFI, CEA) and Li2 TiO3 with additional Li (JA EA) at elevatedtemperatures. In hydrogen atmosphere (10,000 pp m H2 /Ar), the Li mass loss is close to be related to the amount ofadsorbed water and amount of water formation in each Li2 TiO3 , and the Li mass loss of Li2 TiO3 (CEA) is about
three times larger than that of the other Li2 TiO3 pebbles. It is considered that the chemical fo rm of released Li
different whether there is the adsorbed water and the water vapor by the water format ion reaction in the gas phaseor not.
Keywords: ceramic breeders, Li mass loss, surface water, water formation reaction, hydrogen atmosphere

1. Introduction

The tritiu m release behavior fro m the ceramic b reeder
materials packed in the b lanket module under ITER-like
operational condition is estimated based on the tritium
release model using parameters obtained in our studies[2]. It is assumed in this estimation that purge gas is He
gas mixed with 100Pa H2 and that physisorbed water isremoved completely before operation. The operational
condition of ITER is repetition of a burning t ime o f 400
seconds with dwell t ime of 1400 seconds. This model
simp lifies the condition of the ITER test blanket modulesuggested by Japan Atomic Energy Agency (Fig. 2). Fig.
3 shows the tritiu m release behavior fro m Li2 TiO3 (NFI)
under the operational condition of ITER. It was estimatedthat the partial pressure of water vapor in gas phase wasabout 20Pa and it is kept for about 100hour. This water
vapor affects the chemical form of the released trit iu m.
The reaction of released water vapor with the ceramic
breeder materials may pro mote the Li mass transfer.
When the ceramic breeder materials are p laced at h igh
temperature conditions, a certain amount of Li may
evaporate. If non-neglig ible Li transfer is caused, itwould possibly lead to serious issues such as stuffing anderosion of pipes, and drop o f TBR. However, the Li masstransfer in a packed bed of the ceramic breeder materials
has not been sufficiently understood to date. In thepresent study, we investigated the changes in the Li massloss of Li2 TiO3 (NFI, CEA), and Li2 TiO3 with additional
Li (JAEA ) at elevated temperatures in hydrogen
atmosphere, and compared them.

The present authors have developed an estimation model
to represent the tritiu m release behavior fro m ceramic
breeder materials, wh ich is valid for out-of-pile andin-pile conditions [1, 2]. Tritiu m release curves estimatedby the tritiu m release model gave good agreements withexperimental curves fro m the ceramic breeder materials
under various purge gas conditions [3]. Fig. 1 shows thetritiu m release model constructed by the present authors.
In the model, the trit iu m d iffusion in a breeder g rain,
tritiu m transfer to surface water at interfacial layer, andsurface reactions on grain surface as adsorption /desorption of physisorbed and chemisorbed water,
isotope exchange reaction with hydrogen in purge gas(isotope exchange reaction 1), isotope exchange reaction
with water vapor in purge gas (isotope exchange reaction2), and water formation react ion of hydrogen in purge gas
with oxygen in the material are taken into account.

Fig. 1 The tritium release model from ceramic breeder materials
kashimura@aees.kyushu-u.ac.jp

18©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

H. Kashimura et al., Mass Loss of Li2TiO3 Pebbles in Atmosphere Containing Hydrogen

Fig. 2

Fig. 3Fig. 4The ITER-TBM type model

The simulation result of tritium and water release
curve under the operational condition of ITER

The SEM picture of as-received Li2TiO 3 (NFI)

Fig. 5 The SEM picture of as-received Li2 TiO3 (CEA)
was held. The water vapor concentration at outlet wasmeasured by a hygro meter (MAH-50, SHIMAZU Co.).
The weight changes of sample pebbles were measured
before and after the experiment. It is considered that theweight loss of sample pebbles is caused by desorption ofadsorbed water vapor, disassociation of oxygen withwater formation reaction and release of Li-containing
species. In this study, weights of released water vapor andoxygen are subtracted fro m the weight loss of sample
pebbles and the rest is defined as Li mass loss.

2. Experimental
In this experiment, Li2 TiO3 pebbles produced by Nuclear
Fuel Industry (NFI), Co mmissariat à l'énerg ie atomique(CEA ), and Li2 TiO3 with addit ional Li, 2.06 in Li/Ti ratio,
produced by Japan Atomic Energy Agency (JAEA) wereused. The Li2 TiO3 with additional Li is in developmental
stage in JA EA as an advanced tritiu m breeder material for
DEMO blanket. The samp le pebbles of 0.5 g were packed
in a quart z tube, 4 mm in inner d iameter, and 36 mm longfor Li2 TiO3 (NFI) and 29 mm for Li2 TiO3 (CEA ), and29mm for Li2 TiO3 with additional Li (JA EA) in bed
length. Fig. 4, 5 shows a SEM p icture of the inside of
Li2 TiO3 (NFI) and Li2 TiO3 (CEA ) pebble. It wasestimated fro m SEM observation that the grain diameterof Li2 TiO3 is 5μm for NFI, 1μm for CEA, and 30μm for
JAEA. Each pebble diameter is 2mm, 1mm, and 1mm,
respectively. The details of the experimental procedures
were reported in the prev ious paper [4]. After thedesorption of physisorbed water from the sample pebbles
by introduction of H2 /Ar (1000PaH2 /Ar) gas at room
temperature, the sample bed was heated to 900 °C, or
600 °C with a ramp rate of 5 °C/min, and the temperature

3. Result and Discussion
Fig. 6 shows the concentration of water vapor at outlet ofthe quartz tube when Li2 TiO3 (NFI) sample bed was heatedwith a ramping rate of 5 °C/ min to 900 °C. It is observedthat when H2 /Ar gas is introduced, physisorbed water canbe released at room temperature and chemisorbed wateris almost removed by elevating temperature to about
400 °C. At over 500 °C, the water vapor is released by thewater formation reaction between H2 and oxygen in thematerial. The weight of chemisorbed water and o xygen
are obtained to be 6.52×10-3 g  and 1.55×10-3 g  on
average.

19

H. Kashimura et al., Mass Loss of Li2TiO3 Pebbles in Atmosphere Containing Hydrogen

but we use Li2 TiO3 (CEA) which surface area is smaller
than that of Li2 TiO3 (NFI) in this experiment, the Li massloss of Li2 TiO3 (CEA) is promoted. Fig. 8 shows therelationship between the amount of released water and the
Li mass loss. In this figure, the Li mass loss of Li2 TiO3
(CEA) is three times larger than that of the others , and theamount of released water of Li2 TiO3 (CEA) is also about
three times larger. It is suggested that the Li mass loss is
closely associated with the amount of released water fromthe sample pebbles. The water formation reaction is
promoted at high temperatures in hydrogen atmosphere.
Therefore it is predicted that the amount of Li mass loss inhydrogen atmosphere is larger than that in dry gas
condition, but it is not carried out yet.
Fig. 6

Fig. 7

Example of the water release curve from Li2TiO3 (NFI)

Result of the experiment to the weight changes ofsample and mass loss of Li at 900 °C

Fig. 8Fig. 7 shows the weight changes of sample pebbles andthe Li mass loss. In the case of Li2 TiO3 (NFI), the weightloss of sample pebbles is about 1.0 wt% and the Li massloss is about 0.7wt% after 50hours. The weight of sample
seems to be almost steady after 24hours. For the data of
Li2 TiO3 (CEA), a huge mass loss is observed after 8hours.
After 50hours, the weight loss of samp le pebbles is about
4.0 wt% and the Li mass loss is about 3.0 wt%. In thecase of Li2 TiO3 with additional Li, the weight loss ofsample pebbles is 2.0 wt%, and the Li mass loss is about
1.0 wt% after 50hours. It seems that the Li mass loss still
continues.
The Li mass loss per gram of Li2 TiO3 (CEA) is about
three times larger than those of Li2 TiO3 (NFI) and Li2 TiO3
with additional Li. It is considered that the amount of Li
mass loss is related to the difference of surface area and theamount of released water. It is found by the present authorsthat the amount of chemisorbed water per surface area of
Li2 TiO3 (NFI) are about twice larger than that of Li2 TiO3
(CEA) and, the amount of generated water per surface areaof Li2 TiO3 (CEA) and that of Li2 TiO3 (NFI) are almost
same. Therefore it is considered that the Li mass loss of
Li2 TiO3 (NFI) is promoted by the release of water vapor,

Fig. 9Comparison between the amount of released water andthe amount of the Li mass loss

Result of the experiment to the weight changes ofsample pebbles and mass loss of Li at 600 °C

Fig. 9 shows the weight changes of sample pebbles when
Li2 TiO3 with additional Li were heated to 600 °C. In thisfigure, the weight loss of the sample pebbles is 1.0 wt%,
and the Li mass loss is 0.5 wt% after 50hours. It is about
one-half of the Li mass loss at 900 °C. In the previous
studies, it has been found that the amount of generated
water by the water format ion reaction at 600 °C is smaller

20

H. Kashimura et al., Mass Loss of Li2TiO3 Pebbles in Atmosphere Containing Hydrogenthan that at 900 °C for Li2 TiO3 [5,10]. Therefore it can besaid that the Li mass loss at 600 °C is s maller than that at
900 °C because the release of Li by the reaction withgenerated water at 600 °C is smaller. The weight ofsample pebbles seems to be almost steady after 40 hours.
We are also considering that the amount of released
water is associated with the method of Li2 TiO3
manufacturing. By the present authors, Li2 TiO3 (NFI) is
made by the wet process, and Li2 TiO3 (CEA) is made by
the extrusion-spheronization-sintering process. [6, 7] It isconsidered that the amount of released water is related towhether water is used or not in manufacturing process.

It is obtained that the Li mass loss of Li2 TiO3 withadditional Li (JAEA) at 600 °C is half t imes larger thanthat of Li2 TiO3 with additional Li at 900 °C. It is supposedthat the rate of water formation velocity is lo w at 600 °C,
and the amount of released water is less than that of
900 °C.
In hydrogen atmosphere, the chemical form of released Li
is different whether there is the adsorbed water and thewater vapor by the water formation reaction in the gas
phase or not. If there is the absorbed water and the water
vapor by the water formation reaction in the gas phase, thechemical form of released Li is LiOH, and there is nowater vapor in the gas phase, Li could evaporate from thesurface of Li2 TiO3 .

References

Fig. 10[1] T. Kinjyo, et al., Fusion Eng. Des. 83, (2008) 580 587.
[2] T. Kinjyo, et al., J. Nucl. Mater. 367 - 370, (2007)
1361 - 1365.
[3] T. Kinjyo, et al., Fusion Eng. Des. 82, (2007) 15 - 24.
[4] M. Nishikawa et al., J. Nucl. Mater., 325, (2004) 87- 93.
[5] T. Hanada, et al., J. Nucl. Mater., 417, (2011) 735-738.
[6] K. Tsuchiya, et al., J. Nucl. Mater., 258-263, Part 2,(1998) 1985 - 1990.
[7] J. D. Lulewicz and N. Rou x, J. Nucl. Mater. 307-311,
Part 1, (2002) 803-806.
[8] T. Hoshino, et al., Fusion Eng. Des. 81, (2006)
555-559.
[9] T. Kin jyo, et al., Fusion Eng. Des. 83, (2008)
580-587.
[10] K. Katayama, et al., Fusion Eng. Des. 87, (2012)927-931.

Picture of the quartz tube after the experimentFig. 10 shows the picture of quartz tubes when Li2 TiO3
(NFI) was used for sample bed after experiment. In thispicture, the wall of the quartz tube after experiment isstained with the Li compounds. White turbidity is formed
around the sample bed and the downstream region where
the temperature is lower than that at the sample bed. It is
speculated that Li reacts with the water vapor on thesurface of the breeder material and vaporized as LiOH, andthen deposits at lower temperature region. White turbidity
around the sample bed may be caused by just Li
evaporation. It has been observed by Hoshino et al. that Li
partial pressure is higher than LiOD partial pressure at hightemperature under D2 atmosphere [8]. When the sample
bed is held at high temperature for a long time andadsorbed water and the generated water by the water
formation reaction expended completely, and Li would
evaporate as Li atoms from the surface of the breeder
material slowly.

4. Conclusions
The Li mass loss of Li2 TiO3 in hydrogen atmosphere isclose to be related to the amount of adsorbed water andamount of water formation in each breeder material. The Limass loss of Li2 TiO3 (CEA) is about three times larger thanthat of Li2 TiO3 (NFI) or Li2 TiO3 with additional Li(JAEA).

21

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University
Satoshi FUKADA
Department of Advanced Energy Engineering Science, Kyushu University(Received: 7 May 2012 / Accepted: 6 August 2012)

There is increasing interest in using ionic molten-salt Flibe not only as self-cooled tritium(T)-breeding material in afusion reactor blanket but also as fuel solvent of molten-salt fission reactors. Application of Flibe to T-breeding fluid for astellarator-type fusion reactor operated at a high magnetic field brings large simplification of its blanket structure,
allowing continuous operation under high-beta plasma conditions. Using mixed Flibe-ThF4+UF4 fuel in molten salt
fission reactors permits stable long-term operation without fuel exchange. When Flibe or Flinak is irradiated by neutrons,
however, acid and corrosive TF is generated, and some T permeates through structural walls. In order to solve these
problems, chemical conditions of Flibe are changed using the redox-control reaction, Be+2TF=BeF2+T2. In addition,
permeation of hydrogen isotopes is lowered by enhancing T recovery rates. Part of Flibe-tritium researches are
performed at Idaho National Laboratory (INL) under the Japan-US collaboration work of JUPITER-II. Our own
contributions to the topics are shortly introduced in this paper.
Keywords: Flibe, tritium, fusion reactor, molten salt fission reactor, permeation

1. Introduction

TBR>1 [20], which is required for a fusion blanket. Flinak
also gives a higher TBR value with use of Be [21].
Since Flibe heated to above m.p. can dissolve ThF4 andUF4 in any composition, it is used as fuel solvent in molten
salt fission reactors. 232Th is a fertile material which can beconverted into fissile 233U by a sequence ofneutron-absorption reactions and the two beta decays, i.e.
232
Th + n → 233Th → 233Pa → 233U. The earliest molten salt
reactor was set up in Oak Ridge National Laboratory(ORNL) in the 1960’s. China is recently interested in its
application to fuel solvent for the Thorium Molten Salt
Reactor (TMSR), because of their large possible reserves
and less generation of radioactive waste. TMSR can be
operated for extended time without fuel discharge. The
experimental fission reactor will be constructed in Shanghai
Institute of Applied Physics, China Academy of Engineering
Physics (SINAP, CAEP) in the near future. In this context, aFlibe is a stable mixed molten salt composed of
2LiF+BeF2 in molar ratio (2:1) with melting point of 459oC.
Flinak (LiF+NaF+KF) is also a stable molten salt with
0.465:0.115:0.42 in molar ratio and a similar m.p. of 454oC.
The Flibe and Flinak ionic salts are planned to be used not
only for a tritium(T)-breeding blanket in fusion reactors
such as FFHR-2 at National Institute for Fusion Sciences(NIFS) in Japan [11,13,15,17] but also as fuel solvent insmall molten salt fission reactors such as FUJI [21]. The
former is intensively investigated under the direction of Prof.
Sagara of NIFS [11,13,15,17]. Flibe is chemically andphysically stable even at high temperature. Its low electric
conductivity makes more suitable when utilizing instellarator-type fusion reactors, which can be operated underconditions of a high beta magnetic field. Flibe can achieve
the T self-sufficient condition of tritium breeding ratio

Fig. 1 Photos of solid Flibe at room temperature and transparent molten salt when its melts

author’s e-mail: sfukada@nucl.kyushu-u.ac.jp22©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

S. Fukada, Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University

2. Redox Control

collaboration work on tritium release and measurement in
the TMSR was recently initiated between researchers of
SINAP, CAEP and our Kyushu University.
Generally, molten ionic salts of Flibe and Flinak arestable even under air at high-temperature. They become
transparent when they are melted as shown in Fig. 1. When
Flibe or Flinak is used for fusion or fission reactors under ahigh-intensity neutron flux or high-temperature conditions,various advantages and disadvantages are observed. The
advantages of Flibe or Flinak can be summarized asfollows: (i) comparatively low reactivity with O2 or H2O, (ii)low electric conductivity, (iii) low T solubility, and (iv) easy
T recovery. On the other hand, the disadvantages whenusing Flibe or Flinak are (i) generation of chemicallyreactive TF under neutron irradiation, (ii) comparatively
high viscosity (8.6 mPas at 600oC of Flibe compared with
1.0 mPas at 20oC of H2O), which leads to high pumping
power, (iii) the experimental increase of viscosity withincrease of BeF2 content, and (iv) high T permeability
through some metals.
Many experiments have been carried out under theJapan-US collaboration work called JUPITER-II in
2001-2006 and our contributions within the frame of thedomestic collaboration have continued since then. Part of
the tritium research was performed in the Safety and Tritium
Applied Research (STAR) facility of Idaho National
Laboratory (INL). Many achievements are summarized in
the references of the present paper [1-18]. Our T-related
experimental or design-based research on applications of themolten salts Flibe and Flinak to fusion or fission reactors isbriefly introduced in the present paper. The introduction ismainly devoted to our own studies due to the limit of paper
length.

When mixed molten salt of Flibe or Flinak is exposed
to neutron flux in fission or fusion reactors, tritium fluoride(TF) is generated by the reaction of LiF with neutrons asfollows:
6
LiF + n = TF + 4He
7
LiF + n = TF + 4He + n’
The reaction cross-section of 6Li with thermal neutrons
is very large. On the other hand, that with 7Li is small andhas the threshold energy of 2.47 MeV. Therefore,
enrichment of 6Li is often necessary to achieve high TBR,
when applied in a fusion reactor blanket. Since the reaction
product TF is a strong acid, it corrodes metallic walls.
Consequently, TF is converted into molecular T2 or HT,
which is chemically compatible with metallic walls and thefluorine is allowed to react with metallic Be. The
reduction/oxidation (redox) control reaction in the liquid
phase of Be (dissolved in Flibe) + 2TF (HF) → BeF2 + T2
(H2) is used for the purpose. The experimental results are
summarized in Fig. 2. It is assumed in the calculation thatthe Be dissolution rate is constant, the above reaction issecond-order in HF concentration and contributions of thereverse reaction are negligibly small. As seen in the figure,close agreement is obtained between experiment andcalculation regardless of the HF concentration and the Be
dissolution concentration. A reaction rate constant isobtained from fitting between experimental results andcalculations as shown in Fig. 2.
Since the solubility of T2 generated by the reaction in
Flibe is very low, T easily permeates through structural
materials of fusion or fission reactors. Therefore, highly
efficient procedures for T recovery from the Flibe flow infission or fusion reactors and an effective T permeation
protection are necessary in order to operate safely fusion or

Fig. 2 Comparison between experiment and calculation of reaction of 2HF+Be→BeF2+H2 in liquid Flibe

23

S. Fukada, Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University

fission reactors. Our previous work on T recovery andpermeation through Flibe covers the following five topics,(i) diffusivity and solubility of hydrogen isotopes in Flibe or
Flinak, (ii) Flibe purification and redox control by Be, (iii)chemical analysis of T in Flibe, (iv) design of equipment torecover T from Flibe, (v) design of Flibe blanket, etc. Manyof these results are described in the references [1-18]. In thenext section we focus on the solubility and permeability ofhydrogen isotopes in Flibe or Flinak.

3. Solubility and Permeability of H Isotopes in Flibeor Flinak

Fusion or fission communities are much interested in
two liquid breeders Flibe and Flinak, and many of theirvarious properties have been investigated experimentally.
Solubility and permeability of H isotopes in Flibe or Flinak
molten salts are very important factors for the design offusion reactors and to estimate the T-breeding performance.
Solubility is related to the T inventory and permeability to T
leakage through boundaries. Fig. 3 correlates solubility
constants for various gases in Flibe and Flinak. Dissolutionof gases in Flibe or Flinak follows Henry’s law described byc k = K H ,k pk , where ck and pk are the concentration of gas kdissolved in the molten salt and the partial pressure of gas k,respectively. KH,k (k=H2, He, Ar, Xe) is the Henry constant ofgas k. These gases are dissolved in atomic or molecularform in Flibe and Flinak, and the major force acting
between molten salt and dissolved inert gas is surface
tension. Therefore, the solubility constant KH,k is given by
the following equation regardless of Flibe and Flinak:⎛ E + 4 πN r 2γ ⎞
A b,k
⎟⎟
(1),
K H,k = exp⎜⎜− 0
R
T
g
⎝
⎠

Fig. 3 Solubilities of H2, He, Ar, Xe in Flibe or Flinak
The product of diffusivity DH2 and solubility KH,H2 isequal to the permeability when no retardation onmaterial-gas interface occurs. The steady-state H2
permeation rate denoted by jH2 is proportional to the partialpressure difference between the upstream and downstreamsides denoted by pH2,up and pH2,down, respectively, and isexpressed as follows:
K H, H 2 D H 2
j H2 =
p H 2 ,up − p H 2 ,down(2),δ
where δ is the thickness of the molten salt layer and DH2 isthe diffusivity of H2 in Flibe or Flinak.

()

4. Recovery of Tritium from Flibe or Flinak

The equilibrium partial pressure of T dissolved in Flibeor Flinak and vapor pressure of Flibe are shown in Fig. 4
together with other liquid blanket candidates. As seen there,
the T equilibrium pressure over Flibe is the highest, being
about 105 Pa for 1ppm T dissolved in Flibe. Therefore,
simple He gas purge is promising to recover T in Flibe if
direct contact between gas and liquid is possible. If directcontact is unsuitable, the problem is avoided by means of awhere rb,k is the atomic or molecular radius of the gaseous
species k, NA the Avogadro number, Rg the gas law constant,
T temperature and γ the surface tension of Flibe or Flinak.
Although E0 is a constant for inert gases, in case of H2
another attractive force between this molecule and themolten salt plays a role. Therefore the solubility of H2 showsdifferent tendency from that of inert gases (see Fig. 3).

Fig. 4 Comparison of T equilibrium pressure and vapor pressure among Li, Li17Pb83 and Flibe

24

S. Fukada, Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University

permeation window. If the redox control is not successful,the chemical form of tritium remains to be TF. The
equilibrium pressure of TF is, however, about 1/200 of T2.
Therefore the efficiency of gas purge is reduced.
Fig. 5 shows a schematic view of a He-Flibecounter-current extraction tower. Tritium dissolved in Flibeat the tower inlet diffuses from the Flibe bulk to the
Flibe-He interface. When it arrives at the surfaces, T desorbs
to the He side. Then T2 gas diffuses through the He-sideboundary layer. T2 in He is purged out to the gas outlet. In
order to keep good contact between He and Flibe, inert
metal packing is placed inside of the tower. Rashig ring,
Dixon ring or equivalents can be used as usually packed incommercial gas-liquid extraction towers such as distillation
process. Analytical study on T recovery from Flibe using thecounter-current tower is being performed in Japan and
China, but no experimental result has been presented.
Here, we compare Li chemical activity in the liquidblanket candidates among Li, Flibe, Li17Pb83 and Li25Sn75
eutectic alloys. Li has the highest vapor pressure among
them. Since the chemical activity of Li is high, its use infusion reactors may be limited for safety reasons eventhough its TBR is the largest among the four candidates. The
vapor pressure of Li17Pb83 is lower than that of Li. The
chemical activity of Li in the Li17Pb83 eutectic alloy
decreases with increasing temperature as one can see whencomparing the difference in vapor pressure between Li and
Li17Pb83, which slightly spreads with increasing temperature.
Judging from the viewpoint of vapor pressure, Flibe and
Li25Sn75 are the lower group. The chemical activity of Li islow. Since the vapor pressure of Flibe is about 10-2 Pa at 800
K, its vaporization is negligibly small even when its
accidental release happens (see Fig. 4). Therefore, Flibe is
the safest materials among the four candidates. Li25Sn75
cannot constitute any blanket design of TBR>1.

Flibe in

Flibe out
Fig. 5 Overview of He-Flibe counter-current extraction
tower
When a Flibe molten salt mixed with UF4 and TF4 isused as fuel of a molten salt fission reactor, the largest
disadvantage is the strong gamma ray originated from 232U
or its daughter nuclei 208Tl. The gamma ray energy from208
Tl is 2.8MeV. Therefore, the designed molten salt reactor
should be buried under the ground in order to decrease
non-necessary exposure to gamma rays (see Fig. 6). The
figure shows an example for the molten salt reactor cited in
the reference [19]. One generation path of 232U from 232Th
fuel is considered to be 232Th(n,2n)231Th, 231Th→231Pa,
231
Pa(n,γ)232Pa, and finally 232Pa→232U. The first (n,2n)reaction occurs only with fast neutrons. The thresholdenergy is 6.5MeV. Therefore, neutron slow-down using aresonable moderator is effective in order to decrease thegeneration of 232U. Another path for the generation of 231Th
is the neutron capture reaction of 230Th. This isotope ispresent as a natural impurity in 232Th. 230Th is a daughter

5. Thorium as Fertile Material for Molten Salt Reactor

Flibe out

Fig. 6 Conceptual design of underground molten salt fission reactor. Figure is cited from the address [19].

25S. Fukada, Flibe-Tritium Research for Fission or Fusion Reactors at Kyushu University

product of the alfa decay of 238U. When 231Th is generated,as described above, the same path works to generate 208Tl.
Therefore, it is effective to separate natural U from naturalTh (almost 100% 232Th) beforehand in order to minimize thegeneration of high intensity gamma ray. The chemical
separation between U and Th will be performed by a solvent
extraction technique using Tributyl phosphate (TBP) as
solvent.

tritium leak from fusion reactor system”, Fusion Science andTechnology, 48 (2005) 666-669.
[8] S. Fukada, A. Morisaki, “Hydrogen permeability through amixed molten salt of LiF, NaF and KF (Flinak) as aheat-transfer fluid”, Journal of Nuclear Materials, 358 (2006)
235-242.
[9] Y. Edao, S. Fukada, H. Noguchi, A. Sagara, “Tritium releasefrom neutron-irradiated Flibe purged out by Ar-H2 or Ar atelevated temperature”, Fusion Science and Technology, 55(2009) 140-151.
(iii) Design study of tritium recovery from Flibe blankets:[10] S. Fukada, M. Nishikawa, A. Sagara, “Calculation of recoveryof tritium from a Flibe blanket”, Fusion Technology, 39(2001) 1073-1077.
[11] A. Sagara, H. Yamanishi, T. Uda, et al., “Studies on Flibeblanket designs in helical reactor FFHR”, Fusion Technology,
39 (2001) 753-757.
[12] S. Fukada, M. Nishikawa, A. Sagara, T. Terai, “Mass-transport
properties to estimate rates of tritium recovery from Flibeblanket”, Fusion Science and Technology, 41 (2002)
1054-1058.
[13] A. Sagara, T. Tanaka, T. Muroga, et al., “Innovative liquidblanket design activities in Japan”, Fusion Science andTechnology, 48 (2005) 524-529.
[14] S. Fukada, A. Morisaki, A. Sagara, T. Terai, “Control of tritium
in FFHR-2 self-cooled Flibe blanket”, Fusion Engineeringand Design, 81 (2006) 477-483.
[15] A. Sagara, O. Mitarai, S. Imagawa, et al., “Conceptual design
activities and key issues on LHD-type reactor FFHR“, Fusion
Engineering and Design, 81 (2006) 2703-2719.
[16] S. Fukada, K. Katayama, T. Terai, A. Sagara, “Recovery oftritium from Flibe blanket in fusion reactor“, Fusion Scienceand Technology, 52 (2007) 677-681.
[17] A. Sagara, O. Mitarai, T. Tanaka, et al., “Optimization
activities on design studies of LHD-type reactor FFHR”,
Fusion Engineering and Design, 83 (2008) 1690-1695.
[18] S. Fukada, Y. Edao, A. Sagara, “Effects of simultaneoustransfer of heat and tritium through Li-Pb or Flibe blanket”,
Fusion Engineering and Design, 85 (2010) 1314-1319.
[19] http://energyfromthorium.com/2010/12/07/not-so-fast-withthorium/moirteller/[20] H. Yamanishi et aI., “Design studies on nuclear properties on
the Flibe blanket for helical-type fusion reactor FFHR”,
Fusion Engineering and Design, 41 (1998) 583-588.
[21] H.Yamanishi, A.Sagara, Annual Report of National Institutefor Fusion Science, April 1999–March 2000 (2000) 177.

6. Conclusions

Applications of molten salt Flibe or Flinak to fusion or
fission reactors and our previous research on these topics
were presented. Since Flibe and Flinak show not only
proper T breeding ability for fusion reactors but also lessreactivity with O2 or H2O regardless of slightly larger
viscosity than other breeding materials, they can be used asa blanket fluid of a FFHR-2 fusion reactor. The advantagesof molten salts were discussed here. On the other hand,corrosive TF acid is generated when exposed to neutrons. In
order to solve this problem, the Be redox control reaction
was shown to be effective. Since the solubility of T2 (H2) iscomparatively small, although it is larger than that of inert
gases, the T2 recovery is readily carried out with a He-Flibecounter-current extraction tower. When Flibe mixed withfertile Th is exposed to neutrons, strong gamma rays are
emitted from the Th fuel. The problem of strong gamma
radiation is mitigated by placing the whole Th fission
reactor under the ground. Impurity removal in natural Th
can decrease the strong gamma ray originated from 208Tl.

7. References(i) Redox control and Flibe purification study:[1] S. Fukada, R.A. Anderl, Y. Hatano, et al., “Initial studies oftritium behavior in Flibe and Flibe-facing material”, Fusion
Engineering and Design, 61-62 (2002) 783-788.
[2] R.A. Anderl, S. Fukada, G.R. Smolik, et al., “Deuterium/tritium
behavior in Flibe and Flibe-facing materials”, Journal of
Nuclear Materials, 329-333 (2004) 1327-1331.
[3] M.F. Simpson, G.R. Smolik, J.P. Sharpe, et al., “Quantitative
measurement of beryllium-controlled redox of hydrogen
fluoride in molten salt”, Fusion Engineering and Design, 81(2006) 541-547.
[4] D.A. Petti, G.R. Smolik, M.F. Simpson, et al., “JUPITER-II
molten salt Flibe research: An update on tritium mobilization
and redox chemistry experiments”, Fusion Engineering and
Design, 81 (2006) 1439-1449.
[5] S. Fukada, M.F. Simpson, R.A. Anderl, et al., “Reaction rate ofberyllium with fluorine ion for Flibe redox control”, Journalof Nuclear Materials, 367-370 (2007) 1190-1196.
(ii) Solubility, diffusivity, and permeability of hydrogen isotopes
through Flibe and Flinak:[6] S. Fukada, R.A. Anderl, R.J. Pawelko, et al., “Flibe-D2
permeation experiment and analysis“, Fusion Science andTechnology, 44 (2003) 410-414.
[7] S. Fukada, R.A. Anderl, T. Terai, et al., “Diffusion coefficient oftritium through a molten salt Flibe blanket and evaluation of26

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Experimental and computational studies on tritium permeation mechanism
in erbium oxide
Wei Mao1),*, Takumi Chikada1), Akihiro Suzuki2), Takayuki Terai1) and Kenji Yamaguchi3)1) Department of Nuclear Engineering and Management, School of Engineering, The University of Tokyo, 7-3-1 Hongo,
Bunkyo-ku, Tokyo 113-8656, Japan.
2) Nuclear Professional School, School of Engineering, The University of Tokyo, 2-22, Shirakata-shirane, Tokai, Naka,
Ibaraki 319-1188, Japan.
3) Quantum Beam Science Directorate, Japan Atomic Energy Agency, 2-4 Shirakata-shirane, Tokai, Naka, Ibaraki
319-1195, Japan.
(Received: 10 May 2012 / Accepted: 3 October 2012)

A tritium permeation barrier (TPB) is much required in fusion blankets to reduce the loss of fuel and minimize
radiological hazard. However, the detailed mechanism of tritium permeation through TPB coatings has not yet beenclarified because of their complicated crystal structures. To understand the microscopic mechanism, we have not only
prepared and characterized nanostructured ceramic Er2O3 thin films, but also studied the energetics and mobility ofinterstitial deuterium in cubic bulk Er2O3 using ab initio density-functional calculations. The estimated diffusion activationenergy (Ea) of interstitial deuterium is somewhat higher than the diffusion energy barrier observed experimentally at 873 K.
This result shows that grain boundary structures existing in the TPB of Er2O3 significantly lower the diffusion barrier andchange the diffusion pathway. Transport of hydrogen and its isotopes through the Er 2O3 coatings are likely to be dominatedby grain boundaries rather than by grains.
 
Keywords: Hydrogen isotope permeation measurement, Erbium oxide, Density functional theory calculations,
Nudged elastic band

ceramics [6]. Recently, Er2O3 is attracting more and more
attentions in the application of TPB coatings for fusion
systems.
Up to now, we have made rapid progress in the
fabrication of TPB coatings by different methods such as
sputtering deposition (SD) [7] and metal-organic
decomposition (MOD) [8], but complex crystal structuresand many surface defects still exist in the coatings, whichmake it difficult to investigate the detailed mechanism ofhydrogen isotope permeation. At present, a high-quality
ceramic Er2O3 thin film with nanostructures is
successfully synthesized by arc source plasma assisted
deposition (ASPAD). The ceramic shows grain boundary
structures which are regarded as large defects in terms of
permeation of hydrogen isotopes. Consequently, we place
more and more emphasis on investigating hydrogenisotope behavior in this nanostructured thin film becauseof its slightly simpler structures compared to those
prepared by SD and MOD methods.
In addition to experimental studies, ab-initio
calculations based on density functional theory (DFT)
have been applied to complement the established

1. Introduction

There is no doubt that the environmental andeconomic acceptability of presently conceived D-T
fueled fusion power plants will be dependent on theability to contain and handle tritium within the reactor
building and to control tritium releases to theenvironment without incurring into exorbitant costs [1, 2].
The problems of tritium leakage have been studied since1970`s [3], and tritium permeation barrier (TPB) coating
has been one of the most important subjects in thedevelopment of liquid lithium blanket for fusion systems.
As to coating materials, a great many ceramics such as
Y3Al2O12, CaO, MgO, AlN (aluminum nitride), Al2O3,
Y2O3 and Er2O3 were initially considered potentiallyeligible. Of these Er2O3 was selected as one of the most
preferable candidate materials for TPB coatings becauseof its high permeation reduction factor (PRF) [4], highthermodynamic stability at high temperatures and in air,
good compatibility with liquid Li, and high electricalresistivity [5]. Additionally, Er2O3 has one of the lowest
Gibbs free energy of formation among all binary oxide
author’s e-mail: mao@nuclear.jp

27©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

W. Mao et al., Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide> direction. That is, it possesses two vacancies (16c) at
the corners of the oxygen cube, as shown in Fig. 1.
Ab initio density-functional calculations are used toinvestigate the diffusion behavior of interstitial deuteriumin bulk Er2O3. We employ Vienna ab initio Simulation
Package (VASP) for our total energy and DFT
calculations. All calculations were performed using thefull-potential
frozen-core
all-electron
projected
augmented wave (PAW) method within the Generalized
Gradient Approximation (GGA) of Perdew and Wang
(PW91), as implemented in VASP [12]. The cut-offenergy of plane wave was set to 500.0 eV. According to
the Monkhorst-Pack scheme [13], we performed the tests
with the k-point grid of 2x2x2 in the integration over the
Brillouin zone. In the initial configuration of the crystal
lattice, we have chosen the experimental cubic structureof Er2O3 unit cell with a = 1.0543 nm, 32 erbium atoms(occupying the 8a and 24d equipoints) and 48 oxygen
atoms (occupying 48e equipoints). We conducted abinitio DFT calculations in the unit cell of Er2O3 for thegeometry optimization of lattice parameters and atom
positions. The geometry relaxations were terminatedwhen residual forces on each atom and energy changes
on each atom were less than 0.001 eV/Å and 1x10-5 eV,
respectively. All calculations below referring tointerstitial D were performed by inserting an individual D
into the cell optimized by DFT methods. It is worth
noting that the parameters of cut-off energies and k-point
meshes were tested under the convergence criteria ofenergy changes on each atom less than 0.001 eV, andfound sufficient to give reliable convergent results in the
ab initio DFT calculations.

experimental methods to investigate the energetics andmobility of atomic deuterium (D) in cubic bulk Er2O3. In
recent years, DFT calculations have become a very
valuable tool to elucidate the structures and determine thedynamics of interstitial H in metals, alloys, ceramics andother solids such as heterogeneous catalysts [9]. It isreported that DFT can be used to predict the precise
mechanism of H diffusion or reaction events incondensed phases [10]. It is, therefore, essential to apply
DFT calculations to the diffusion of hydrogen and its
isotopes in Er2O3 for the verification of experimental data.
As a result, a comparison between diffusion experiments
and DFT simulations can provide useful information on
the mechanism of hydrogen isotope permeation.
In this study, the objective is to compare thediffusion activation energies of interstitial D derived by
DFT calculations with the diffusion energy barriers
obtained from the experiments, and to study the
relationship between the microstructure and hydrogenisotope behavior in Er2O3 for the mechanism of tritium
permeation.

2. Methodology
2.1. Preparation and characterization of Er2O3 coatings
A mirror polished disk-shaped substrate of 20 mm
in diameter and 0.5 mm in thickness was coated by arc
source plasma assisted deposition (ASPAD), described inRef. [7]. The substrate materials were made from reduced
activation ferritic/martensitic (RAFM) steels, including
JLF-1 ( 0.10% C; 0.05% Si; 0.45% Mn; 0.003% P;
0.002% S; 0.003% Al; 8.85% Cr; 1.99% W; 0.20% V;
0.080% Ta; 0.0231% N; 0.0002% B; < 0.01% Ni; <0.05% Cu; <0.001% Mo; < 0.002% Nb; balance Fe) andF82H (7.65% Cr; 2.00% W; 0.50% Mn; 0.20% V; 0.04%
Ta; 0.10% Si; 0.10% C; balance Fe) steels. Er2O3 coating
about 1.3ȣm thick was deposited on one side of thesubstrate. Subsequently, deuterium permeation throughthe coated sample was carried out in the temperature
range 573-973 K at driving pressures of 103-105 Pa [7].
2.2 Ab initio density-functional calculations in bulk
Er2O3
At room temperature, Er2O3 exhibits cubic rareearth bixbyite (C phase, space group Ia3) structure with alarge complex unit cell comprising 16 Er2O3 sequioxide
units with an experimental lattice parameter of 1.0543
nm [11]. In this structure, the O anions approximately
occupy three-quarters of all the tetrahedral interstices
while the Er cations lie in a face centred cubic
arrangement. Apart from rotation symmetry, the bixbyite
structure is different from that of the fluorite in missing
one-fourth of the sites in the anion sublattice in the < 111

Fig. 1. Cubic bixbyite structure of Er2O3 unit cell. The cell is
subdivided into eight small units, in which tetrahedral
vacancies (16c) are virtually depicted as obtained from acomparison between the bixbyite and fluorite structures.

3. Results and discussions

28

W. Mao et al., Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide

3.1 Microstructures of Er2O3 coatings
Er2O3 coatings of 1.3 ȣ m in thickness wereprepared using JLF-1 substrate and the ASPAD method.
The coatings were found to possess a high-quality grainboundary nanostructure, as shown in Fig. 2. This grainboundary is regarded as the largest defect because ofalmost complete absence of other defects such asvacancies and dislocations. The average grain size of thecoatings is about 200 nm.

3.2 Calculations of Er2O3 lattice constant
As initial configuration we have selected theexperimental cubic structure of Er2O3 with a lattice
constant of 1.0543 nm. The lattice parameter corresponds
to a volume of 1.1719 nm3. After relaxations of thestructure with a different lattice constant approaching theexperimental one (forces on atoms are less than 0.001
eV/Å), we obtained the cubic unit cell with a series ofvalues of lattice constant and corresponding volume as afunction of total energy shown in Fig. 4. After fitting with
Birch-Murnaghan 3rd-order equation of state (EOS) [15],we obtained the balanced volume of 1.1727 nm3
corresponding to the lowest total energy and the optimized
lattice constant equal to 1.0545 nm, in good agreementwith the experimental ones [11].

Fig. 2. STEM image of the Er2O3 coating fabricated by arc
sputtering plasma assisted deposition (ASPAD) at room
temperature and annealed at 973 K.

Given no change of grain sizes with instantincrease of temperatures, the relationship between grain
size and deuterium permeation flux was investigated in
the Er2O3 coatings by permeation measurements carriedout at 773, 873, and 973 K, as shown in Fig. 3. The
average grain sizes of the coatings as deposited, after themeasurements at 773, 873, and 973 K are 20, 90, 250,
and 280 nm, respectively. Apparently, the deuterium
permeation flux decreases with increasing grain size
when the temperatures are kept constant, in good
agreement with previous results [14]. In other words, adecrease of grain sizes results in an increase of thepermeation flux of deuterium through the Er2O3 coatings.

Fig. 4. Total energy configurations as a function of volume in theunit cell of Er2O3. The configurations are performed by ab initio
calculations and Birch-Murnaghan 3rd-order EOS fitting.

3.3 Individual D atom in bulk Er2O3
In the lattice model described in Section 2.2, wetheoretically occupied the initial positions of tetrahedral
vacancies in the oxygen anion sublattice, based on acomparison of the fluorite and bixbyite structures. Afterrelaxing the cell consisting of two primitive units, we
introduced an individual neutral hydrogen atom into the
cubic bixbyite structure to simulate dilute loadings ofinterstitial D, and calculate directly its motion by DFT.
Considering a great number of interstitial sites (thetetrahedral vacancies of oxygen anion are included) in the
structure, we found that D binding energy is mainly
dependent on the type of the surrounding atoms, especiallythe nearest neighboring ones. The binding energy of aninterstitial D atom in host lattice can be defined as follows:

Fig. 3. Grain size dependence of deuterium permeation flux
through the 1.3ȣm thick Er2O3 coatings. The grain sizes of thecoatings were estimated by TEM images. The permeation
measurements were carried out at a pressure of 8.00 x 10-4 Pa. In
the figure, X is the permeation flux-grain size point in the plot,
and the red rash arrows represent the change at constanttemperatures of 773, 873, and 973 K.

Ebind

E( D / S )  E( S )  E( D )(1)where E(D/S) and E(S) are the total energy of the system withand without an interstitial D atom, respectively. E(D) is theenergy of D atoms in free space with spin polarization.
With regard to the DFT calculations, the D binding energy

29

W. Mao et al., Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide

in the tetrahedral vacancies of oxygen anion sublatticewere found to be strongest among the interstitial sites,
equal to - 10 kJ/mol representing an exothermic reaction.
This indicates that an individual interstitial deuterium atom
is easily trapped in the tetrahedral vacancies. In otherwords, the tetrahedral interstitial site of oxygen anion
sublattice belongs to the local minima position surroundedby the interstitials filled with O anions.
It is worthwhile to look into the effect of spinpolarization because the total energy goes down 14~18
kJ/mol when a deuterium atom is inserted into theinterstitial site of the tetrahedral. We therefore include spinpolarization in the present computational scheme andapply it to an individual H2 molecule. The calculated
binding energy / cohesive energy and bond length of H2
are 4.55 eV and 0.740 Å, corresponding well to theexperimentally obtained values of 4.48 eV and 0.741 Å
[16], respectively. Such calculations confirm the fact thatthe computational scheme considering spin polarization isreliable and sufficient to give fully converging results.

The overall minimum energy pathway (MEP)
shown in Figure 5 indicates another new local minima
position (C) and barriers (A-C*, C-B*) along the pathconnecting the initial local minima position (A) and its
crystallographic equivalent position (B). The migration
barriers from A to C and from C to B are 41.5 kJ/mol(0.43 eV) and 39.6 kJ/mol (0.41 eV), respectively.
Furthermore, we have as well calculated a migration
barrier of 157.3 kJ/mol (1.63 eV) in the minimum energy
pathway of deuterium diffusion along the direction fromC to C`s crystallographic equivalent site. The maximum(157.3 kJ/mol) in the MEPs corresponds to the position
where D diffuses through the highest number of bondmaking/breaking events in Er2O3, in which thecorresponding activated state locates the O-layer at thesite equidistant from O anions.
3.5 Relationship between diffusion and permeation ofhydrogen and its isotopes through erbium oxide
The permeability of hydrogen isotopes is generally
defined as the steady state diffusional transport of atoms
through a material supporting a pressure difference.
According to the theory reported in the paper of Marchi
et al. [18], the permeability and diffusivity of hydrogenisotopes through single-layer material can be expressedin the general form as follows:

3.4 Diffusion behavior of deuterium in bulk Er2O3
To investigate the diffusion behavior of deuterium inbulk Er2O3, we employed the DFT calculations to examine
trapping binding sites, diffusion pathways and diffusion
activation energies for an individual D atom hopping fromone local minima position to its nearest neighboring site.
The diffusion activation energies (Ea) were determined by
the energy difference between the local minima position
and the activated state site. The hopping distance between
the two energy minima in the bulk Er2O3 was measured tobe 3.787 Å. The activated states were found by the nudged
elastic band (NEB) [17] method using nine images as wellas the symmetry consideration, in which forces on atoms
are zero and any infinitesimal change of deuterium initialcoordinate causes a finite change of atomic coordinates.
The identical position of the activated states lies in the Er
layer at the point equidistant from Er cations, derived by
the above two methods.

E)
)
(2)
RT
ED
)
D D0 exp( (3) 
RT
where Φ (mol/m/s/Pa1/2) and D (m2/s) are thepermeability and diffusivity of hydrogen, respectively. R(8.314 J/mol/K) is the gas constant; T (K) is the sample
temperature; and EΦ (J/mol) and ED (J/mol) are theactivation energy of permeation and diffusion ofhydrogen and its isotopes, respectively. Φ0 and D0 are thecorresponding coefficients of diffusion and permeation.
Equations (2) and (3) indicate that diffusivity canrepresent permeability to a large extent since they are
similar in magnitude over a wide range of temperaturesin the transport of hydrogen and its isotope atoms
through the material.
According to Ref. [14], the permeation flux understeady-state conditions through a single layer barrier canbe expressed as
1
P2
J KD
(4)
d
where J is the permeation flux per unit area of material of
thickness d (m); K is the Sieverts` constant and P thedriving pressure of hydrogen and its isotopes. Since Φ =KD, Eq. (4) in combination with Eq. (2) gives therelationship between the permeation flux (J) and theactivation energy (ED) of diffusion under steady-state
conditions, i.e.
)

Fig. 5. Energy versus reaction coordinate obtained by the NEB
method in bulk Er2O3. The initial separation distance between the
two local minima sites is 3.787 Å.

30) 0 exp( W. Mao et al., Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide

J1)P 2 / d) o exp( E D  ES 12)P / dRTcorrespondingly. Actually, GB’s existing in Er2O3
coatings degrade the grain sizes due to loss of bulk areaand lead to an increase of permeation flux, in good
agreement with the experimental data given in Fig. 3. It
is intuitively considered that diffusion and permeation ofhydrogen and its isotopes through erbium oxide are likely
to occur predominately through grain boundaries rather
than grains. In addition, it is worth noting that theactivation energy obtained in the deuterium permeation isstill large and high concentration to that derived by the
DFT calculations. This suggests that the Er2O3 coatings
prepared by the ASPAD method possess highperformance as required for TPB coatings. As shown in
Fig. 3, the growth of grain size enhances the ability of thecoatings to suppress deuterium diffusion and permeation.
As the grains become very large, the permeability of
deuterium will decrease because of the absence of grainboundaries. That is, the activation energy is maximized
when the Er2O3 coatings become a perfect crystal without
any grain boundary as assumed in the DFT calculations.

(5)
where Es (J/mol) is the heat of solution of hydrogenatoms in the single-layer material or oxide with athickness of d (m), and EΦ is defined as the sum of ES
and ED. It is obvious that diffusion plays an importantrole in the process of permeation of hydrogen and its
isotopes based on the relationship of permeation flux J
and diffusion activation energy ED given in Eq. (5). With
the aim of understanding the mechanism of permeation itis therefore reasonable to study first the diffusion
behavior of hydrogen and its isotopes through erbiumoxide.
In experiments on the permeation of deuterium, thediffusion activation energy was estimated from an
Arrhenius plot of temperature versus diffusivity, D being
obtained by measuring the time-lag tl (s) with relationshiptl (s) = d2/6D, as described in [19]. The measured
activation energy of approximate 20.0 kJ/mol [14] is ofsimilar order of magnitude with the D diffusion energy
barrier of 157.3 kJ/mol evaluated by the NEB method as
implement in ab initio DFT calculations. The
discrepancy in activation energies is partly attributed to
the fact that nanostructured Er2O3 thin films were used in
the experiments while bulk Er2O3 was adopted in thecomputational simulations. Although the Er2O3 thin films
were prepared by the ASPAD method with the aim ofachieving good performance as TPB coatings, they hadvery large defects, i.e. grain boundaries (GB`s) within thenanostructure as revealed by Fig. 1. A GB represents anenvironment where the symmetry of a crystal isterminated. Apart from the diffusion of hydrogen and its
isotopes through the bulk of erbia, diffusion is likely tooccurs also via the large defect GB`s.
In order to elucidate the hydrogen behavior in
erbium oxide, we purposely conducted ab initio DFT
calculations in the bulk Er2O3 assuming an ideal crystal
structure. The volume and geometry relaxations in thecalculations show that an individual interstitial D prefers
to bind in the tetrahedral interstitial sites of O anion
sublattice (local minima). MEP found by the NEB
method shows the diffusion pathway of deuterium
hopping from one local minima to the nearest
neighboring one. Most important, the maximum energy
difference between the minimum site and the estimatedtransition state (saddle point) determined in this way is
157.3 kJ/mol, and somewhat higher than theexperimental value of 20.0 kJ/mol. This indicates that at
the grain boundary the activation energy for hydrogenisotope (deuterium) diffusion decreases because of thedifferent diffusion pathway as predicted by MEP. In otherwords, the diffusivity of deuterium will be increased if
we introduce GB into the perfect bulk Er2O3. On thebasis of Eq. (5), which relates the permeation flux J tothe diffusivity D, the permeability will increase

4. Conclusions

In this study, the mechanism of tritium permeation
through Er2O3 has been investigated by the experimental
measurements and the computational simulations. We
have prepared Er2O3 coatings of 1.3ȣm in thickness by
the ASPAD method and characterized the coatings via
deuterium permeation measurements. From theexperiments, we concluded that a grain boundary
nanostructure exists in the Er2O3 coating and estimatedthe diffusion activation energy to be 20.0 kJ/mol.
Additionally, ab initio DFT calculations based ontransition state theory were carried out for interstitial
deuterium in an ideal supercell assuming cubic bixbyite
Er2O3. Calculated and fitted lattice constants of Er2O3
showed good agreement with the experimental data. Also
the minimum energy sites, diffusion pathways anddiffusion activation barriers for an individual D atom
were found by MEP calculations. The local minima
position is found located in the tetrahedral interstitial siteof oxygen anion sublattice. The results indicate that thepreferred hopping path of individual D atoms is from one
local minima position to its nearest neighboring site. The
maximum diffusion activation energy resulting from the
MEP calculation is 157.3 kJ/mol and higher than thatobtained experimentally (20.0 kJ/mol). This discrepancy is
attributed to grain boundary structures found in the Er2O3
coatings. GB`s lead to pathways of hydrogen diffusion
different from MEP derived by DFT calculations. They
provide useful information on tritium permeation
mechanism in erbium oxide. Additionally, it appears that
diffusion and permeation of hydrogen and its isotopes
through the Er2O3 coatings will probably take placethrough grain boundaries rather than grains. This

31

W. Mao et al., Experimental and Computational Studies on Tritium Permeation Mechanism in Erbium Oxide

conclusion is drawn from i) the comparison of theactivation energy between diffusion experiments and DFT
simulations ii) the interdependence of hydrogen and its
isotope diffusion and permeation, and iii) the relationshipbetween the permeation flux and grain size during themeasurement of deuterium permeation.

Acknowledgements

This work was supported in part by GoNERI, the Global
Center of Excellence (G-COE) Program of the Nuclear
Education and Research Initiative of the Japan Society for
the Promotion of Sciences (JSPS).

References[1] V.A. Maroni, Tritium Processing and Containment
Technology for Fusion Reactors: Perspective and status,
Proc. Second Topical Meeting on the Technology ofControlled Nuclear Fusion, USERDA Report CONF
-760935-P3 (1976) P. 799.
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Muroga, T. Terai, J. Nucl. Mater. 386-388, 700 (2009).
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329-333, 1414 (2004).
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database 2), Version 1.06, The Japan Society ofCalorimetry and Thermal Analysis ( Kagaku Gijutsu Sha,
Japan, 1992).
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Muroga, Fusion Eng. Des. 82, 2572 (2007).
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Fusion Eng. Des. 85, 1537 (2010).
[9] K. Honkala, A. Hellman, I.N. Remedeiakis, Á. Logadóttir,
A. Carlsson, S. Dahl, C.H. Christensen, J.K. Nørskov,
Science 507, 555 (2005).
[10] D.S. Sholl, J. Alloy. Compd. 446-447, 462 (2007).
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Nevyadomskaya, Russ. J. Inorg. Chem. 39, 211 (1994).
[12] G. Kresse, J. Furthmuller, Phys. Rev. B 54, 11169 (1996).
[13] H.J. Monkhorst, J.D. Pack, Phys, Rev. B 13, 5188
(1976).
[14] T. Chikada, A. Suzuki, H. Maier, T. Terai, T. Muroga,
Fusion Sci. Technol. 60, 389 (2011).
[15] F. Birch, Phys. Rev. 71, 809 (1947).
[16] B.G. Johnson, P.M.W. Gill, J.A. Pople, J. Chem. Phys. 98,
5612 (1993).
[17] H. Jónsson, G. Mills, K. W. Jacobsen, Classical andQuantum Dynamics in Condensed Phase Simulations,edited by B. J. Berne, G. Ciccotti and D. F. Coker (World
Scientific, Singapore, 1998), pp. 385-404.
[18] C. San Marchi, B.P. Somerday, S.L. Robinson, Int J
Hydrogen Energy, 26, 597 (2001).
[19] T. Chikada, A. Suzuki, T. Kobayashi, H. Maier, T. Terai,
T. Muroga, J Nucl. Mater. 417, 1241 (2011).

32

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Solubility of hydrogen isotopes in zirconia ceramics
Kenichi HASHIZUME1), Kosuke OGATA1), Satoshi AKAMARU2), and Yuji HATANO2)
1) Interdisciplinary Graduate School of Science and Engineering, Kyushu University, Hakozaki 6-10-1, Fukuoka 812-8581,
Japan.
2) Hydrogen Isotope Research Center, University of Toyama, Gofuku 3190, Toyama 930-8555, Japan.
(Received: 25 May 2012 / Accepted: 5 October 2012)

The hydrogen solubility in zirconia (yttria stabilized zirconia, YSZ) using a tritium imaging plate (IP) technique
has been measured in a temperature region ranging from 673K to 973K. The tritium distribution in the zirconia
specimens was nearly uniform suggesting an equilibrated dissolution. The hydrogen solubility in specimens wererather low (around ppm or below) and showed Arrhenius type temperature dependence, while it seemed to be
independent of Y contents.
Keywords:

Zirconia, YSZ, Tritium Gas Exposure, Imaging Plate Technique, Hydrogen Solubility

the disks exceeded 95 % of theoretical ones for all grades.
The crystal structure of the disk specimens characterizedby X-ray diffraction was tetragonal for TZ-3Y, themixture of tetragonal and cubic for TZ-4Y, and cubic for
TZ- 6Y, 8Y and 10Y as shown in Fig.1. The single- andpoly-crystalline specimens were heated at 873 K for 1 h ina vacuum to remove an initially dissolved hydrogen, andthen exposed to 133 Pa of tritium-deuterium gas mixture(T/(T+D)~ 0.17) at temperatures ranging from 673 to 973
K for 1-5 h. No apparent pressure change was observed
during the tritium gas exposure. After that, the specimens

1. Introduction

Oxide materials are used in fusion reactors as plasma
diagnosis windows, electric insulators and tritium
permeation barriers [1-7]. Tritium dissolution andpermeation in the oxide materials are a significant issue
from a stand point of tritium retention and leakage.
However, compared with metals and alloys, tritium(hydrogen) solubilities in oxides have not been studied
well, because very low hydrogen solubility in oxides madedifficult to detect hydrogen in them [8].
In this study, we have tried to examine thehydrogen solubility in Y2O3 stabilized ZrO2(YSZ) using arather high concentration tritium gas and an imaging plate(IP) technique. The single- and poly-crystalline YSZ
specimens were used in order to clarify the influences ofgrain boundary and Y content on hydrogen solubility.

Intensity (arbitrary uint)

3000

2. Experimental

Mirror polished single crystal plates (1050.5mm3) of cubic zirconia (13%Y2O3-ZrO2(100)) from Dalian
Keri Optoelectronic Technology Co. Ltd. were used as
specimens. Polycrystalline specimens were prepared withconventional sintering process using PSZ powders
manufactured by Toso Co. Ltd (grades: TZ-3Y, TZ-4Y,
TZ-6Y, TZ-8Y and TZ-10Y, Y contents of which ranged
from 6 % (TZ-3Y) to 20 % (TZ-10Y)). The each grade
powder was die-pressed into disks and then isostaticallypressed at 200 MPa. The disks were sintered at 1773 K
for 20 h. The size of the disk specimens was about 8
mm in diameter and 1 mm in thickness. The densities of

TZ-10Y

2500
2000

TZ-8Y

1500

TZ-6Y

1000

TZ-4Y

500
TZ-3Y

0

20

30

40

50

60

70

80

2 theta (degree)

Fig.1 XRD spectra of polycrystalline specimens

author’s e-mail:hashi@nucl.kyushu-u.ac.jp

33©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

K. Hashizume et al., Solubility of Hydrogen Isotopes in Zirconia Ceramics(tritium activities) seemed to be rather uniform except formuch intense parts which should be surface contaminations,particularly in specimen edges where were not polished.
Since the high PSL intensity area was limited around theedges, these PSD data were omitted. The errors without
the contaminations were estimated to be r30 % for all
specimens. The PSL images of the cross sections suggested
an equilibrated dissolution of tritium.
Figure 3 shows the tritium activities (hydrogensurface densities) on the specimens exposed to the tritiumgas. The amount of tritium uptaken on the specimen
surface increased with increasing temperature up to 2105
Bq/cm2 (equivalently ~1015 H/cm2) and almost saturated at
873 K. At 973 K the tritium density on the single
crystalline surface decreased, where tritium desorption
should be more dominated than its adsorption. In any
case, no clear difference appeared among the specimens.
Figure 4 shows the tritium activities in the singleand poly-crystalline specimens which increase with theexposure temperature, which appear to be very similar tothat of its surface density. Figure 5 also shows the tritium
activities in poly-crystalline specimens against an oxygen
deficiency x which was simply calculated from Y contents
of YSZ. No clear difference in tritium activities appeared
for the single- and poly-crystalline specimens. Several
studies previously suggested that the hydrogen solubility ingrain boundary was high, compared with the oxide grain
inside. Unfortunately, the spatial resolution of IP technique
is not high enough to detect tritium only in the grainboundary. The influence of grain boundary would be
examined in the future study. In any case, it should besafely concluded that no clear influence of grain boundary
and O deficiency (Y content) in the present experimental
condition.

were quenched down to fix tritium distribution. After the
specimen surfaces were exposed to IP (TR2025, GE HealthCare Co.) for 15 min in order to measure surface tritium
densities, the specimens were cut in halves with a diamond
saw, and the cross sections of the specimens were exposed
to IP for 15 h. The details of this procedure were describedin elsewhere [7]. Photo-stimulated luminescence (PSL)
intensities were obtained from IP Reader (FLS7000,
Fujifilm). The PLS values were converted to tritium
activities by calibration with a tritium standard sample(ART 123A, American Radiolabeled Chemicals, Inc.).
Hydrogen surface density and solubility were evaluatedfrom the tritium activity and isotope abundance ratio.
3. Results and Discussion
Figure 2 shows the examples of the PSL images of
specimen surface and cross section. The PSL intensities
High T

Low T
5mm

Fig. 2 PSL images of surface and cross section.
Single crystal plate (left) and polycrystalline disk(right) exposed to D-T gas at 873 K for 1 h.
53h exposure15

1h exposure

Tritium activity (MBq/cm3)

H surface density (H/cm2)

105h exposure

5x10

T surface activity (Bq/cm2)

T/(T+D)~0.17

10

T /͠

5

1x10

5x104
14

4

<T1J$SEO

1.0

1.1

1.2

1.3

2x10
1.4

10

2

10

1

10

0

1000/T (1/K)

40013mol%Y O YSZ(single crystal)233mol%Y O YSZ(polycrystal)236mol%Y O YSZ(polycrystal)2

310mol%Y O YSZ(polycrystal)

10

1.5

500

600

2

-1

1.0

1.1

1.2

3

1.3

1.4

1.5

1.6

1000/T (1/K)

Fig. 3 Hydrogen isotope surface density (Tsurface activity) on single crystal specimen.

Fig. 4 Tritium activities in single- and polycrystalline specimens (5 h exposure to D-T gas).

34

K. Hashizume et al., Solubility of Hydrogen Isotopes in Zirconia Ceramics

crystal structures of zirconia (monoclinic, tetragonal andcubic) and hydrogen containing gas species (H2 or H2O).
The present study focused on the solubility ofhydrogen in zirconia ceramics. The dependence ofhydrogen amounts uptaken on surface and bulk in theoxides becomes clear using the IP technique. This also
indicates that the IP technique was a useful tool to study
the kinetics of hydrogen dissolution and hydrogensolubility in various oxides.

The color of the specimens after the tritium gas
exposure slightly changed to be brownish. Not only thecharacteristics of surface but that of bulk might beinfluenced by the reductive reaction. However, the effect of
this reaction is not discussed here because the adsorption
reaction of hydrogen is too complicated to make aquantitative analysis.

Tritium activity (MBq/cm3㪀

10

2
873K
773K
673K

10

1

10

0

-1

10

0

0.02

0.04

0.06

0.08

0.1

4.

Summary
Using an imaging plate (IP) technique and highconcentration tritium gas (T/(T+D)~0.17), tritium
solubility in zirconia ceramics (YSZ) was measured in atemperature region ranging from 673 K to 973 K. The
surface activity of tritium on the specimens increased withincreasing temperature and almost saturated at 873 K,
which reached 2105 Bq/cm2 (equivalently ~1015 H/cm2).
The hydrogen solubility in the oxides also increased
with increasing temperature. No clear difference in thesolubility appeared between single crystal andpolycrystalline specimens and among the specimens withvarious Y contents.

0.12

x in Zr(Y)O2-x

H mol / 1 mol ZrO2 at 133Pa

Fig. 5 Tritium activities vs O deficiency x inpoly-crystalline specimens.
10

-4

10

-5

10

-6

10

-7

5.

Acknowledgements
The present study was partly supported by theusage/research program of the Hydrogen Isotope ResearchCenter, University of Toyama (HRC2010-04) and by the
Sohoko Research Program of National Institute for Fusion
Science (NIFS11KUR014) .

6.  References[1] F. W. Clinard, Jr., G. F. Hurley, J. Nucl. Mater.,
103&104, 705 (1981).
[2] W. Dienst, T. Fett, R. Heidinger, H.D. Rhörig, B.
Shultz, J. Nucl. Mater., 174, 102 (1990).
[3] R. Heidinger, J. Nucl. Mater., 179-181 (1991) 64.
[4] E.R. Hodgson, Nucl. Instr. Met. Phys. Res., B191,
744 (2002).
[5] A. Ibarra, E.R. Hodgson, Nucl. Instr. Met. Phys.
Res., B218, 29 (2004) .
[6] Y. Hatano, K. Zhang, K. Hashizume, Physica Scripta,
T145, 0140144 (2011).
[7] K. Hashizume, K Ogata, M. Nishikawa, T. Tanabe,
A. Abe, S. Akamaru, Y Hatano, Proceeding of
ICFRM-15, to be published.
[8] T. Norby, Solid State Ionics, 40/41, 857 (1990) .
[9] V.C. Wagner, Ber. Bunsenges. Physik Chem., 72,
778 (1968) .
[10] K. Park, D.R. Olander, J. Am. Ceram. Soc., 74, 72(1991) .
[11] S. Yamanaka, T. Nishizaki, M. Uno, M. Katsura, J.
Alloys Compd., 293-295, 38 (1999) .
[12] M. Miyake, M. Uno, S. Yamanaka, J. Nucl. Mater.,
270, 233 (1999) .

present study (c-ZrO2 in H2)
Yamanaka (m-ZrO 2 in H2O)
Wagner (c-ZrO2 in H2O)

10

-8

0.4

Park (t-ZrO2 in H2)

0.6

0.8

1

1.2

1000/T (1/K)

1.4

1.6

Fig. 6 Comparison of hydrogen solubility in zirconia(5h-exposed single-crystal for the present data).
The present data on hydrogen solubility are comparedwith literature data [9-12] in Fig. 6. The all literature data
have been converted to those at a pressure of 133 Pa ofhydrogen containing gases (H2 or H2O), based on relationss v p1H/ 22 and s v p1H/ 22O . As shown in Fig. 6, thevalues of solubility and its temperature dependence arerather scattered. Reasons of the scattering are not clear atpresent stage, but might be attributed to differences inexperimental conditions such as dissolution temperatures,

35

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Numerical Simulation of Detritiation System for NIFS with
Commercial Catalyst and Adsorbent
K. Munakataa,*, K. Haraa, T. Sugiyamab, K. Kotohc, M. Tanakad and T. Udad
*kenzo @gipc.akita-u.ac.jpaFaculty of Engineering and Resource Science, Akita University, Tegata-gakuen-machi 1-1, AKITA 010-8502 Japanb
Nagoya University, Furo-cho Chikusa-ku, Nagoya464-8603, Japanc
Kyushu University, Motooka 744, Nishi-ku, Fukuoka 819-0395, Japand
National Institute for Fusion Science, Oroshi-cho 322-6, Toki, Gifu 509-5292 Japan(Received: 9 May 2012 / Accepted: 18 August 2012)

Establishment of tritium recovery system is necessary to prevent its leakage to the working area. The catalyticoxidation and adsorption is the most reliable method to recover tritium released into the working area of thosefacilities. Therefore, there is a necessity that the tritium recovery system with large-scale and higher integrity isdeveloped and constructed. For this purpose, it is required to obtain and accumulate updated database for chemicalengineering design based on most recently commercial and available catalysts and absorbents. In this work, we
selected a catalyst and an adsorbent, and examined their adsorption characteristics for water vapor and performanceof catalytic oxidation of hydrogen. Adsorption isotherms were studied with a volumetric gas adsorption instrument.
Various adsorption models were tested to correlate the experimental isotherms. Breakthrough experiments werealso performed to assess adsorption performances. Furthermore, numerical simulations were performed based on
the obtained experimental results.
Keywords: detritiation, catalyst, adsorbent, membrane dehumidifier, numerical simulation, mass balance
1. Introduction

The authors went through previous literature andinquired about catalysts or adsorbents used for
detritiation systems that are being used in several
research institutes. However, the result of such survey
reveals that the production or model number of thematerials such as catalyst and adsorbent is not clear, andthey are not considered to be available in the present
industrial market. Additionally, conduction of new
screening tests, in order to search feasible catalysts or
adsorbents, is time consuming. It should be also notedthat feasibility tests of catalysts and adsorbents are notlimited within catalytic activity or adsorption
performance; durability in long-time use, mechanicalstability, ease in their procurement with large quantities
are important points as well.

A large amount of tritium would be handled as fuelin nuclear fusion power plants. In these facilities, tritium
is confined in multiple confinement barriers. However, if
unexpected accidents take place, tritium leaks to theworking areas. Thus, establishment of tritium recovery
system is necessary to prevent its leakage to the working
area. The catalytic oxidation and adsorption is the mostreliable method to recover tritium released into theworking area of those facilities. In this method, tritium
gas is oxidized with catalysts, and then tritiated water
vapor is collected by adsorbents.
It is planned to conduct deuterium plasma
experiments in National Institute for Fusion Science(NIFS). In these experiments, D-D fusion reactions takeplace and tritium is produced, and thus the release oftritium to the environment has to be minimized.
Therefore, there is a necessity that the tritium recovery
system with large-scale and higher integrity is developed
and constructed. However the previous experimental andengineering database [1] obtained by several researchers
on absorbents cannot be applied to the design of updated
tritium recovery system since the absorbents used in theprevious studies are not been produced any longer in theindustry. Therefore, it is required to obtain andaccumulate updated database for chemical engineering
design by testing most recently commercial andavailable catalysts and absorbents in experiments. In thiswork, we selected an adsorbent and examined its
adsorption characteristics for water vapor.

For this reason, the authors selected a catalysts andan adsorbent, which are ones of the most widely andgenerally used in industries and are credited and reputedby engineers and designers as well.
In terms of catalyst, the authors selected a Pt/alumina
catalyst, DASH520, manufactured by N. E. Chemcat Co.,
which is widely used in industries for various purposes.
The catalyst was deposited with 4.1g/L of platinum. The
average diameter of the catalyst is 3.25 mm. the packing
density was 770 g/L.
With regard to the adsorbent, the author selected a
MS5A adsorbent that is produced by Union Showa Co.
in Japan (subsidiary company of Union Carbide Co.) andis widely used in industries. The production of zeolites inUnion Showa Co. is licensed by Union Carbide Co, andthe pellet-type MS5A adsorbent could be one of theworld-standard adsorbents.

2. Catalyst and adsorbent_______________________________________________________________________________
Corresponding author: kenzo@gipc.akita-u.ac.jp

36©2013 by The Japan Society of Plasma
Science and Nuclear Fusion ResearchK. Munakata et al., Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent

3. Experimental
3.1 Catalytic oxidation of hydrogenk r ,ox ,0a 1.16×105 exp 25400 /(RT )(4)

The experimental adsorption isotherms of water
vapor on the DASH520 catalyst and MS5A adsorbent
were found to be successfully correlated using a sort ofmulti-site Langmuir-Freundrich equation:nqad ,s ,mbm pv m

¦ 1  b pv

qad

m 1

m

(5)

m

bm b0,m exp(E m / RT )(6)

Parameters (optimized by non-linear least squares
analysis) for DASH520 catalyst (n=5) and MS5A
adsorbent (n = 4) are listed in Table 1.
Experimental results on the catalytic oxidation ofhydrogen in wet gases reveal that coexistence of water
vapor severely inhibits the catalytic oxidation ofhydrogen. It was found that the mass transfer coefficientfor catalytic oxidation in wet gases is expressed as afunction of the amount of water adsorbed on the catalyst
substrate:

The authors studied adsorption behavior of water
vapor on the DASH520 catalyst and MS5A adsorbent, aswell. The adsorption isotherms of water vapor on the
DASH520 catalyst and MS5A adsorbent were measured
using a volumetric gas adsorption instrument,
BELSORP-max, manufactured by BEL Japan Inc. The
instrument is designed for measurement of wide range ofadsorption isotherms on surface area and pore size
distribution analysis.

k r ,ox a

3.3 Summary of experimental results

k r ,ox ,0aConversion rate Rc and over all mass transfer
capacitance for oxidation (based on surface reaction) in
dry condition kr,ox,0a were computed to assess catalytic
performance. RC and KF,ox are defined as
SV Q /V(1)

100 u (C in C out )/C in(3)

The mass transfer capacitances for oxidation ofhydrogen (based on surface reaction) of DASH520 isexpressed as

The experiments were performed under the steadystate condition. The catalysts were packed in a reactor
made of quartz. The temperature of the reactor wascontrolled with the constant temperature bath. Argon gas
was used as a carrier gas. The argon gas containinghydrogen (about 300-800 ppm) and oxygen (20%) was
introduced to the reactor.
Experiments were also performed using wet processgases which contained water vapor (about 300-1000 Pa)
in order to study the influence of coexistent water vapor
on the catalytic activity for oxidation of hydrogen. The
concentrations of hydrogen at inlet and outlet stream of
the reactor were measured with a gas-chromatograph(GC-8A) manufactured by SHIMAZU Co.
3.2 Water adsorptionRCk r ,ox ,0a SV ln C in /C out

0.476exp ª  643qad
¬«

25.1

º
¼»

(7)

5.00
º
0.524exp ª  456qad¬«¼»

4. Simulation procedure and mass balance
The basic assumption on catalyst and adsorption bed
is one-dimensional dispersed plug flow in packed bed [2].
Chemical reactions such as catalytic oxidation ofhydrogen isotopes, adsorption of water vapor, isotope
exchange reaction were taken into consideration. The
mass balance equations in catalyst beds are expressed asfollows (for adsorption bed, catalytic reactions are notconsidered):(2)

Table 1 Parameters optimized for L-F equation.
m DASH520
MS5A
qs,m [mol/g]
1.52×10-4
19.8
3.86×10-13
2.02×10-15
b0,m [(mol/(gPa))1/ui]
1
39100
44700
Em [J/(molK)]vm [-]
3.16
0.908
qs,m [mol/g]
4.58
0.00703
1.49×10-22
8.90×10-20
b0,m [(mol/(gPa))1/ui]
2
75100
92000
Em [J/(molK)]vm [-]
1.76
2.67
qs,m [mol/g]
3.57×10-40.00191
1.43×10-7
1.42×10-9b0,m [(mol/(gPa))1/ui]
3
38800
36100
Em [J/(molK)]vm [-]
3.25
4.05
qs,m [mol/g]
1.70×10-50.000485
6.62×10-10
4.98×10-12b0,m [(mol/(gPa))1/ui]
4
34800
64400
Em [J/(molK)]vm [-]
1.93
6.33
qs,m [mol/g]
2.12×10-3-91/u
9.85×10b0,m [(mol/(gPa)) i]
532800
Em [J/(molK)]vm [-]
2.09-

㺃Mass balance of Q2 (i = H, D or T) in gas phasew(uC Q 2,i )wz

HwC Q 2,iwt

k F ,ox aC Q 2,i1

k F ,ox a1k r ,ox aHw § wC Q 2,i ·¨D¸wz ¨© L wz ¸¹

1

(8)

(9)

kga

㺃Mass balance of Q2O (i = H, D or T) in gas phasew(uC Q 2O ,i )
wz

37

H

wC Q 2O ,iwt

Hw § wC Q 2O ,i ·¨D¸wz ¨© L wz ¸¹(10)

K. Munakata et al., Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent

vapor in equilibrium increases almost linear with theactivity of water vapor [5]. In the present study, water
vapor content in the matrix of support layer in
equilibrium is expressed by the following equation.

k F ,ox aC Q 2,i k F ,ad ,i a cQ 2O ,H  cQ 2O ,D  cQ 2O ,T T1
k F ,ex aT2

̅ � = ����� ���,� /�� ���,�㺃Mass balance of Q2O (i = H, D or T) in solid phaseJ

wqQ 2O ,iwz

qQ 2O ,ik F ,ad ,i a cQ 2O ,H  cQ 2O ,D  cQ 2O ,T T1

k F ,ex aT2

qad ,Q 2O ,i  qst ,Q 2O ,i

T1 cQ 2O ,i 

Sorption rate of water vapor is, then, expressed as�̅ � � ��,�
� = ���� ����� ���,�
�.

(11)

qQ 2O ,H  qQ 2O ,D  qQ 2O ,T(13)��,� = ��,�,� � �� ���,� /�� �

when cQ 2O ,H  cQ 2O ,D  cQ 2O ,T t 0

T2 cQ 2O ,i / cQ 2O ,H  cQ 2O ,D  cQ 2O ,T���= 0 = ��,� � ��,��� � ���,� �� ,
��
���,�
���,�
= ���,� �� � ���,� �� � ����T2 qQ 2O ,i / qQ 2O ,H  qQ 2O ,D  qQ 2O ,T��,�

The heat balance equation is given as

w u U g c p , gT


wz

H

w §
wT ·
¨ Oeff¸
wz ©
wz ¹

w Ug c p ,gTwt

Jw c p ,sT

��,�

(15)

���,�
= ��,� ��,� � ��,��� ��,��� � ���,�,� �� ,�����,�,�
���,�
= ���,�,� �� � ���,�,� �� � �� ,�����,�,�
���,�
= ���,�,� �� � ���,�,� �� ,��k F ,ad ,T a(C Q 2O ,H  CQ 2O ,D  CQ 2O ,T C )'H ad4hw (T Tair )/ Db

The assumptions used in the modeling of membrane
dehumidifier are: (i) the feed gas is a binary mixture ofwater vapor and air, (ii) the pressure drops and diffusive
mass transfer in the feed side and permeation side are
negligible, (iii) laminar films on the surface of themembrane are negligible and (iv) the system is adiabaticand the temperature of the gas mixture is constant [3].
The molar fluxes of penetrants through the dense andporous layers can be expressed as,��,� = ��,�

���,�
.
��

���
= 0 = ��,��� � ��,� � ���,� �� .
����,� ��*= ��,� ���,� � ��,� �,

(22)

(23)

The mass balances for each component are, then,(k F ,ox ,H C Q 2,H  k F ,ox ,DC Q 2,D  k F ,ox ,T C Q 2,T )a 'H ox

���,� ��,� ��=���,� � ��,� �
��
��

(21)

and

wt

��,� = ��,�(20)

The membrane is divided into small sections along
the axis, and a kind of plate model was applied. The total
mass balances are expressed as(14)

when cQ 2O ,H  cQ 2O ,D  cQ 2O ,T  0(19)

It is known that a very low partial pressure of acondensable species, such as water vapor in the feedstream, can significantly reduce the permeability of theother constituent relative to its permeability as a purecomponent. In the present study, the permeance of air isexpressed by the following equation at relatively low
water vapor activity.

(12)

cQ 2O ,H  cQ 2O ,D  cQ 2O ,T qQ 2O ,i(18)(1 � �)��,�
and��,� ��

����,�,�
= ��
��

���,�
= ��,��� ��,��� � ��,� ��,� � ���,�,� �� .
��

(24)
(25)
(26)
(27)

(28)

Equations (25) and (26) represent the mass transfers in
the pore of the support layer. On the other hand, Eq. (27)represents the concentration change of water vapor in thematrix of the support layer. Equations (29)-(31) and therelationships, ���,�,� = ��,� = ��,� and ���,�,� = ��,�give Eqs. (32) and (34) as fluxes in which the massconservations across the membrane is taken into account.
In the same way, Eqs. (33) and (35) are obtained for the
air.

(16)

(17)

In Eq. (16), a constant pressure gradient is assumed anda permeance is applied because the thickness of thedense layer is very thin and the value is not disclosed by
the vendor. The diffusion coefficient ��,� in Eq. (17) isestimated as Knudsen’s diffusion coefficient.

��,� =��,������� �� ��,� � �� ��,�,�
���
��

���
���
= ��,� ��� ��,� � ��,�,�

Polyimide membranes absorb moisture easily at highpressures of water vapor. The sorption capacity of water

38

(29)

K. Munakata et al., Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent���� =

����� ���������� − ������ ����� ����
������������ −������������ =����� ������ � − ������ ������� �� �� � ����� � ���������� =����� �������� �� �� � �����

������� =��������� − �� ���� ��� �� �����

������� =� ��� �� − ���� � − ������ ���������� �� − �� �� − ���� ���� �� �����

����� = ������� � �������

����� = ������� � �������

Table 2 Operational condition of detritiation system.

(30)

Detritiation of vacuum
vessel purge gas
CQ2,H: 0.5 ppm
CQ2,D: 90.0 ppm
CQ2,T: 1×10-8 ppm
CQ2O,H: 1500.0 ppm
CQ2O,D: 90.0 ppm
CQ2O,T: 1×10-8 ppm
Amount of catalyst:
100 kgTemp. of catalyst䠖30Υ,
100Υ, 200Υ
Heating䠖㻌 heatexchanger, adiabatic
Recovery䠖 Polymer
membrane-type
dehumidifier UM-XC5
䠄Φ90 mm×710 mm䠅, 60
sets䠈high pressure 0.6
MPa䠈low presasure0.1
MPa, 30 Υ䠈100 L/min
Dehumidifier cooling䠖㻌
heat exchanger, adiabatic

(31)
(32)

(33)

(34)
(35)
(36)
(37)

The mass balance equations discredited and thenmatrices generated were solved by direct methods such
as LU decomposition method. The generalized CrankNicholson scheme was applied for time-marchingprocedure.
5. Results and discussion

Detritiation of plasma
exhaustCQ2,H: 130.0ppm
CQ2,D: 2983.3 ppm
CQ2,T: 9.25 ×10-6 ppm
CQ2O,H: 3.6667 ppm
CQ2O,D: 28.667 ppm
CQ2O,T: 6×10-8 ppm
Amount of catalyst:
100 kgTemp. of catalyst䠖30Υ,
100Υ, 200Υ, 400Υ
Heating䠖㻌 heatexchanger, adiabatic
Recovery䠖㻌 MS5A
adsorbent,100 kg

Adsorbent cooling䠖㻌
heat exchanger,
adiabatic

plasma exhaust gases are treated using catalyst andadsorption beds. One of the major differences isthroughput of the gases processed. The concentrations oftritium and deuterium are considerably higher in theplasma exhaust gas. Numerical simulations were
performed for both cases described above.

5.1 Detritiation system design of NIFSIn D-D experiments in NIFS, (a) detritiation ofdownstream of outgas generated in cleaning of plasma
container (vacuum vessel purge gas) and (b) plasma
exhaust gases is required. Thus, it is necessary todevelop a detritiation system to recover lowconcentration tritium generated in LHD (Large Helical
Device). The schematics of the detritiation systemdesigned at present are shown in Fig. 1. Representative
operational conditions are summarized in Table 2. With
regard to the vacuum vessel purge gas, the process gas isfirst introduced to 100 kg of catalyst bed and thentritiated water vapor is collected by condenser andpolymer membrane-type dehumidifier. In contrast, the

5.2 Results of numerical simulation for detritiation of
vacuum vessel purge gas

CQ2,T,out [ppm]Figure 2 shows changes in the concentration ofmolecular form of tritium in the outlet stream of thecatalyst bed. With regard to the temperature of inlet gas,
three cases (30 qC, 100 qC and 200 qC) were tested. It
can be seen in the figure that catalytic activity is
considerably decreased in the cases where the inlet gas
temperatures are 30 qC and 100 qC. This is because thatwater vapor contained in the process gas inhibits
10

-8

10

-9

10

-10

10

-11

10

-12

10

-13

10

-14

㻌

30 qC

100 qC㻌���� =

200 qC

0

5

10

15

20

Fig. 1 Detritiation systems designed in NIFS

Time [h](a) outgas generated in cleaning of plasma container(b) downstream of plasma exhaust gases.

Fig. 2 Change in concentration T2 in outlet stream of
Pt/alumina catalyst beds.

39

K. Munakata et al., Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent

catalytic activity with increase in the amount of water
adsorbed on the catalyst substrate. However, this effectcould be considerably avoided if the inlet gas
temperature is raised to 200 qC as shown in the figure.
Thus, preheating of the process gas with heat exchangers
is necessary for the treatment of vacuum vessel purgegas.

increases in the same trend. Figure 4 shows changes inconcentrations of molecular form of tritium and tritiatedwater vapor at the outlet of the duct. The concentrationof tritiated water vapor increases almost linearly withlapse time. The molecular form of tritium sharply
increases 20 h after the commencement of operation of
the detritiation system, since the inhibition of catalytic
activity by water vapor becomes a dominant factor.
However, the result of this simulation indicates that
regulatory level of tritium concentration at the outlet ofthe duct is fulfilled.

Subsequent numerical simulation for condenser andpolymer membrane-type dehumidifier was performed in
the case where the inlet gas temperature is raised to 200
qC. The transient variations of chemical components
such as hydrogen isotopes and isotopes of water vapor in
the outlet stream of the catalyst bed were used as theinlet condition of the polymer membrane-type
dehumidifier. Figure 3 shows changes in amounts ofbulk water and concentration of tritiated water in draintank. The amount of bulk water increases almost linearlywith lapse time. The concentration of tritiated water

5.3 Results of numerical simulation for detritiation ofplasma exhaust gas
Figure 5 (a) shows changes in T2 concentration inoutlet stream of the catalysts bed. In this case, preheatingof the process gas before the catalyst bed is unnecessary
because of the help of higher concentration of D2; its
reaction can generate heat and raise the reactor

(a)

(b)

Fig. 3 Change in (a) amount of bulk water and (b) concentration of tritium in drain tank.

(a)

(b)

Fig. 4 Change in (a) concentration of molecular form of tritiumand (b) concentration of tiritiated water vapor on outlet stream of duct.
-5

10

(a)

-10

10

-15

10

-15

10

-16

10

-17

(b)

-20

10

CQ2O,T,out [ppm]

CQ2,T,out [ppm]

-14

30 qC

10

-25

10

-30

10

-35

10

-40

10

-45

10

-50

10

-55

200 qC

100 qC

400 qC

30 qC

100 qC

10

-18

-60

10

10

0

5

10

15

10

20

Time [h]

0

5

10

15

20

Time [h]Fig. 5 Change in (a) concentration of molecular form of tritium in outlet gas of catalysts bed
and (b) concentration of tiritiated water vapor on outlet stream of adsorption bed.
(Temperatures shown in the figures represent the temperature of the process as at inlet of the catalyst bed.)40

K. Munakata et al., Numerical Simulation of Detritiation System for NIFS with Commercial Catalyst and Adsorbent

t
T
u
V
V

temperature. Figure 5 (b) shows breakthrough curve forT2O vapor, which indicates that breakthrough starts at 10
h after the start of operation of 100 kg of MS5A
adsorbent regardless of variation in temperature ofprocess gas at the inlet of the catalyst bed.

XZ

6. Conclusion
Catalytic oxidation rate of hydrogen and water
adsorption characteristics are investigated with regards
to commercial catalyst and adsorbent. Based on theexperimental results, simulation models wereconstructed for catalyst bed, adsorption bed andpolymer-type membrane dehumidifier. Test processsimulations were carried out on the basis of the
detritiation systems designed in NIFS.

Greek lettersΔ
Thickness of the layer [m]ΕVoid ratio in the porous layer
Oeff
Effective heat conductivity in axial direction [J
m-1 s-1]
Ug
Density of gas [g m-3]
Overscript
Int
Interface between dense and porous layer¯
Equilibrium state'
Matrix of the porous layer

Nomenclature

a
asorp
ap
b0

cp,g
cp,s
C
D
Db
'Had
'Hox
kF,ad
kF,ox
Kg

kr,ox

kr,ox,0
ksorp
l
N
p
P
q
qad
qst
q

Q
r
R
Rc
S
SV
Sw

Time [s]Temperature [K]Superficial gas velocity [m s-1]
Volume [m3]Constant related to Langmuir-Freundrichequation [-]Mole fraction of water vapor [-]Length in axial direction [m]Specific surface area [m2 m-3]
Slope of the adsorption isotherm [mol m−3]
Coefficient in Eq. (20) [mol m−2 s−1 Pa−1]
Constant related to Langmuir-Freundrichequation [arbitrary]Specific heat of gas [J g-1 K-1]
Specific heat of solid phase [J kg-1 K-1]
Molar concentration [mol m−3]
Diffusion coefficient [m2 s−1]
Bed diameter [m]Adsorption heat [J mol-1]
Reaction heat for catalytic oxidation [J mol-1]
Overall mass transfer coefficient for adsorption[m/s]Overall mass transfer coefficient for catalyticoxidation [m/s]Mass transfer coefficient for gas film layer[m/s]
Mass transfer coefficient for catalytic oxidation
based on surface reaction [m/s]Mass transfer coefficient for catalytic oxidation
based on surface reaction in dry condition[m/s]Rate constant of the adsorption [s−1]Position in the permeate direction [m]Permeation flux [mol m−2 s−1]
Pressure or partial pressure [Pa]Permeance [mol m−2 s−1 Pa−1]
Flow rate [mol s−1]
Concentration of water adsorbed [mol g-1]
Concentration of water strongly bound tosurface (structural water) [mol g-1]
Net Concentration of water species on thesurface [mol g-1]
Volumetric gas velocity [m3 s-1]
Adsorption rate of water vapor [mol s−1]
Gas constant [J K−1 mol−1]
Conversion rate [-]Outside wall area of a calculation cell [m2]Space velocity [s-1]
Solubility coefficient of water [mol m−3 Pa−1]

Subscripts
0Initial value
1
Feed side
12
From feed side to the porous layer
2
Porous layer
23
From porous layer to the permeation side
3
Permeation sideA
AirA
Indicator in Fig. 2
B
Indicator in Fig. 2
C
Indicator in Fig. 2
D
Dense layer
HHigh pressure sideI
Component (H, D or T) or (w or a)
In
inlet
J
Number of the calculation cell along the axis
L
Low pressure sideM
Adsorption site number
N
Maximum adsorption site numberOutoutletQ
H, D, or T (1H, 2H or 3H)
S
Saturated state
W
Water vapor

References
[1]

[2]

[3]

41

K. Kotoh and K. Kudo, “Applicability of Commercial
Adsorbents to Multi-column Adsorption Systems with
Condenser for Tritiated Water Vapor Removal”, Journal of
Nuclear Science and Technology, 34, 1099(1997)
K. Munakata, M. Nishikawa, T. Takeishi, N. Mitsuishi
and M. Enoeda, “Recovery of Tritium in Room Air by
Precious Metal Catalysts with Hydrophilic Substrate”,
Journal of Nuclear Science and Technology. 25 , 383(1988)
T. Sugiyama, N. Miyahara, M. Tanaka, K. Munakata andI. Yamamoto, “A Simulation Model for Transient
Response of A Gas Separation Module Using A Hollow
Fiber Membrane”, Fusion Engineering and Design, 86,2743(2011).

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Effects of La2O3, Cu and Fe addition on the catalytic performance of
Ni-SiO2 catalysts for methane decomposition
Heyi Wang,* Wenhua Wang, YongYangˈShuming Peng
Institute of Nuclear Physics and Chemistry, China Academy of Engineering Physics, Mianyang, 621900, PR China(Received: 9 May 2012 / Accepted: 6 November 2012)

Methane decomposition over four kinds of catalysts such as Ni-SiO2, Ni-La2O3-SiO2, Ni-Cu-SiO2 andNi-Fe-SiO2 prepared by co-impregnation method and sol-gel method, respectively, has been studied. The results
indicated that Ni-La2O3-SiO2 catalyst with La/Ni = 0.3 prepared by co-impregnation method showed better catalytic
performance compared with all the other Ni-La2O3-SiO2 catalysts during methane decomposition at 823 K.
However, the temperature resistance of this catalyst was not very good. The introduction of Cu and Fe into Ni-SiO2
enhanced the temperature resistance of the catalysts during methane decomposition at 923 K and the overall
catalytic performance of the catalysts was Ni-Cu-SiO2 (sol-gel method) > Ni-Fe-SiO2 (sol-gel method) >75%Ni-25%SiO2 (sol-gel method) > Ni-La2O3-SiO2 (co-impregnation method). XRD studies of Ni-La2O3-SiO2,
Ni-Cu-SiO2 and Ni-Fe-SiO2 catalysts indicated that Ni species existed in the form of Ni metal in Ni-La2O3-SiO2,
while Ni-Cu alloy and Ni-Fe alloy appeared in Ni-Cu-SiO2 and Ni-Fe-SiO2 catalysts. The structure of the carbon
filaments formed during methane decomposition changed a lot with the introduction of Cu and Fe. At 923 K,“bifurcate carbon filaments” formed over Ni-Cu-SiO2 during methane decomposition, while “bamboo-shaped
carbon filaments” formed over Ni-Fe-SiO2 catalysts.
Keywords: Methane decomposition, Ni-SiO2, Ni-La2O3-SiO2, Ni-Cu-SiO2, Ni-Fe-SiO2

1. Introduction

were hardly investigated. Above all, the unsupported
alloy catalysts in the dynamic systems will deactivate
quickly. Methane decomposition over the supported
catalysts to prepare H2 and carbon filament in the
dynamic system has been studied by many researchersand many achievements have been obtained in recent
years [12-16]. However, the effects of the rare metaloxide, Fe and Cu addition on the temperature resistanceof the Ni-based catalysts were little investigated. In
view of this, methane decomposition over supported
catalysts such as Ni-SiO2, Ni-La2O3-SiO2, Ni-Cu-SiO2
and Ni-Fe-SiO2 catalysts in dynamic systems has beenstudied in this work. The influences of the preparation
method, the reaction temperature and the contents of La,
Cu and Fe in Ni-SiO2 on methane decomposition were
investigated.

If graphite and carbon fiber composite are employed
as plasma facing components for ITER, the deuterated
and tritiated methane may be a considerable constituentof the impurity gas stream [1-6]. Methane decomposition
is a safe and simple method for recovering deuterium andtritium from the deuterated and tritiated methane and an
optional catalyst bed for methane decomposition iscurrently included into the design of the tokamak exhaustprocessing system (TEP) [7,8].
Lots of research on methane decomposition has been
done by many researchers [9-11]. The ZrCo powder has
already been found to decompose methane slowly at
573 K. A product of 0.05 molC/mol ZrCo was obtained
after 80 h when the getter was exposed to 430 torr CH4[9]. Watanabe [10] et al investigated the decompositionof methane on Zr7Ni10, ZrNi, Zr2Ni, Zr4Ni and Zr in thetemperature range 573K–823 K and the alloy of Zr4Ni
removed methane over 99% efficiently in 105 s at 673 K.
Baker [11] et al discovered that a St 909 getter(Zr–Mn–Fe) could effectively remove 99% of NH3, CO2
and O2 at 673 K, CO at 873 K, and CH4 at 973 K from aninert gas stream.
However, methane decomposition for recovering
deuterium and tritium which was studied by the above
researchers all focused on the static systems and theunsupported alloy catalysts, while the dynamic systems

2. Experimental
2.1. Catalyst preparation

Ni-La2O3-SiO2 catalysts were prepared byco-impregnation method [17] which was done throughimpregnating SiO2 with the mixed aqueous solution ofNi(NO3)2 and La(NO3)3 for 1.5 h, and then drying up the
impregnated samples at 413 K for 12 h, and finallycalcined them at 873 K for 6 h. The Ni content in all thecatalysts was Ni/Si = 0.67 (molar ratio) and the La
contents in the catalysts were La/Ni = 0, La/Ni = 0.1,

author’s e-mail:hywang@caep.ac.cn

42©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

La/Ni = 0.3, La/Ni = 0.45 and La/Ni = 0.6 (molar ratio),
respectively.
75%Ni-25%SiO2, 65%Ni-10%Fe-25%SiO2, 35%Ni40%Fe-25%SiO2, 65%Ni-10%Cu-25%SiO2, and 35%
Ni-40%Cu-25%SiO2 catalysts (mass fraction) wereprepared by sol-gel method [18,19]. The method wasbased on mixing the active components such as NiO or
the mixtures of NiO and CuO (or Fe2O3) in this case, withalcosol containing silica in a certain amount. The
suspends of the active components with alcosol weredried in flowing air at room temperature and were finallycalcined at 923 K for 3 h. Alcosol was prepared bymixing 50 mL of TEOS, 40 mL of ethanol, 2 mL of waterand 0.5 mL of 40% HCl. Silica content in the alcosol was
0.142 g/mL and could be diluted by ethanol. The mixtureof NiO and CuO (or Fe2O3) was originated from calcining
the mixture of Ni(NO3)2 and Cu(NO3)2 (or Fe(NO3)3) at
723 K and the Ni(NO3)2 and Cu(NO3)2 (or Fe(NO3)3)
mixture was obtained by evaporating the mixed solutionof Ni(NO3)2 and Cu(NO3)2 (or Fe(NO3)3).

35m(catalysts)=60mg

La/Ni=0
La/Ni=0.1
La/Ni=0.3
La/Ni=0.45
La/Ni=0.6

Methane conversion/%

30
25
20
15
10
5

0

50

100

150

200

250

Time/min

300

350

400

450

Fig.1 Kinetic curves of methane decomposition over
Ni-La2O3-SiO2 catalysts at 823 K.
35

773k
823k
873k

m(catalysts)=60mg

Methane conversion/%30

2.2 Activity tests and catalyst characterization

Methane catalytic decomposition reactions werecarried out in a fixed-bed quartz reactor (10 mm i.d.)under atmospheric pressure. Before the activity tests,
Ni-La2O3-SiO2 catalysts were subjected to a reduction
pre-treatment using pure hydrogen of 20 mL/min at 873K
for 1h and Ni-Cu-SiO2 and Ni-Fe-SiO2 catalysts werereduced with pure hydrogen of 20 mL/min at 923 K for 1
h. The feed stream was high-purified methane (99.99%)
of 15 mL/min. The gaseous reaction products weremonitored by an off-line gas chromatography (GC)combined with a 5A molecular column, using a thermal
conductivity detector (TCD) for hydrogen and methane
analysis.
XRD patterns were recorded on a Bruker
D8Advance diffractometer with Cu KD radiation at
40 kV and 40 mA.
The detailed structure of the deposited carbons wasobserved using the FEI Inspect F scanning electron
microscope (SEM) operated at 20 kV and the JEOL
JEM-100CX Transmission electron microscopy (TEM),
operated at 80 kV.

25
20
15
10
5
0

0

100

200

300

400

Time/minFig.2 Effect of reaction temperature on methane decomposition
over Ni-La2O3-SiO2 catalysts with La/Ni = 0.3.

of Ni-SiO2 catalyst obviously. For the Ni-La2O3-SiO2
catalyst with La/Ni = 0.1, the methane conversion
decreased from 16% to 7% after 450 minutes. With theincrease of La2O3 content from La/Ni = 0.1 to La/Ni = 0.3,
the catalytic performance of the catalyst was greatly
improved and the methane conversion decreased from
17% to 13% with only 4% dropped after 450 minutes.
However, when La2O3 content in the catalyst exceeded
La/Ni = 0.3, the catalytic performance of the catalysts
dropped slightly. The methane conversion decreased from
15% to 7.5% for Ni-SiO2 with La/Ni = 0.6 after 450
minutes. Thus, a conclusion can be drawn that an
optimum La2O3 content exists in Ni-La2O3-SiO2 catalysts
which can improve the catalytic performance of Ni-SiO2
the most efficiently, while the introduction of more La2O3
into Ni-SiO2 catalyst will impair the catalytic activity andthe lifetime of the catalysts and this optimum La2O3
content is La/Ni = 0.3.
Fig.2 shows the effect of reaction temperature on
methane decomposition over Ni-La2O3-SiO2 catalyst with
La/Ni = 0.3. The catalytic activity of Ni-La2O3-SiO2
with La/Ni = 0.3 increased remarkably from 773 K to 823
K. When the reaction temperature increased from 823 K
to 873 K, the initial methane conversion increased to 20%,

3. Results and discussion

3.1 The catalytic performance of the catalysts
Fig.1 depicts the kinetic curves of methane
decomposition over Ni-La2O3-SiO2 catalysts withdifferent contents of La2O3 at 823 K. Only hydrogen was
obtained as a gaseous product for all the catalysts. It
could be observed that the catalytic activity of Ni-SiO2
catalyst without La2O3 decreased gradually with the timeon stream and it was completely deactivated after
450 minutes. The addition of La2O3 extended the lifetime

43

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

Ni-Cu-SiO2 didn’t change obviously during methane
decomposition.
The
methaneconversion
of
65%Ni-10%Cu-25%SiO2 catalyst was kept at around40%, while when the Cu content increased to 40%, themethane conversion dropped to about 19%. Therefore, it
was different from Ni-La2O3-SiO2 catalyst that
Ni-Cu-SiO2
catalysts
could
catalyzemethanedecomposition at a higher reaction temperature such as923 K and the introduction of Cu into Ni-SiO2 enhanced
the temperature resistance of the catalysts.
Fig.4(b) shows the kinetic curves of methane
decomposition over Ni-SiO2 and Ni-Fe-SiO2 catalysts at
923 K. The catalytic behaviors of Ni-Fe-SiO2 catalysts
were similar with that of Ni-Cu-SiO2 catalysts. The
introduction of Fe also improved the lifetime of Ni-SiO2
catalysts obviously at 923 K. However, a deactivation
period occurred for 65%Ni-10%Fe-25%SiO2 catalyst andwhenmethanedecomposition
over
65%Ni-10%Fe-25%SiO2
catalyst
proceeded
to200 min, the methane conversion decreased from 44% to25%. While, the methane conversion of 35%Ni-40%Fe-25%SiO2 was maintained at 19% all the time whichindicated that with the increase of the Fe content in the
Ni-Fe-SiO2 catalysts the catalytic activity of the catalyst
dropped and the stability of Ni-Fe-SiO2 catalyst wasenhanced. Therefore, Ni-Fe-SiO2 could also catalyze
methane decomposition at a higher reaction temperaturesuch as 923 K which was similar with Ni-Cu-SiO2.
However,thecatalytic
performanceof
65%Ni-10%Fe-25%SiO2 was a little poorer than that of
65%Ni-10%Cu-25%SiO2 for the poorer stability.
The research on the temperature resistance of thecatalysts is very important for methane decomposition isan endothermic reaction and the conversion rises with theincrease of the reaction temperature [20]. Thus, a catalyst
with high temperature resistance can catalyze methane
decomposition more efficiently at a higher reaction
temperature [21]. According to the above results, a
conclusion can be drawn that the order of the catalytic
performance of the four kinds of catalysts above is
Ni-Cu-SiO2 (sol-gel method) > Ni-Fe-SiO2 (sol-gel
method) > 75%Ni-25%SiO2 (sol-gel method) >Ni-La2O3-SiO2 (co-impregnation method). Also, whenmethane decomposition over Ni-Cu-SiO2 or Ni-Fe-SiO2
catalysts is applied to TEP to recover deuterium andtritium from the deuterated and tritiated methane, a muchhigher reaction temperature can be adjusted to utilize thetemperature resistance of the catalysts to improve thecatalytic efficiency.

however, the lifetime of the catalyst was reduced and themethane conversion decreased to 2% after 300 minutes.
Therefore, a higher reaction temperature reduced thelifetime of the catalyst and another conclusion can also be
drawn that the temperature resistance of Ni-La2O3-SiO2
was not so good even if the introduction of La2O3 into
Ni-SiO2 greatly extended the lifetime of Ni-SiO2 at alower reaction temperature such as 823 K.
Fig.3 compared the catalytic performance of
Ni-La2O3-SiO2 prepared by co-impregnation method withthat of 75%Ni-25%SiO2 prepared by sol-gel method and
it demonstrated that the effect of the preparation methodof the catalysts on methane decomposition was important.
The catalytic activity of Ni-La2O3-SiO2 with La/Ni = 0.3
which showed the best catalytic performance among theNi-La2O3-SiO2 catalysts was a little lower than that of
75%Ni-25%SiO2. The methane conversion of
75%Ni-25%SiO2 catalyst at 823 K remained 21% anddidn’t change during the reaction. However, the methane
conversion of Ni-La2O3-SiO2 with La/Ni = 0.3 decreased
from 17% to 14% after 400 minutes. Therefore, both thecatalytic activity and the stability of 75%Ni-25%SiO2
were better than that of Ni-La2O3-SiO2 with La/Ni = 0.3
which indicated that the Ni-based catalysts prepared bysol-gel method showed better catalytic performance
than that prepared by co-impregnation method during
methane decomposition.
35

Ni-La2O3-SiO2(La/Ni=0.3)
75Ni-25SiO2

Methane conversion/%

30
25
20
15
10
5

0

50

100

150

200

250

300

350

400

Time/minFig.3 Effect of preparation method on methane decompositionat 823 K (m(Ni-La2O3-SiO2) = 60 mg, co-impregnation
method; m(75%Ni -25%SiO2) = 30 mg, sol-gel method).

Fig.4(a) depicts the kinetic curves of methane
decomposition over Ni-Cu-SiO2 catalysts with different
contents of Cu at 923 K. The lifetime of 75%Ni-25%SiO2
catalyst was very short at 923 K unlike the catalytic
behaviors at 823 K of which the methane conversion was
kept at 21% for 400 minutes (Fig.3). After 70 minutes,
the methane conversion of 75%Ni-25%SiO2 decreased toaround 5% and remained at this value. The addition oflower content of Cu improved the lifetime of Ni-SiO2
catalyst apparently and the catalytic activity of

3.2 The characterization of the catalysts before andafter methane decomposition
Fig.5(a) shows the XRD spectra of Ni-La2O3-SiO2

44

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

7070(a) m(catalysts)=30mg

75%Ni-25%SiO2
65%Ni-10%Cu-25%SiO2

60

Methane conversion/%

Methane conversion/%

35%Ni-40%Cu-25%SiO2
50
40
30
20
10
0

(b) m(catalysts)=50mg

75%Ni-25%SiO2
65%Ni-10%Fe-25%SiO2

60

35%Ni-40%Fe-25%SiO2

50
40
30
20
10

50

100

150

200

250

300

0

350

50

100

150

200

Time/min

Time/min250

Fig.4 Kinetic curves of methane decomposition over Ni-Cu-SiO2 and Ni-Fe-SiO2 catalysts at 923 K.

catalysts reduced by hydrogen. The diffraction peaks at
2T = 44.46°, 51.81° and 76.36° in the XRD spectra of
Ni-SiO2 revealed that the Ni species in the Ni-SiO2
catalyst reduced by hydrogen existed mainly in the formof Nimetal. When lower content of La2O3 was introduced
into the catalysts such as Ni-La2O3-SiO2 with La/Ni = 0.1
and La/Ni = 0.3, no La2O3 peaks were detected whichmeant that the lower contents of La2O3 existed in an
amorphous phase. However, as the contents of La2O3
exceeded La/Ni = 0.3, the La2O3 diffraction peaks
appeared. In addition, with the increase of the La2O3
content, the intensity of the diffraction peaks due to Ni
metal turned weaker and weaker and the FWMH of thediffraction peaks turned wider and wider which indicatedthat the Ni crystallite size turned smaller and smaller.
Therefore, the Ni dispersion degree was enhanced with
the increase of La2O3 content.
Fig. 5(b) shows the XRD spectra of Ni-SiO2 andNi-Cu-SiO2 catalysts reduced by hydrogen. The
diffraction peaks at 2T = 44.46°, 51.81° and 76.36° in the

XRD spectra of 75%Ni-25%SiO2 also demonstrated that
the Ni species in the Ni-SiO2 catalysts reduced by H2
existed mainly in the form of Ni metal. The diffraction
peaks due to Cu were not detected in Ni-Cu-SiO2
catalysts even if the Cu content was very high. In fact,
when Ni-Cu-SiO2 was reduced by hydrogen, the Ni
species was associated with the Cu species through theformation of Ni-Cu alloy in the catalysts and only Ni-Cu
alloy metal phase existed in the Ni-Cu-SiO2 catalyst.
Fig.5(c) shows the XRD spectra of Ni-SiO2 andNi-Fe-SiO2 catalysts reduced by hydrogen. The Fe in the
catalysts also appeared as Ni-Fe alloy which was similar
with Ni-Cu-SiO2 catalysts. However, the Ni-Fe-SiO2
catalysts reduced by hydrogen had two metal phases: theNi metal phase and the Ni-Fe alloy phase which wasdifferent from Ni-Cu-SiO2. The diffraction peaks due to
Ni-Fe alloy and Ni metal overlapped with each otherwhich indicated that the Ni-Fe alloy had identicalstructure with the Ni metal.

(b) Ni-Cu-SiO2(a): Ni-La2O3-SiO2(c)Ni-Fe-SiO2

65%Ni-10%Cu-25%SiO2

La/Ni=0.3
La/Ni=0.45

30

40

50

2T/ q

60

70

80

75%Ni-25%SiO2

65%Ni-10%Fe-25%SiO2

35%Ni-40%Cu-25%SiO2

La/Ni=0.6
20

Intensity/cps

intensity/cps

La/Ni=0.1

Intensity/cps

75%Ni-25%SiO2
La/Ni=0

20

30

40

50

2T q

60

70

80

35%Ni-40%Fe-25%SiO2
20

30

40

50

2T q

60

70

80

Fig.5 XRD spectra of the Ni-La2O3-SiO2, Ni-Cu-SiO2 and Ni-Fe-SiO2catalysts reduced by hydrogen(Ÿ-Ni, Ƒ-La2O3, Ŷ-Ni-Cu alloy, 佐-Ni-Fe alloy).

Fig.6 shows the SEM images of the Ni-SiO2,
Ni-La2O3-SiO2, Ni-Cu-SiO2 and Ni-Fe-SiO2 catalysts
after methane decomposition at different reaction
temperatures. It could be clearly observed that thesurfaces of all the catalysts were completely covered with
the interlaced carbon filaments and the structure of the
carbon filaments was seriously influenced by the reaction
temperature and the introduced Cu and Fe. The

45

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

catalysts showed better catalytic performance during
methane decomposition.
The carbon filaments formed over 75%Ni-25%SiO2
at 923 K turned shorter and thinner compared with thoseformed at 823 K. Thus, the increase of the reaction
temperature lowered the carbon capacity of
75%Ni-25%SiO2. When Cu was introduced into Ni-SiO2,
the carbon filaments turned much thicker and lesshomogeneous. The introduction of Fe into Ni-SiO2 alsochanged the fundamental structure of the carbon
filaments formed during methane decomposition at 923 K.
They were much longer and thinner.

fundamental structure of the carbon filaments formed
over Ni-La2O3-SiO2 with La/Ni = 0 and La/Ni = 0.3 at
823 K (Fig.6(a) and Fig.6(b)) was similar with that of
75%Ni-25%SiO2 at 823 K (Fig.6(c)) except the thickness
and the thickness of the carbon filaments formed over thethree catalysts was 75%Ni-25%SiO2 > Ni-La2O3-SiO2
with La/Ni = 0.3 > Ni-SiO2 with La/Ni = 0.
From the performance of methane catalytic
decomposition in Fig.1 and Fig.3, the order of thecatalytic performance of the three catalysts was
75%Ni-25%SiO2 > Ni-La2O3-SiO2 with La/Ni = 0.3 >Ni-SiO2 with La/Ni = 0. It could be concluded that whenthicker carbon filaments formed over the catalysts, theFig. 6 SEM images of carbon filaments formed during methane decomposition (a: Ni-SiO2 with La/Ni=0, 823e:K; b:Ni-La2O3-SiO2 with La/Ni=0.3, 823 K; c: 75%Ni-25%SiO2, 823 K; d:75%Ni-25%SiO2, 923 K;
65%Ni-10%Cu-25%SiO2, 923 K; f: 35%Ni-40%Fe-25%SiO2, 923 K).

Considering the great changes of the carbon
filaments with the increase of the reaction temperatureand the introduced Cu and Fe, TEM images of thesecarbon filaments were also investigated to analyze thedifferences of the detailed structure of these carbon
filaments.
In the TEM images of 75%Ni-25%SiO2 after
methane decomposition at 823 K (Fig.7(a)), the carbon
filaments with diameters from 70 nm to 110 nm wereobserved. The “pear-shaped” Ni metal particles were
present at the tip of the carbon filaments. While in the
TEM images of 75%Ni-25%SiO2 catalyst formed during
methane decomposition at 923 K (Fig.7(b)), the diameter
range of the carbon filaments was 50 nm-90 nm andspherical Ni metal particles appeared at the tip of the
carbon filaments. The diameters of both types of carbon

filaments were in agreement with their diameters of thecatalyst particles at the tips.
The different shapes of the Ni metal particles atdifferent reaction temperatures reflected the different
carbon diffusion mechanisms during methane
decomposition. Takenaka [22] et al pointed out that theformation of the “pear-shaped” Ni metal particles wasdue to the diffusion of carbon atoms through the bulk of
Ni metal particles, while the formation of the spherical Ni
metal particles resulted from the diffusion of carbon
atoms on the surface of Ni metal particles. Thus, it could
be speculated that the carbon atoms formed during
methane decomposition was preferential to diffuse
through the surface of the Ni metal particles at a higher
temperature.
Another phenomenon which should be paid attention

46

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

Fig.7(c) and Fig.7(d) show the TEM images of the
carbon filaments formed during methane decomposition
over 65%Ni-10%Cu-25%SiO2 at 923 K. It can beobserved that “bifurcate carbon filaments” formed over
65%Ni-10%Cu-25%SiO2 catalysts and the Ni-Cu alloy
particle which was active for methane decomposition was
not at the tip of the carbon filaments as it was in the
carbon filaments formed over Ni-SiO2 catalyst. A few
carbon filaments grew from one Ni-Cu alloy particle andthe Ni-Cu alloy particle was positioned among the carbon
filaments.

to was that when the formed carbons during methane
decomposition diffused through the bulk of the active
metal at 823 K, the catalyst could keep active for methane
decomposition for a long time, while, when the formedcarbons during methane decomposition diffused throughthe surface of the active metal, the catalyst lost its activity
in a short time (Fig.3 and Fig.4). As a result, the diffusionof the formed carbons through the surface of the catalysts
wasn’t beneficial for the long time methane
decomposition.

Fig.7 TEM images of carbon filaments formed over Ni-SiO2, Ni-Cu-SiO2 and Ni-Fe-SiO2 during methane
decomposition (a: 75%Ni-25%SiO2,823 K; b:75%Ni-25%SiO2, 923 K; c,d: 65%Ni-10%Cu-25%SiO2, 923 K;e,f:35%Ni-40%Fe-25%SiO2, 923 K).

Compared with the Ni particles in Fig.7(b) which
was formed over 75%Ni-25%SiO2 after methane
decomposition at 823 K, the Ni-Cu particles were also
“pear-shaped” after methane decomposition overNi-Cu-SiO2 catalysts at 923 K. Thus, it could bespeculated that the addition of Cu into Ni-SiO2 catalystchanged the diffusion mechanism of the formed carbons.
At a higher reaction temperature such as 923 K, thediffusion of the carbons formed over Ni-Cu-SiO2 catalyst
remained through the bulk of Ni-Cu particles. This might
be the reason why the lifetime of Ni-Cu-SiO2 catalysts
was also very long at a higher reaction temperature suchas 923 K.
Fig.7(e) and Fig.7(f) shows the TEM images of the
carbon filaments formed over Ni-Fe-SiO2 catalysts at 923
K. Carbon filaments of “bamboo-shaped” structure
formed over Ni-Fe-SiO2 catalyst. The walls of the“bamboo-shaped carbon filaments” were uneven and thehollows of these carbon filaments were divided into manycells which was different from those formed over Ni-SiO2
catalyst. The Ni-Fe alloy particles which were active formethane decomposition were not only present at the tip ofthe carbon filaments after methane decomposition, butalso existed in the hollow cells of the “bamboo-shaped”
structure.
Apparently, the addition of Fe also changed thecarbon diffusion mechanism of Ni-SiO2 catalyst during
methane decomposition. The results of Ermakova and
Takenaka [18, 23] et al indicated that the products ofmethane decomposition over Fe-SiO2 catalysts were also“bamboo-shaped carbon filaments”. Thus, the carbon
diffusion mechanism of Ni-Fe-SiO2 catalyst during
methane decomposition was similar with that of Fe-SiO2
catalyst.
Since the carbon diffusion mechanism of Ni-Fe-SiO2
is similar with that of Fe-SiO2 catalyst, it can also be

47

H. Wang et al., Effects of La2O3, Cu and Fe Addition on the Catalytic Performance of Ni-SiO2 Catalysts for Methane Decomposition

Acknowledgements

believed that the characterization of methane
decomposition over Ni-Fe-SiO2 catalysts should be inagreement with that of Fe-based catalysts. As theactivation temperature of the Fe-based catalysts is muchhigher than that of the Ni-based catalysts [23], theintroduction of Fe also enhanced the temperature
resistance of Ni-based catalysts and the lifetime of Ni-Fe
alloy catalysts was much longer than that of the single
Ni-based catalysts at a higher reaction temperature suchas 923 K.
In addition, the temperature resistance of Ni-Cu
alloy and Ni-Fe alloy catalysts was both stronger thanthat of the single Ni-based catalysts, while the stability of
65%Ni-10%Fe-25%SiO2 catalyst during methane
decomposition wasn’t as good as that of
65%Ni-10%Cu-25%SiO2 (Fig.4). The reason for thismight be that not all of the Ni species in
65%Ni-10%Fe-25%SiO2 catalyst existed in the form of
Ni-Fe alloy (Fig.5), while Ni and Cu could form alloys
over a wide composition range and this was proved by
the results of Fig.5 and other researchers [24]. When the
Fe content in Ni-Fe-SiO2 increased to 40%, the stabilityof the catalyst was enhanced for the Ni species in
35%Ni-40%Fe-25%SiO2 mainly existed in the form of
Ni-Fe alloy although the catalytic activity of the catalyst
decreased.

This work was ¿nancially supported by the National
Magnetic Confinement Fusion Science Program of China(2011GB111005).

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4. Summary

1) The introduction of La2O3 into Ni-SiO2 catalysts
obviously extended the lifetime of the catalysts during
methane decomposition at 823 K. However, thetemperature resistance of Ni-La2O3-SiO2 didn’t change
obviously and the catalytic activity of Ni-La2O3-SiO2
with La/Ni = 0.3 disappeared at a very short time at 873
K. While the introduction of Cu and Fe into Ni-SiO2
enhanced the temperature resistance of the catalyst. The
catalytic performance of the catalysts during methane
decomposition was Ni-Cu-SiO2 (sol-gel method) >Ni-Fe-SiO2 (sol-gel method) > 75%Ni-25%SiO2 (sol-gel
method)> Ni-La2O3-SiO2 (co-impregnation method).
2) XRD studies of Ni-La2O3-SiO2, Ni-Cu-SiO2 andNi-Fe-SiO2 catalysts indicated that the Ni species existed
in the form of Ni metal in Ni-La2O3-SiO2 catalyst andNi-Cu alloy and Ni-Fe alloy appeared in Ni-Cu-SiO2 andNi-Fe-SiO2 catalysts, respectively.
3) The fundamental structure of the carbon filaments
formed over Ni-La2O3-SiO2 with La/Ni = 0 and La/Ni =0.3 was similar with that of 75%Ni-25%SiO2 at 823 K.
While the carbon filaments formed over 75%Ni-25%SiO2
at 923 K turned shorter and thinner than those formed at
823 K. “Bifurcate carbon filaments” formed overNi-Cu-SiO2, while “bamboo-shaped carbon filaments”
formed over Ni-Fe-SiO2 catalysts during methane
decomposition at 923 K.

48

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Measurement of Tritium Distribution in Nickel and Vanadium by Means of
A Combined Technique of An Imaging Plate and Thin Absorbers
Hiroko YOSHIDA-OHUCHI, Yuji HATANO1), Akram MOHAMMADI2) and Takao KAWANO3)
Radioisotope Research and Education Center, Graduate School of Pharmaceutical Science,
Tohoku University, Sendai 980-8578, Japan

Hydrogen Isotope Research Center, University of Toyama, Toyama 930-8555, Japan

1)

Radiation Research Center, Shiraz University, Shiraz 85111-71946, Iran2)

3)

National Institute for Fusion Science, Toki-city 509-5292, Japan(Received: 10 May 2012 / Accepted: 21 October 2012)

A combined technique of an imaging plate (IP) and thin absorbers was applied to tritium in nickel and vanadium
specimens using copper, aluminum, and gold foil as the absorber. Copper and aluminum foil are used as a K-edge
filter with X-ray absorption at 9.0 keV and 1.56 keV, respectively. Gold has L-edges X-ray absorption at around 13
keV. With this technique, photostimulated luminescence (PSL) decay curves are obtained by changing absorber’s
thickness. In the nickel specimen, the difference in PSL decay curves between for the copper and gold absorber was
clearly observed 20 days after loading, however, all curves became similar single pattern after 388 days. The same
curve pattern was obtained in vanadium for all absorbers. The cross section images and depth profiles, which weretaken at 468 days and 3.9 years after loading for the nickel and vanadium specimen, respectively, show no
significant inclination of tritium concentration for both specimens. These results indicate that uniform tritium
distribution in the specimen provides the similar single PSL decay curves pattern for the copper or aluminum andgold absorber.
Keywords: tritium, imaging plate, thin absorber, bremsstrahlung X-rays, depth distribution, non-destructive     measurement
bremsstrahlung X-rays generated by the interaction
between the beta particles from tritium and matter [7].
In this study, the combined technique of the IP andthin absorbers was applied to tritium in nickel andvanadium specimens using copper, gold, and aluminum
foil as the absorber. Copper and aluminum foil are used asa K-edge filter with X-ray absorption at 9.0 keV and 1.56
keV, respectively. Gold has L-edges X-ray absorption ataround 13 keV [7].

1. Introduction
Non-destructive and quantitative measurement of theamount of tritium retained on/in plasma-facing materials(PFMs) of magnetic fusion devices are of great importancefor the control of fuel particles and to ensure safety for
maintenance work and waste processing of PFMs in fusion
systems. The tritium imaging plate (IP) technique fordetecting beta particles is useful to non-destructively
determine the surface tritium distribution on PFMs [1-5].
An IP made of europium-doped BaFBr(I), a
photostimulated luminescence (PSL) material, is a
two-dimensional radiation sensor. This IP has many
excellent properties making it suitable for this purpose,including high sensitivity, wide dynamic range over fiveorders of magnitude, a high degree of spatial resolution,and reusability by exposing the IP to visible light between
uses [6].
Tritium emits beta particles with a maximum energyof 18.6 keV, with an average of 5.7 keV. These have arange of about several micrometers in graphite, and hence
the IP technique is sensitive to tritium up to a depth of afew microns. In order to obtain tritium distribution inregions deeper than the escape depth of beta rays fromtritium, we have been developing a combined technique ofan imaging plate (IP) and thin absorbers using the

2. Experimental
Tritium was introduced into sheet type specimens (1.5
x 1.5 x 0.05 cm3) of nickel by gas absorption method at
673 K on 24 Dec 2010 in the manner described in theprevious paper [8]. Specimens were loaded with deuterium
containing 17% tritium. For vanadium specimen (4.8 x 1.0
x 0.05 cm3), tritium was introduced by the same method at
673 K on 27 May 2008 [9]. The tritium concentration in
the specimen was controlled to be 1.0 appm; the total
concentration of tritium and deuterium was 200 appm.
Then all specimens were kept at room temperature.
The initial distribution of tritium in the nickel
specimen was examined by β-ray induced X-ray
spectroscopy (BIXS) under Ar atmosphere in the manner
described in the previous paper [10]. Tritium distribution

author’s e-mail: hiroko@m.tohoku.ac.jp

49©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

H. Yoshida-Ohuchi et al., Measurement of Tritium Distribution in Nickel and Vanadium by Means of a Combined Technique of an Imaging Plate and Thin Absorbers

3. Results and Discussion

was measured twice, 20 and 388 days after tritium loading
for the nickel specimen and 3.5 years for the vanadium
specimen, respectively by the combined technique of a
BAS-MS type-IP (Fujifilm Co., Ltd.) and absorbers. At thesame time, tritium in the surface regions was measured byusing a BAS-TR type-IP (Fujifilm Co., Ltd.). The
BAS-MS type-IP was used to detect bremsstrahlung
X-rays from tritium. This IP consists of a 9-μm-thickpolyethyleneterephthalate protective film and a
115-μm-thick photostimulable phosphor layer affixed to a
12-μm-thick plastic back layer and a 190-μm-thickpolyethyleneterephthalate base layer [11]. It has a highsensitivity to photons. The BAS-MS type-IP can detect
bremsstrahlung X-rays, but not the beta particles emitted
from tritium because of its thick protective film. The
BAS-TR type-IP lacking a protective surface layer wasdeveloped for detection of the beta particles from tritium.
This IP consists of a 50-μm-thick photostimulable
phosphor layer affixed to a 250-μm-thick base layer.
The BAS-MS type-IP was irradiated withbremsstrahlung X-rays by placing the specimen directly on
the IP or with inserting absorbers for 1 h. To avoid
contamination of the IP during irradiation, each specimen
was sandwiched with a quite thin 1.2-μm-thick film. The
schematic diagram of the irradiation setup is shown in Fig.
1. During irradiation, the IP was kept inside a refrigerator
with temperature controlled to 4°C, in order to decrease thefading effect, although the BAS-MS-type IP shows goodfading characteristics at room temperature [12]. The IP was
kept for 20 min after irradiation and then read to avoid
fading effects caused by the short half-life component. A
model FLA7000 IP reader (Fujifilm Co., Ltd.) was usedwith a spatial resolution of 100 μm and 8 bits of digital
resolution dpi, followed by an image analysis with theMulti Gauge (ver.3.2). For each irradiation, the IP was
simultaneously irradiated with a 137Cs disk-shaped
gamma-ray source and measured PSL values obtained
from the specimen were corrected by that obtained by the137
Cs source to minimize errors induced by the readout
system and the procedure. To detect tritium in the surface
regions, the BAS-TR type-IP was irradiated with the beta
particles by placing the specimen directly on the IP without
using the film for 20 min.

Fig.1

Fig.2 shows X-ray spectrum from the nickel specimen
containing tritium. A sharp intense peak of Ar Kα X-rays
is at around 3 keV. Another sharp peak at 7.48 keV can be
assigned to Ni Kα X-rays. It should be noted that relatively
small intensity of Ni Kα peak in comparison with Ar Kα
peak. Ar Kα X-rays are induced by tritium in shallow
surface region, while Ni Kα X-rays are induced mainly bytritium in the bulk. The spectrum in Fig.2 indicates that
tritium is concentrated in the surface regions, which isconsistent with the result shown in the previous paper [8].

Ar Kα

Ni Kα
N

Fig.2 Characteristic BIXS spectrum of a tritiumcontaining nickel specimen.
Fig.3(a) shows a BAS-TR image for the nickel
specimen taken at 20 days after loading, describing tritium
two-dimensional distribution in the surface region. The
PSL intensity along the line A and B are shown in Fig.3(b),respectively, expressing flat distributions along both lines
at the same PSL intensity level except the edge. The
(a)

(b)

Fig.3 (a) BAS-TR image for the nickel specimen taken at
20 days after loading. (b) PSL intensity along
the line A and B.

Schematic diagram of the irradiation setup. 
  

50

H. Yoshida-Ohuchi et al., Measurement of Tritium Distribution in Nickel and Vanadium by Means of a Combined Technique of an Imaging Plate and Thin Absorbers(g/cm3), and t the thickness of absorbers (cm). The scaleof the effective thickness for aluminum absorber isexpressed as 1/10 in Fig.4(b). When X-rays energy
spectrum distribution spreads over some energies, theslope of PSL decay curve is expected to change as the
thickness of the absorber increases.
PSL decay curves pattern for copper and aluminumabsorbers (shown as 佑, 伽, respectively) are similar, whilethat for gold (shown as ‫��ڧ‬is quite different fromothers. In the previous paper [7], we described theprinciple of the combined technique of the IP andabsorbers and how each absorber works when thebremsstrahlung energy spectrum distribution changes astritium migrates into matter, that is, PSL values derived
from the lower energy range rapidly decrease as thecopper and aluminum absorber thickness increase. On theother hand, the gold absorber works oppositely, rapidly
reducing PSL derived from the higher energy range. Thus,
PSL decay curve consists of two components, with fastand slow rates of decay, and they can be expressed in thefollowing equation [7]:f (x) =a · eíμ1x +b · eíμ2x          (2)where a and b are the relative ratio of two components, μ1
and μ2 penetrating coefficient of each component, and x
is the thickness of the absorber. Two components of PSL
decay curves for copper and gold absorbers are shown as
solid lines in Fig.4(b). The differences in relative ratio of
two components and penetrating coefficient of eachcomponent between two PSL decay patterns for copperand gold absorbers are summarized in Table 1. The
relative ratio of two components, a and b is considered toindicate the ratio of higher and lower energy
bremsstrahlung X-rays in the bremsstrahlung energy
spectrum distribution [7].

average PSL intensity per arbitrary unit is 8.4 on the flat
curves. The image taken at 388 days after loading wasexact the same one with that at 20 days in Fig.3(a) but witha decrease of 16.3% in the average PSL intensity.
Fig.4(a) shows BAS-MS images for the nickel
specimen taken at 20 days after loading when the thicknessof each absorber varied, exhibiting quite different images
compared with the BAS-TR image in Fig.3(a). Each PSL
value in the area of 85 mm2 at the center of the specimen
was obtained, and then PSL decay curve was obtained bynormalizing the measured PSL values to the PSL value
without absorber. Fig.4(b) shows PSL decay curves as afunction of the effective thickness for three absorbers. The
effective thickness, μ·t, which normalizes the thickness ofeach absorber by 1/ȝ, was obtained by the following
equation:  
μ·t = (μ/Ȩ) ·Ȩ· t       (1)where μ/Ȩis the mass attenuation coefficient (cm2/g) at
K or L-edge X-ray absorption for copper and aluminumabsorbers and gold absorber, respectively [13],Ȩdensity(a)(b)‫ڧڹۍ‬:Cu, Al, Au at
20 days

Fig.4

Table 1 Differences in relative ratio of two components, aand b and penetrating coefficient of eachcomponent, μ1 and μ2 between two PSL decay
patterns for copper and gold absorbers.
Absorber
a
b
w
w
Copper
0.98
0.020
4.0
0.15
Gold
0.90
0.10
1.7
0.19
The same measurements were repeated at 388 days
after loading. BAS-MS images show a large decrease of
98% in PSL intensity in comparison with that at 20 days,
which can be caused by chronic release from nickel underambient air at room temperature [8]. PSL decay curves asa function of the effective thickness for three absorbers
taken at 388 days are shown in Fig.5(a) and all PSL decay
curves at 20 and 388 days after loading are summarized
in Fig.5(b). It should be noted that Fig.5(a) shows nodifferences in the PSL decay curves pattern for threeabsorbers and these curves became close and similar to the(a) BAS-MS type-IP images for the nickel
specimen taken at 20 days after loading when the
thickness of each absorber varied. (b) PSL decay
curves as a function of the effective thickness forthree absorbers. The scale of the effective thickness
for aluminum absorber is expressed as 1/10. PSL
decay curves for copper and aluminum absorbers
are shown as 佑, 伽, respectively and that for gold isshown as ‫ڧ‬Solid lines express two componentsof PSL decay curves for copper and gold
absorbers .

51

H. Yoshida-Ohuchi et al., Measurement of Tritium Distribution in Nickel and Vanadium by Means of a Combined Technique of an Imaging Plate and Thin Absorbers

curve for the gold absorber at 20 days in Fig.5(b). Thechange in the PSL decay curves for copper and aluminumabsorbers between 20 and 388 days can be explained by
the diffusion process. Diffusion coefficient (D) at 293 K is
calculated for tritium through nickel to be 2.5 x 10-14 (m2/s)
using literature values [14]. At 20 days after loading,tritium concentration is considered to distribute with agradient as Dt/l2 is calculated to be 0.3 (l=0.25 mm) andthen a diffusion process ends at 388 days as Dt/l2 is
calculated to be 13 (>> 1.0) [15]. Based on the
preliminary results obtained by simulating thebremsstrahlung energy spectrum with Monte Carlo
N-Particle Transport Code (MCNP), the spectrum
distribution spreads over lower to higher range at 20 days
because tritium concentration distributes with the gradient,and the spectrum distribution changes with time, shiftingits peak to higher energy with a decrease of lower energy
bremsstrahlung X-rays because tritium concentration

distribution loses the gradient as the diffusion processproceeds and the uniform tritium distribution is formed in
the specimen after 388 days. The relative ratio, a, thecomponent with fast rate of decay for the copper absorberat 388 days was determined as 0.72, which became smaller
than that at 20 days (0.98 in Table 1), indicating PSL
derived from the lower energy range decreased with time.
This result is consistent with the assumed theory. On thecontrary, the PSL decay curve for the gold absorber does
not seem to have changed with time. This is because PSL
derived from the higher energy range does not decrease
through the diffusion process. Then, at a result, all PSL
decay curves for three absorbers are considered to become
similar at 388 days.
The combined technique of the IP and thin absorbers
was applied to tritium in the vanadium specimen and PSL
decay curves as a function of the effective thickness forthree absorbers were obtained (shown in Fig.6),
expressing the same pattern with those in the nickel
specimen for three absorbers at 388 days.

(a), ,: Cu, Al, Au
at 388 days

(b)

Fig.5

,:

:,Cu, Al, Au

Cu, Al, Au

at 20 days

at 388 days

Fig.6

PSL decay curves as a function of the effective
thickness for the vanadium specimen. The scale ofthe effective thickness for aluminum absorber isexpressed as 1/10.

In order to obtain tritium distributions in the nickel
and vanadium specimen, cross section images and depth
profiles were obtained by autoradiography. They weretaken at 468 days and 3.9 years after loading for the nickel
and vanadium specimen, respectively. Both specimens
were cut through the middle and the BAS-TR type-IP was
irradiated by contacting the cross section of the specimen
directly on the IP for 47 h. To avoid high concentrated
tritium on the nickel surface, the side of the cross section
was covered with tape during irradiation. The cross section
images and depth profiles so obtained are shown in
Fig.7(a) and (b) for the nickel and vanadium specimen,
respectively. Through the cross section, no significant
inclination of tritium concentration was observed in both
images and depth profiles. The observation for the nickel
specimen in Fig.7(a) can be explained by the diffusion(a) PSL decay curves for the nickel specimen as afunction of the effective thickness for threeabsorbers taken at 388 days after loading. (b) All
PSL decay curves at 20 and 388 days after loading.
Curves for copper, aluminum, and gold absorbers
are shown as ♦,▲, and at 388 days and ◇, △,and  at 20 days, respectively. The scale of theeffective thickness for aluminum absorber isexpressed as 1/10 in both figures.

52

H. Yoshida-Ohuchi et al., Measurement of Tritium Distribution in Nickel and Vanadium by Means of a Combined Technique of an Imaging Plate and Thin Absorbers

tritium distribution in the specimen and PSL decay curves
pattern by simulating the bremsstrahlung energy spectrumwith MCNP Code.

process. The process ended at 388 days as mentioned
above and then tritium distribution is supposed to beuniform in the specimen except high concentrated tritiumon the surface. For the vanadium specimen, the diffusivityof tritium is much larger than that in the nickel [9] andthe diffusion process ended at the time whenautoradiography was obtained, resulting in uniform tritium
distribution in the vanadium specimen. Theseobservations for two specimens indicate that uniform
tritium distribution in the specimen provides the similar
single PSL decay curves pattern for the copper oraluminum and gold absorber.
From comprehensive results, this technique can beused usefully to observe tritium depth distribution inmatter nondestructively. Further investigation should beneeded for the analysis about the relationship between(a)Acknowledgments
This study has been supported in part by KAKENHI
on Priority Areas, 476, Tritium for Fusion, from MEXT,
Japan, and the National Institute for Fusion Science under
the Joint Work contract #NIFS12KUHR017, as a part of
Bilateral Collaboration Project with University of
Toyama. The author is grateful to Professor Yasuhiro
Kondo of Ishinomaki Senshu University for his advice.

T. Tanabe et al., Fus. Eng. and Des. 54, 147 (2001).
K. Miyasaka et al., J. Nucl. Mater. 307-311, 1441 (2002).
K. Masaki et al., J. Nucl. Mater. 313-316, 514 (2003).
T. Tanabe et al., J. Nucl. Mater. 313-316, 478 (2003).
K. Sugiyama et al., J. Nucl. Mater. 313-316, 507 (2003).
J. Miyahara et al., Nucl. Instr. and. Meth. A 246, 572(1986) .
[7] H. Ohuchi-Yoshida et al., Fus. Eng. and Des. 87, 423(2012).
[8] M. Saito et al., Fus. Sci. Technol. 60, 1459 (2011)[9] Y.Hatano et al., Proceedings of 2nd Japan-China Workshop
on Blanket and Tritium Technology 100 (2008)[10] Y.Torikai et al., J. Nucl. Mater. 363-365, 462 (2007)[11] Fuji Film Co., Ltd. Homepage,
http://www.fujifilm.com/products/life_science_systems/scie
nce_imaging/imaging_plate/faq/[12] H. Ohuchi et al., Nucl. Instr. and. Meth. A 490, 573(2002).
[13] NIST Homepage,
http://physics.nist.gov/PhysRefData/XrayMassCoef/tab3.html.
[14] F. Reiter et al., EUR 15217 EN (1993).
[15] J.Crank, The mathematics on diffusion 2nd edition(Oxford university press, New York, 1975) p.50.
[1]
[2]
[3]
[4]
[5]
[6]

(b)

Fig.7

BAS-TR autoradiography cross section images anddepth profiles for the nickel (a) and vanadium (b)specimen, respectively. In each figure, the average
profile obtained from 10 profiles is shown in a
line.

53

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

The R&D Status of ITER SDS
Sei-Hun YUN1)*, Chang Shuk KIM1), Min Ho CHANG1), Hyun-Goo KANG1), Seungyon CHO1),
Hyun Gon LEE1), Ki Jung JUNG1), Hongsuk CHUNG2), Dae Seo KOO2), Kyu-Min SONG3),
Soon-Hwan SHON3) and Duk Jin KIM4)
1)ITER KO, National Fusion Research Institute, Gwahangno 113, Daejeon, 205-806, Korea

2)
3)

Korea Atomic Energy Research Institute, Daedeokdaero 1045, Daejeon, 305-353, Korea

Korea Hydro & Nuclear Power Co., Central Research Institute, 70-1312-gil, Yeseong-daero, Daejeon, 305-343, Korea
4)

TUV SUD KOCEN, 13F Kranztechno, Sangdaewon-dong, Jungwon-gu, Seongnam, 462-729, Korea(Received: 27 May 2012 / Accepted: 24 August 2012)

The ITER tritium storage and delivery system (SDS) is still in pre-conceptual phase. The major functions of
ITER SDS are composed of tritium storage in a safe way, rapid fueling based on the plasma scenario and tritium
inventory accountancy in a day schedule. The ITER SDS will be delivered to the ITER site as a set of glovebox
modules which have every tritium compatible components in its space for tritium confinement and are to beassembled together at the ITER tritium building. Meanwhile there have been some reports for the ITER SDS’ R&D
activities and progress, the current R&D status is still on-going at the pre-conceptual design (CD) phase. The scope ofcurrent R&D has been rearranged with the depleted uranium (DU) oriented experiments for the future ITER
application. This report deals with the current status and activities, especially setup of handling procedure, test devices
and facility design based on DU getter material, for CD completion of the ITER SDS next year.
Keywords: ITER SDS, R&D status, getter material, ZrCo, DU, process verification, 1:1 scale test.

SDS in tritium plant fuel cycle as well. The R&Ds on this
ITER SDS lies in the starting point of newly programmed
systematic R&D phase. In this period the referencematerial for ITER SDS has been changed from zirconium
cobalt (ZrCo) to depleted uranium (DU), because ofdifferent shortfalls in material properties and limits intechnical applications [4].
There have been some reports for the SDS R&D
activities and progress [1-3]. The content in the first phase,pre-conceptual R&D was focused on ZrCo as the tritium
storage material; R&Ds for type and design concept of thetritium storage bed and characteristics of storage material
were reported in Ref. [1]; R&D’s for bed performance interms of delivery rate, thermo hydraulic analysis of in-bed
calorimetry, studies on hydride reaction using visual cell
reactor, unit process verification test having 1:1 scaled
ZrCo bed, transfer pump performance and its combination,and calorimeter design for the tritium accountancy at thetritium loading station (TLS) were reported in Ref. [2]; andR&D’s for highly concentrated and enhanced results on
every ITER SDS’ component and interrelationship werealso reported in Ref. [3].
From the turning point of the tritium storage material,
DU, this report deals with the current status and activities

1. Introduction
ITER tritium storage and delivery system is one of
key technologies, which are to be utilized for the fusion
tritium plant construction, to provide the fusion fuelingsystem by means of tritium storage, supply and inventory,
in a safe way, in accordance with the plasma scenario andaccountancy availability in a day basis, respectively. While
the SDS is composed of major equipment needed further
R&D for detailed design, it will be followed by thesystematic R&D program. The main objective of the ITERSDS R&D—even though still on-going—is to verify not
only each component’s requirements in accordance withthe system function, but also performance availabilities as
the whole of the SDS’ requirements physically andtechnically combined with each other components thathave individual characteristics and performance in most
aspects of tritium safety and compatibility. The overall
project schedule of the ITER SDS is divided into three
distinguishable phases: (1) pre-conceptual R&D phase, (ii)systematically programmed R&D phase (conceptual
design phase, preliminary design phase), (iii) and the final
design & procurement phase. This report is focused on thecurrent R&D status in pre-conceptual phase, in addition to
the further consideration on the interface relatives to the54©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

S.-H. Yun et al., The R&D Status of ITER SDS

their performance. The rig is used for a measurement of
the hydrogen recovery and delivery rates. A hydrogendelivery scroll pump is connected to the manifold piping.
The scroll pump has a desirable pumping characteristic.
Its leakage rate is negligibly small. But the leakage rate ismeasured every time before each hydrogen delivery
experiment. Because of its leakage characteristics thispump is used only for an experimental purpose. Metal
powders in two bed models are activated through several
repetitions of hydriding and dehydriding. The amount ofhydrogen used for the initial activation and hydriding &dehydriding runs is measured by the hydrogen pressure
filled in the loading vessel and the manifold. The control
and data acquisition system is provided.
Fig. 3 shows a temperature transient and heating
pattern of the primary vessel during the heating test. The
measured and setting values of the primary vesseltemperature are almost the same. Overshooting of theheater exists but is negligible. Fig. 4 shows thetemperature transient of the primary vessel during thecooling test [6]. Secondary vessel was either undervacuum (case A in Fig. 4) or filled with helium (case B in
Fig. 4). A series of hydriding and dehydriding tests for the
ZrCo and DU beds is going to be performed further.

mainly issued topics, especially DU handling procedures,experimental devices and facility design, in R&D of ITER
SDS.

2. Small Scale Bed Experiments and Basic Study
2.1 DU bed preparation & experiment

Various SDS beds including beds using DU has beendeveloping [5]. Fig. 1 shows a metal hydride bed in the
glovebox. The argon in the glovebox is pumped through aseries of treatment devices, which remove water andoxygen from the gas. Finely divided heated copper metal
is used to remove oxygen, and this oxygen removing
column is regenerated by passing 4% hydrogen in an
argon mixture through it while it is heated, and the water
formed is passed out of the box with the excess hydrogenand argon. A 13X molecular sieve is used to remove
water by adsorbing it in the pores of the molecular sieves.
The hermetically sealed glovebox maintains an argon
atmosphere with an H2O concentration at 0.0 ppm, andO2 concentration at about 0.0–0.4 ppm. The glovebox is
kept at a higher pressure of about 1.25-in. H2O than thesurrounding air, so that any microscopic leaks are mostly
leaking inert gas out of the box instead of letting air in.
The metal hydride bed is introduced into the mainchamber through a sealed antechamber on the right sideof the glovebox. In the glovebox the Swagelok nut on the
primary vessel is open to introduce the depleted uranium
into the primary vessel of the metal hydride bed. The
quantity of the depleted uranium has already been exactly
measured by using a precision balance in the glovebox.
The primary vessel that contains depleted uranium sealedby the Swagelok nut is withdrawn from the glovebox
through the antechamber. Finally the primary vessel isassembled with the secondary vessel of the metal hydridebed.

2.2 Visual DU hydriding reaction

The DU hydride is going to be analyzed by using avisual cell reactor as shown in Fig. 5 and thermo physicalproperty analyzers. This kind of experimental apparatus
was utilized in ZrCo hydriding system before [7-8].
The visual cell apparatus in the right-hand side ofFig. 5 was used to look at the hydride material’s behavior
at hydriding/dehydriding reaction and the compressibilityof the expanded hydride micro-powder [7-8]. Table 1
shows the summary of experimental thermo physicalpropertiesȢpacking density, heat capacity, thermal
diffusivity and thermal conductivityȢof ZrCo hydridematerial [8]. This result was also applied to thetheoretical analysis and simulation of heat & masstransfer phenomena in the complicated getter bed systems,
and engineering design as well.
The DU hydride material is to be tested using similar
apparatus (left-hand side of Fig. 5). However, the maindifference of DU hydriding reaction system including avisual cell reactor comparing to ZrCo hydriding case is to
apply an inert gas atmosphere using a glovebox thatcovers all equipment for the containment of air andmoisture.

Fig.1 Hydride bed in the glovebox [5].

2.3 Related engineering studies

Fig. 2(a), 2(b) show the equipped system picture andflow diagram of the metal hydride bed performance test
rig, respectively [5]. The bed performance test rig canequip two different beds at the same time to compare

Fig. 6 shows the detailed mathematical approach toanalyze and simulate the getter bed system that iscomposed of thin double-layered, annulus-type, ZrCo bed

55

S.-H. Yun et al., The R&D Status of ITER SDS

gross to cool down. The cooling time, therefore, of theheated getter bed after desorption requires much more
hours to get ready for the next cycle at absorption process.
The next getter bed model should be limited to minimize
the thermal mass and to maximize the heat and masstransfer rate for the ITER SDS optimization, theoretically.

to be anticipated getting a rapid absorption/desorption
behavior than any other geometric type of getter bed. In
comparison of thin annulus bed with the typicalcylindrical bed that has same lengthy in Fig. 6 theabsorption and desorption rate of the thin annulus bed
shows remarkable speed in more than two times. On theother hand, the thermal mass of the getter bed was too(a) Equipped test rig of (b)(b) Flow diagram of ZrCo and DU bed system
Fig.2 Hydride bed performance test rig [5].

56

S.-H. Yun et al., The R&D Status of ITER SDS

Table 1 Thermo physical properties of ZrCoHx [8].

Fig.3 Heating test of the primary vessel [6].

Fig. 4 Cooling test of the bed [6].

Fig. 5 Metal hydriding apparatus and a visual cell reactor.

57

S.-H. Yun et al., The R&D Status of ITER SDS(1) Mesh Configuration(2) Geometric details and mesh reconfiguration(3) Illustration of temperature and H/M ratio evolution profiles.
Fig. 6 Theoretical approach and simulation result of the annulus-type ZrCo bed (upper), andcomparing absorption & desorption rate with cylindrical bed (down).

58

S.-H. Yun et al., The R&D Status of ITER SDS

At one fuel supplying scenario a group of getter beds
has to be operated simultaneously to satisfy the ITER
fueling requirement. Here, another complicated problem
might be occurred, such as control logic development forthe optimization of getter bed system operation. Fig. 8
shows a similarity test apparatus using multiple batch-wise
efflux system that seems to behave the multiple getter bed
system: Efflux/desorption, recharging/ absorption, andsuccessive cyclic operations.

Fig. 7 shows an illustration of ITER SDS module forthe tritium building installation. In general, SDS moduleand every component in the module glovebox keep thefire-sector rule (The maximum tritium content in every
component is limited under 70 gram, and in every
container under 700 gram of tritium). ITER SDS module
is composed of multiple getter beds, buffer vessels,transfer pumps and many instruments as shown in Fig. 7.
Tens of gloveboxes are to be assembled to compose thewhole ITER SDSĜs functions and requirements.

Fig. 7 Illustration of glovebox inside in SDS module (4~6 modules).

1 Series of Tanks
2 Buffer Tank
3 Reservoir Tank
4 Frame5 Transfer Pump
6 Water Tank (Bed)
7 Level Indicator

Fig. 8 Similarity test apparatus using multiple batch-wise efflux flow system instead of getter bed system.
(Number means a detailed component in this system.)

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S.-H. Yun et al., The R&D Status of ITER SDS

2.4 He-3 recovery process
Recovery or separation of He-3 from the tritium
storage system is one of major function in ITER SDS. Fig.
9 shows the schematic drawing of He-3 recovery systemfrom the tritium storage getter bed as a basis of ITER
SDS. Since the release of He-3 from the tritium absorbed
in uranium is known to start at about 300 days age, theactual recovery of He-3 in ITER SDS’ tritium storage
does not seem to be realized. Moreover, the blanket effectis also well known phenomena in recovery of He-3 byconcentration of tritium. However, the cyclic circulationof the He-3 mixture is recommended to increase thetritium concentration at the metal getter in Fig. 9.
Thereafter, the initial He-3 blanket effect is expected tobe diminished in gradual operation. There has been noquantitative result for the cyclic operation of this ITER
concept (this concept will be studied also).

gases, therefore, the beds must be cooled down in a fast
way. If the heating and cooling time of bed is too long,
the turnover is delayed and results in enormous increaseof SDS beds.
  As a part of preparatory, systematic R&D work forthe development of SDS bed, experimental study on
various cooling of metal hydride bed is to be performed.
A revised mock-up of metal hydride bed having thecapacity of 70 g tritium storage is developed. ZrCo is
used as a dummy material for the getter. The thermal
mass is reduced and several active cooling methods areapplied to the revised bed. Experimental apparatus for the
independent circulation of several cooling circuit isinstalled as shown in Fig. 10 and Fig. 11. Fig. 10 shows
three active cooling circulations. The first is circulation ofhelium loop which is inserted in the metal hydride layer
for the in-bed calorimetry. The second is gas circulation
inside outer jacket of the bed. And the last is watercooling jacket which is attached on the surface of outer
jacket. Even though the direct water cooling of the getterbed for the purpose of safe tritium storage in priority is tobe evaded in feelings, otherwise an indirect coolingshould be necessary to maintain the constant glovebox
temperature by removing the heat accumulation of heatedgetter beds. Preliminary test results shows that the outer
jacket helium circulation reduces 25 % of cooling timecomparing to the natural cooling by helium filled outer
jacket. Various combination of active cooling method isunder experiments and its results will be presented.

3. Verification of SDS Process
3.1 Cooling rate enhancement test
Metal hydride bed is one of the key components in
ITER SDS and fusion fuel cycle. As mentioned in section
2.3 (experimental plan for the control logic development),
SDS beds are heated for desorption of deuterium andtritium gases. At the same time, SDS beds must be readyfor the absorption of DT gases recycled from the fuelcycle. In addition to some flocks of beds have to be readyfor the emergency uptake of DT gases. For the better
performance of absorption and emergency uptake of DT

Fig. 9 Schematic He-3 recovery system in ITER SDS.

60

S.-H. Yun et al., The R&D Status of ITER SDS

Fig. 10 Schematics of active cooling for hydride bed.

Fig. 11 Installation of improved design 1:1 mock-up bed with active cooling system.

61

S.-H. Yun et al., The R&D Status of ITER SDS

installation drawing, an overall flow diagram, and one of
the typical operation schemes for absorption/desorption
performance test, respectively.

3.2 1:1 scaled DU hydride bed system

SDS beds and major components including DU
system has to be experimented to verify the current SDS
technology. Fig. 12(a), 12(b), 12(c) show a schematic 3D(a) Schematic 3D installation drawing(b) Overall flow diagram for the SDS verification(c) Typical operation for absorption/desorption test
Fig. 12 DU hydride system (1:1 scaled system installation).

62

S.-H. Yun et al., The R&D Status of ITER SDSThe SDS verification test apparatus (see Fig. 12(b))
has various kinds of SDS operation processes that shouldbe verified for the technical completion of SDS. Moreover,through this verification test a clear distinction has to beshown the operability of this facility in two safety aspects:
one for tritium and the other for DU.
Fig. 12(c) is one typical operation mode in this test
apparatus. The other operation modes are to be tested asfollows:
DU bed experiments with;. Fuel receiving rate in absorption phase. Fuel delivery rate in desorption phase (one bed)
. Fuel delivery rate with multiple beds
. Tritium accountancy using in-bed calorimetry. Bed cooling with in-bed calorimetry circulation. Bed cooling with outer jacket circulation. Bed cooling with external cooling & circulation. Bed cooling with overall combination. He-3 recovery and collection. Transfer pump performance and combinations. PVT-c measuring and mass flow meter test
. Other instruments test.

[6]
[7]
[8]

4. Conclusion

The ITER SDS is on-going at pre-conceptual R&D
phase. The SDS R&D is currently focused on DU getterbed system which requires careful approach to master andverify every component’s performance and application
skills. Basic approaches and 1:1 scaled process verification
apparatus in SDS R&D are to be prepared on this basis.
The current status and activities, especially setup ofhandling procedure, test devices and facility design basedon DU getter material was described in this report. Every
R&Ds will be proceeding for the CD completion of the
ITER SDS next year.

Acknowledgment

This research was supported by National R&D
Program through the National Research Foundation of
Korea (NRF) funded by the Ministry of Education, Scienceand Technology & Ministry of Knowledge Economy(2012-0000293).

5. References
[1]
[2]
[3]
[4]
[5]

H. Chung et al., Fusion Engineering and Design 84, 599(2009).
S. Cho et al., IEEE Transactions on Plasma Science 38, 3,
425 (2010).
S. Cho et al., Fusion Science and Technology 60, 1077
(2011).
M. Glugla et al., presented in Tritium 2010 Conference,
Nara, Japan.
H. Chung et al., Fusion Engineering and Design,63

doi:10.1016/j.fusengdes.2011.12.002 (2012).
D. Chung et al., KNS spring conference, Jeju (2012).
S.-H. Yun et al., Fusion Engineering and Design 86, 2282
(2011).
S.-H. Yun et al., Fusion Science and Technology 60, 373(2011).

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Trapping of Tritium by Stainless Steel Exposed to Plasmas inExperimental Campaigns of LHD
M. Matsuyama1), S. Abe1), K. Nishimura2), Y. Ono3), Y. Oya4), K. Okuno4), T. Hino5) and A. Sagara2)* masao@ctg.u-toyama.ac.jp
1)

Hydrogen Isotope Research Center, Univ. of Toyama, Gofuku 3190, Toyama 930-8555, Japan
2)

3)

National Institute for Fusion Science, 322-6 Oroshi-cho, Toki 509-5292, Japan

Center for Research and Development in Natural Sciences, Univ. of Toyama, Gofuku 3190, Toyama 930-8555, Japan
4)

5)

Radiochemistry Research Laboratory, Shizuoka Univ., Shizuoka 422-8529, Japan
Laboratory of Plasma Physics and Engineering, Hokkaido Univ., Sapporo 060-8628, Japan(Received: 7 May 2012 / Accepted: 5 October 2012)

Retention behavior of tritium by stainless steel exposed to plasmas in the Large Helical Device (LHD) has beenstudied by E-ray-induced X-ray spectrometry (BIXS) and an imaging plate (IP) technique. Plasma-exposed stainlesssteel samples were prepared by exposing the small plates at the location of 6.5L in LHD. X-ray photoelectron
spectroscopy (XPS) was also applied to analyze deposition layers on the surface of a sample. XPS analyses showedthat the deposition layers contain B, C, O and Ti as well as constituent elements of stainless steel. Chemical form of
metallic elements was mainly oxides, but a part was metallic state. The plasma-exposed sample was exposed totritium gas along with a bare stainless steel plate for comparison. The BIXS and IP measurements showed that thetritium retention largely increased by plasma exposure and it strongly depended on sample temperatures during
vacuum heating and tritium exposure. In addition, it was seen from the IP images that non-uniform tritium
distribution on the surface was formed even tritium exposure at room temperature.
Keywords: tritium retention, plasma exposure, BIXS, imaging plate, plasma-facing materials
samples were fixed at four different locations, which are
1.5U, 5.5U, 6.5L and 9.5L in LHD [2], and those samples
were exposed to plasmas for one cycle in the 13th
experimental campaign. After plasma exposure, they
were applied to tritium exposure at 623 K for 4 hours
after heating in vacuum at 673K.
It was seen that all of the stainless steel samples werecovered with deposition layers, although their thickness
was different from each other. Tritium retention in eachsample was strongly dependent on the sample location [1].
In addition, significant non-uniform tritium distribution
appeared in the plasma-exposed surface of the 6.5L
sample, although other samples showed almost uniform
distribution. Therefore, the present study was focused on
the trapping behavior of tritium in the 6.5L sample.

1. Introduction

Reduction of tritium retention in the plasma-facing
materials (PFMs) and the structural materials of fusion
experimental device such as ITER is one of great
important issues from viewpoints of the safety andeconomy of tritium as well as the fuel particle balance inthe reactor core, since a huge amount of tritium is used in
the device. These materials are exposed to various energy
particles, and surfaces of them are modified. Deposition
layers are formed on the surfaces as a result of plasma
exposure. They will affect tritium retention. The tritium
retention strongly depends on surface states such as
chemical state and composition of constituent elements inthe deposition layers. Namely, it is important tounderstand the correlation between tritium retention andsurface states of deposition layers.
We have examined so far the trapping and release
behavior of tritium in the PFM’s of the Large Helical
Device (LHD) as an example of the effects of exposure toplasmas for tritium retention [1]. Stainless steel plate isused as a protection plate in LHD. Small stainless steel

2. Experimental

The samples used in this study are stainless steel type
316 (SS316), and size of the sample plates was 10×10mm and 0.5 mm in thickness. These sample plates were
finally electropolished, and they were fixed at four64©2013 by The Japan Society of Plasma
Science and Nuclear Fusion ResearchM. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD

0.10

A

0.08

Sample: 6.5L(13 cycle)
Evacuation temp. : 296 K
Exposure temp. : 296 K
Exposure Time : 4 hrs
Tritium gas : 2.66 kPa

0.06
0.04
0.02
0.00

0

1.0

2

4

6

Counting rate, N / counts min-1 ('E)-1

B

0.8

8

10

12

0.6
0.4

16

18

0.4

Ar

Ti

0.3

0.2

0.1

0.0

Cr
0

2

4

Fe

6

8

Energy of X-rays, E / keV

10

12

Sample: 6.5L(13 cycle)
Evacuation temp. : 673 K
Exposure temp.: 296 K
Exposure Time : 4 hrs
Tritium gas : 2.66 kPa

0.2
0.0

14

Energy of X-rays, E / keV

0

10

2

4

6

C

8

8

10

6
4

14

Ar

1.0

16

Cr
Mo

0.6

Ti

0.4

Cr

0.2
0.0

18

Fe

0.8

0

2

4

6

Fe Ni
8

Energy of X-rays, E / keV

10

12

Sample: 6.5L(13 cycle)
Evacuation temp. : 673 K
Exposure temp.: 623 K
Exposure Time : 4 hr
Tritium gas : 2.66 kPa

2
0

12

Energy of X-rays, E / keVCounting rate, N / counts min-1 ('E)-1

Counting rate, N / counts min-1 ('E)-1

Counting rate, N / counts min-1 ('E)-1

Counting rate, N / counts min-1 ('E)-1

different walls in LHD before beginning of the 13th
plasma campaign, which was conducted for 13 weeks
from October 2009. After plasma exposure, only 6.5L
sample was provided for surface analyses and tritium
exposure experiments in this study.
As a first step, the plasma-exposed surface wasobserved by a digital microscope. This sample was also
used to analyze chemical states of surface elements by
X-ray photoelectron spectroscopy (XPS). The XPS
analysis was performed on a Thermo Fisher Scientific
Model ESCLAB250Xi XPS system using Al(KD) X-rays
with an energy of 1486.6 eV.
Tritium exposure tests were conducted by thefollowing procedures. The 6.5L sample and a bare SS316
plate with the same size as a reference sample were set in
the tritium exposure device, and it was evacuated below
2x10-6 Pa at room temperature. The latter sample wasused for comparison. In order to examine the effects of
pre-heating in vacuum of the samples, at first, they were
exposed to tritium gas at room temperature. After
E-ray-induced X-ray spectrometry (BIXS) and imaging
plate (IP) measurements of them, as the second
examination, those samples were again set in the tritium
exposure device and then it was evacuated below 6x10-6
Pa at 673 K. After this evacuation, the samples were
exposed to tritium gas at room temperature for 4 hours.
Subsequently, as the third examination, the samples wererepeatedly set in the tritium exposure device, and thosewere heated at 673 K in vacuum and exposed to tritium
gas at 623 K for 4 hours. Total pressure and tritium
concentration of tritium gas used in this study were 2.66
kPa and 4.8%, respectively.
After each procedure of the tritium exposure, toestimate the amount and distribution of tritium retained insurface layers, both of 6.5L and bare SS316 sample were
contacted with an IP along with a set of standard tritium
sources, which were made of a polymer material labeledwith tritium. In this procedure, a thin poy-phenylenesurfide film (PPS: 1.2 Pm) was inserted between thesamples and IP to avoid contamination of the IP surface.
After exposure to IP for a given time, two-dimensional
mapping of E-ray dose was obtained by measuring theintensity of photo-stimulated luminescence with theimaging analyzer (Fujifilm FLA-7000). The spatial
resolution was 25×25 Pm. Some examples of tritium
measurements by a similar IP technique have been
recently reported by Hatano [3-4]. The BIXS was also
employed to evaluate the tritium retention in surface
layers and the tritium distribution in depth of a sample.
The details of BIXS are described elsewhere [5].

0

2

4

6

8

10

12

14

Energy of X-rays, E / keV

16

18

Fig. 1 X-ray spectra observed by BIXS for 6.5L sample.
Each of upper, middle and bottom figures was observed after
vacuum heating and tritium exposure under the following
conditions. (A) ; heated at 296 K and exposed at 296 K, (B);heated at 673 K and exposed at 296 K, and (C); heated at 673
K and exposed at 623 K. The dashed line in the figure (A)shows the background level of the present X-ray detector.

3. Results and discussion

3.1 Tritium measurements by BIXS
Figure 1 shows the X-ray spectra observed by BIXS
measurements. All measurements by BIXS were carriedout under the argon atmosphere. Each X-ray spectrum
consists of characteristic X-ray peaks and a bremsstrahlung
X-ray peak. Number of characteristic X-ray peak observed
was dependent on conditions of the vacuum heating andexposure temperature. In the case that both temperatures

65M. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD

2

4

6

Length, L / mm

8

10

0.002
0.000

0

2

4

6

Length, L / mm

8

0.04
0.02
0.00

10

0

2

4

6

Length, L / mm

8

0.06
㻌

㻌

0.06

0.04
0.02
0.00

10

Sample: 6.5L(#12)(13 cycle)
Vac.: 673 K
Expos.: RT
Time: 4 hr

0.08

0

2

4

6

Length, L / mm

810Fig. 3 IP images and line analyses for bare SS316 and 6.5L
samples. Vacuum heating was 673 K and exposure
temperature was 296 K.

Fig. 2 IP images and line analyses for bare SS316 and 6.5L
samples. Both temperatures in vacuum heating and tritium
exposure were 296 K.

(bare SS316)

during vacuum heating and exposure processes were room
temperature as shown in Fig. 1 (A), only a single peak ofAr(KD) X-rays was observed. On the other hand, pluralcharacteristic X-ray peaks appeared as shown in Figs. 1 (B)and (C), when vacuum heating and/or exposure
temperature were raised. Those X-ray peaks were assigned
to Ti, Cr, Fe, Ni and Mo from energies of X-rays as shownin the inset of Figs. 1 (B) and (C). It is considered thatappearance of a rather large titanium peak in Fig. 1 (B) is
due to the use of a titanium sublimation pump as one of
vacuum pumps in LHD. It should be noted that intensity of
the Ti(KD) peak in Fig. 1 (B) is 5 times larger than that ofFe(KD) peak. Other characteristic peaks are due toconstituent elements deposited on the surface by sputteringof stainless steel walls with plasmas in LHD.
It is known that intensity of an Ar(KD) peakcorresponds to the amount of tritium retained in surface
layers of a sample because of a short escape depth of
E-rays [6]. Intensity of an Ar(KD) peaks in Figs. 1 (A), (B)and (C) was 0.13, 3.5 and 39 counts/min, respectively. The
amount of tritium retained in surface layers largely
increased with increasing temperatures of vacuum heating
and tritium exposure: namely, tritium activity increased 27
times by elevating the temperature of vacuum heating, andit increased about 300 times by elevating both temperaturesof vacuum heating and tritium exposure. The former
increase may be due to removal of impurities adsorbed on
the surface and diffusion of metallic species from bulk tothe surface since the tritium exposure temperature was thesame. On the other hand, it is expected that the latter
increase was caused by acceleration of tritium diffusion
into the deposited layers since the exposure temperature
increased.
In addition, the shape of the bremsstrahlung X-ray peakin Fig. 1 (C) was different from that in Figs. 1 (A) and (B).
Namely, the shape in Fig. 1 (C) was convex above 3.5 keV.

㻌

1.0

Intensity of PSL, I / hr-1

Plasma-exposed SS316 (6.5L)

0.8
0.6
0.4
0.2
0.0

㻌

1.0
Sample: SS316
Evacuation temp.: 673 K
Exposure temp. : 623 K
Exposure Time: 4 hr
Tritium gas: 2.66 kPa

0

2

4

6

Length, L / mm

810Sample: 6.5L(13 cycle)
Evacuation temp.: 673 K
Exposure temp. : 623 K
Exposure Time: 4 hr
Tritium gas: 2.66 kPa

0.8
0.6

㻌

0

0.004

0.08

Intensity of PSL, I / hr-1

0.002

0.006

㻌

0.10
Sample: SS316
Vac.: 673 K
Expos.: RT
Time: 4 hr

㻌

0.004

0.008

Plasma-exposed SS316 (6.5L)

㻌

0.10
Sample: 6.5L(#12)
13 cycle
Vac.: RT
Expos.: RT
Time: 4 hr

㻌

㻌

0.006

Intensity of PSL, I / hr-1

0.008

0.000

㻌

0.010
Sample: SS316
Vac.: RT
Expos.: RT
Time: 4 hr(bare SS316)

Intensity of PSL, I / hr

㻌

0.010

Intensity of PSL, I / hr-1

Plasma-exposed SS316 (6.5L)

Intensity of PSL, I / hr-1

(bare SS316)

0.4
0.2
0.0

0

2

4

6

Length, L / mm

810Fig. 4 IP images and line analyses for bare SS316 and 6.5L
samples. Vacuum heating was 673 K and exposure
temperature was 623 K.

This indicates the diffusion of tritium into the bulk as wellas adsorption on the surface. It can be seen, therefore, thatthe amount of tritium retained in surface layers is strongly
affected by both of vacuum heating and exposure
conditions even though the surface is covered with thedeposited layers.
3.2 Measurements of surface tritium distribution by IP
techniqueSurface tritium distribution was measured by an IP
technique. To determine tritium activity in surface layers,
seven standard polymer plates labeled with a given
concentration of tritium were used. Main deposition of thesurface of plasma-exposed sample was carbon describedlater and its thickness was thinner than about 1Pm.
Therefore, escape depth of E-rays in the sample wasassumed to be almost same as the standard samples. Figure
2 shows the observed IP image (upper) and the result of66

Intensity, I / arb. units

M. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD

21 nm
Fe(2p)

O(1s)Intensity, I / arb. units

1000

Fe(2p)C(1s)
Cr(2p)Mo(3p)

800

600

400

Fe(3p)

200

Binding energy, EB / eV0Fig. 5 An example of XPS spectra observed by wide scan.

[B]

[A]

center A

O(1s)
Cr(2p)

10 nm

C(1s)
Mo(3d)
B(1s)w/o etching

red
21 nm

B

10 nm

w/o etching

1000

800

600

400

200

Binding energy, EB / eV0Fig. 7 Changes in XPS spectra for two spots with depth.

This is clearly shown by the result of line analysis. Tritium
activity in the green region was 4.06 GBq/m2, which isabout twice of the average concentration. These indicatethat high tritium retention appears even though exposure
temperature is low, when the sample was pre-heated in
vacuum at high temperature. It is considered, therefore,that the large increase in tritium retention is due to thechange in adsorption properties of deposition layers byvacuum heating as mentioned above.
Further examination for the vacuum heating effects was
carried out: namely, both samples were simultaneously
exposed to tritium gas at 623 K after vacuum heating at
673 K. The results are shown in Fig. 4. Tritium retention insurface layers of the 6.5L sample was estimated to be 56.0
GBq/m2 as an average. This activity was 437 times larger
than that observed in Fig. 2, and it was close to the result
observed by BIXS. On the other hand, tritium retention ofbare SS316 was 3.23 GBq/m2. Tritium retention of the
6.5L sample was about 17 times larger than that of bare
SS316 sample. Furthermore, remarkably non-uniform
tritium distribution appeared clearly in the left side of 6.5L
sample and the tritium activity in the yellow region was ashigh as119 GBq/m2.

Fig. 6 Photograph of the sample surface and the analyzed
spots by XPS.

line analysis (lower) at the red broken line in the image,where both the vacuum heating and exposure temperatures
were room temperature. In comparison, similar
measurements for a bare stainless steel that is not exposed
to any plasma are also shown. Image of tritium distribution
measured by IP technique is represented by a color
measure. Red color shows the highest activity, while blue
color does the lowest activity. Namely, tritium activity
decreases relatively with the following color change: red,
yellow, green and blue. Surface tritium activity of the 6.5L
sample was estimated to be 128 MBq/m2, while that ofbare SS316 was 16 MBq/m2. Difference in tritium activity
was as large as 8 times. It is clear that such a large
difference in tritium retention is due to the plasma
exposure.
To examine the effects of vacuum heating, exposure totritium gas was carried out at room temperature after
vacuum heating at 673 K. After this, surface tritium
distribution was similarly measured by IP technique. The
results are shown in Fig. 3. Tritium retention in surface
layers of the 6.5L sample was estimated to be 1.97 GBq/m2
as an average, while that of bare SS316 sample was 0.16
GBq/m2. The difference in tritium retention of both
samples became larger by increasing in temperature of
vacuum heating. In addition to this, high retention area
appears in the left side of the 6.5L sample as green color.

3.3 XPS analyses of the plasma-exposed surface
In order to examine causes of non-uniformity in thetritium distribution on the deposition layers of the 6.5L
sample, the plasma-exposed surface and its depth profile
was analyzed by XPS without any heating treatment.
Figure 5 shows an example of XPS spectra observed by67

M. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD

Fe(2p)

Fe(2p3/2)

A

Cr(2p)

Cr(2p3/2)

Cr(2p1/2)

Fe(2p1/2)A

Ti(2p)21 nm

A

Ti(2p3/2)

21 nm

w/o etching

B
21 nm

10 nm

w/o etching

B
21 nm

Intensity, I / arb. units

Intensity, I / arb. units

10 nm

㻌

Intensity, I / arb. units

21 nm

10 nm

10 nm
w/o etching

590

585

580

575

570

Binding energy, EB / eV

Fig. 8 Changes in Fe(2p) peak withdepth.

Fig. 9 Changes in Cr(2p) peak withdepth.
(-CH2-)
Carbide A

21 nm

475

Intensity, I / arb. units

10 nm

w/o etching

B
21 nm

470

465

460

455

450

Binding energy, EB / eV

Fig. 10 Changes in Ti(2p) peak withdepth.
(OH-)(O2-)

O(1s)

21 nm

Intnsity, I / arb. unitsB

w/o etching

Binding energy, E / eVC(1s)w/o etching

10 nm

w/o etching

740 735 730 725 720 715 710 705 700

10 nm

A

21 nm
10 nm
w/o etching

B

21 nm

10 nm

10 nm
w/o etching

w/o etching

295

290

285

545

280

Binding energy, EB / eV
Fig. 11 Changes in C(1s) peak withdepth.

540

535

530

Binding energy, EB / eV

525

Fig. 12 Changes in O(1s) peak with depth.

two analyzed spots of the deposition layers were measured
with etching by argon. The results are shown in Fig. 7.
Etching rate by argon ions was assumed to be 1.0 nm/s,
which was obtained from the sputtering of Ta2O5 by argon.
Difference in peak intensity was observed between two
spots: namely, intensity of a carbon peak in the spot [B]was less than that in the spot [A], while that of oxygen was
contrarily larger. It should be noted that intensities ofmetallic Fe, Cr and Mo peaks in spot [B] is larger thanthose in spot [A], although oxygen peak in spot [B] is more
intense than that in spot [A]. Such a remarkable difference
in both measurement spots was observed for the
as-received surface only, and it became small with depth.

wide scan from 0 to 1200 eV. The analyzed spot wasalmost center of the sample surface. As clearly seen fromthe spectrum, major elements on the surface are carbon andoxygen. In addition to these elements, small amount ofelements such as iron, chromium and molybdenum also
deposited on the surface. However, titanium was scarcely
deposited in surface layers.
Depth profiles of two spots on the sample surface, [A]and [B], were measured as an example. Location of two
spots is shown in Fig. 6. The former spot was the center ofthe surface, while the latter one is the region of left side,which corresponds to the region in which much tritium wasretained as shown in Fig. 4. Changes in XPS spectra for68

M. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD

that in point [A] as is shown in Figs. 3 and 4. It isconsidered, therefore, that the amount of carbon on thesurface of deposition layers is one of key factors for tritium
retention.
Figure 12 shows the changes in O(1s) peak with depth.
At least, the peak of O(1s) consists of two kinds of oxygen
species: O2- (530.1 eV) and OH- (531.5 eV) [11].
Assignment of these peaks coincides with that of oxide
peaks of iron and chromium observed in Figs. 8 and 9.
There are not large differences in a chemical state ofoxygen for both analyzed spots.

Figs. 8 to 12 show the spectra observed by narrow scanof each element for two spots of [A] and [B]. The observed
spectra for iron are illustrated in Fig. 8. It is seen frombinding energy that chemical forms of iron consist of
metallic iron (707.1 eV (2p3/2) and 720.1 eV (2p1/2)) andthe mixed oxides of FeO, Fe2O3 and FeO(OH) (710.1 eV(2p3/2) and 723.2 eV(2p1/2)) [7]. Peak of oxide assigned to
the latter species shifted to the lower binding energy sidewith depth. This indicates that decrease in the amounts of
FeO(OH) and Fe2O3 is larger than that of FeO with depth.
Great difference in two analyzed spots is the fraction of
metallic iron as mentioned above. Metallic iron, which is
known as one of catalysts available for the adsorption ofhydrogen, will contribute to the adsorption of tritium [8].
In addition, existence of iron carbide was also suggested
from a C(1s) spectrum as described later.
The observed spectra for chromium are illustrated inFig. 9. It was seen that chemical forms of chromium aremetallic chromium (574.1 eV (2p3/2) and 583.6 eV (2p1/2))
and its oxides such as Cr2O3 and CrO(OH) (576.6 eV(2p3/2) and 586.2 eV (2p1/2)), although intensities of thesepeaks are fairly weaker than that of iron peaks. Fraction of
metallic chromium increased with the increase in depth.
This is a similar tendency to the changes in a peak intensityof iron.
The observed spectra for titanium are illustrated in Fig.
10. Only a weak peak appeared in the spectra. It was seenthat a chemical form of titanium is TiO2 (458.5eV (2p3/2)),
but metallic titanium (453.8 eV (2p3/2) can not be foundwithin a depth of 21 nm. However, an intense X-ray peakof titanium appeared as is shown in Fig. 1 (B), where thesample was heated at 673 K in vacuum and exposed totritium gas at 296 K. It is considered that titanium atoms
diffuse from bulk to the surface by vacuum heating, andtritium concentration near titanium atoms becomes high byannealing at high temperature in vacuum. Further
investigations by XPS are required about the diffusion oftitanium after vacuum heating at 673 K.
As is shown in Fig. 5, most of the surface was covered
with carbon and oxygen atoms. Figure 11 shows thechanges in C(1s) peak with depth. Two peaks apparently
appeared in each spot. The peak of 283.2 eV was assigned
to iron carbide [9], which may not contribute to the tritium
adsorption. The higher energy peak (284.9 eV for the spot[A], 284.5 eV for the spot [B]) consists of plural peaks, andthey were assigned to chemical forms such as graphite/aromatics and (-CH2-) [10]. There was a large difference inthe as-received surface: namely, carbon content in the spot
[A] is appreciably larger than that in the spot [B], andcarbon species in the spot [B] consist of chemical forms ofgraphite, (-CH2-) and carbide, although the content andchemical form of carbon in bulk is similar to each other.
Tritium retention in point [B] was several times larger than

4. Summary

Tritium retention behavior on the surface of a stainlesssteel plate, which was previously exposed to plasmas at thelocation of 6.5L in the Large Helical Device (LHD) during
one cycle, has been studied. To clarify the effects of plasma
exposure, the plasma-exposed sample along with a baresample were exposed to tritium gas under the given
temperature conditions. Three different temperature
conditions of the vacuum heating and tritium exposure
were applied to these samples.
It was seen from the BIXS and IP measurements thatthe amount of tritium retained in the plasma-exposed
surface was larger than that in the bare surface irrespectiveof temperature conditions and it remarkably increased inincreasing temperatures of vacuum heating and tritium
exposure. Large heating effects for tritium retention
appeared in the plasma-exposed sample than the baresample of stainless steel. Furthermore, significantly intense
X-ray peak of titanium was observed by vacuum heating at
673 K, although it did not appear at 296 K.
Non-uniformity of tritium distribution on theplasma-exposed surface was confirmed by the images of IP
irrespective of temperatures of the vacuum heating andtritium exposure. XPS analyses of the as-received surfaceof the 6.5L sample indicated the existence of metallic
chromium and iron, although the surface was exposed toambient atmosphere after exposure to plasmas. It isconsidered, therefore, metallic species in the depositedlayers are responsible for the non-uniformity of tritium
distribution. This will be clarified by further XPS analysesof the sample processed at high temperatures.

Acknowledgement

This work was supported by the Collaboration
Research Program of the National Institute for Fusion
Science (NIFS10KUHR001).

References[1] M. Matsuyama, et al., Plasma and Fusion Research 7,2405091-1 (2012).

69

M. Matsuyama et al., Trapping of Tritium by Stainless Steel Exposed to Plasmas in Experimental Campaigns of LHD[2] T. Hino, Y. Nobuta et al., Fusion Eng. Design 82,
1621 (2007).
[3] Y. Hatano, et al., J. Nucl. Mater. 386-388, 569 (2009).
[4] Y. Hatano, et al., Fusion Sci. Technol. 60, 982 (2011).
[5] M. Matsuyama, et al., J. Nucl. Mater. 290-293,
437(2001).
[6] M. Matsuyama, et al., Fusion Sci. Technol. 41, 505(2002).
[7] C.R. Brundle, et al., Surface Sci. 68, 459 (1977).
[8] G.C. Bond, “Catalysis by Metals“, Academic Press,
London, (1962), p.65.
[9] P.H-Fierro, et al., J. Vac. Sci. Technol. A11,354(1993).
[10] E. Desimoni, et al., Surface and Interface Analysis
15, 627 (1990).
[11] T. Hanawa, et al., Materials Trans. 43, 3088 (2002).

70

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Simulating Tritium Retention in Tungsten with a Multiple Trap Modelin the TMAP Codea

Brad J. Merrill, Masashi Shimada, Paul W. Humrickhouse

Fusion Safety Program, Idaho National Laboratory, Idaho Falls, ID, USA 83402(Received: 10 May 2012 / Accepted: 1 September 2012)

Accurately predicting the quantity of tritium retained in plasma facing components is a key safety issue for licensing
future fusion power reactors. Retention of tritium in the lattice damage caused when high energy neutrons collide withatoms in the structural material of the reactor’s plasma facing components (PFCs) is an area of ongoing experimental
research at the Idaho National Laboratory (INL) under the US/Japan TITAN collaboration. Recent experiments with the
Tritium Plasma Experiment (TPE), located in the INL’s Safety and Tritium Applied Research (STAR) facility,
demonstrate that this damage can only be simulated by computer codes like the Tritium Migration Analysis Program(TMAP) if one assumes that the lattice damage produced by these neutrons results in multiple types of hydrogen traps(energy wells) within the material, each possessing a different trap energy and density. Previous attempts to simulate thequantity of deuterium released from neutron irradiated TPE tungsten targets indicated that at least six different traps are
required by TMAP to model this release. In this paper we describe a recent extension of the TMAP trap site model toinclude as many traps as required by the user to simulate retention of tritium in neutron damaged tungsten. This model has
been applied to data obtained for tungsten irradiated to a damage level of 0.025 dpa in the High Flux Isotope Reactor(HFIR) at the Oak Ridge National Laboratory (ORNL) after exposure to a plasma in TPE.
Keywords: fusion, safety, tritium, retention, trapping models, TMAP.

1. Introduction

In this paper we describe recent extensions of the
TMAP trap site model to treat as many traps as requiredby the code to simulate the retention of tritium in neutron
damaged tungsten material. In the sections that follow,
we describe the equations for these new trap site models,the experiments conducted at the INL to develop the dataneeded to benchmark the implementation of thesemodels, and the application of these models to deuterium
retained in neutron irradiated tungsten after exposure to a
TPE plasma.

Accurately estimating the quantity of tritium retainedin plasma facing components (PFCs) is a key safety
requirement for licensing future fusion power reactors,
as well as for the International Thermonuclear
Experimental Reactor (ITER) [1]. These components
not only see the very harsh environment produced by aplasma, such as surface erosion, severe surface heat flux,
and high rates of plasma ionic specie implantation, but
the 14 MeV neutrons generated by the fusion processcreate significant damage to the PFC material lattice byatomic displacements and elemental changes throughtransmutations [2]. This damage acts as ideal traps fortritium retention as implanted tritium diffuses towards
the surfaces of the PFC, where it is either released backinto the vacuum vessel, or into the PFC’s coolant.

2. Trapping Models in TMAP
The TMAP calculates the time-dependent response
any number of gaseous species in a system of solid
structures or walls, and related gas filled enclosures by
including:

Tritium retention in PFC materials is an area ofexperimental research at the Idaho National Laboratory(INL) under the US/Japan TITAN collaboration. Recent
experiments with the Tritium Plasma Experiment (TPE),
located in the Safety and Tritium Applied Research(STAR) facility at the INL, demonstrate that this damage
can only be simulated by codes like the Tritium
Migration Analysis Program (TMAP) [4] if multiple trapsites are used to simulate the damage generated byneutrons [3]. The TMAP code has been released as
various versions, the most recent of which has thecapability of modeling at most three user specified trapsite types. However, even with this number of sites,
TMAP was unable to adequately simulate deuterium
release from tungsten targets exposed to TPE plasmas
that had first been irradiated to a damage level of 0.025
dpa in the High Flux Isotope Reactor (HFIR) at the Oak
Ridge National Laboratory (ORNL).

•

Implantation of species in structures,•

Movement of species through structuresurfaces, governed by molecular dissociation
and atomic recombination, or by a solution law
such as Sieverts’ or Henry’s Laws•

Movement in the structure as treated by Fick’s
law of bulk diffusion with the possibility of
specie trapping in material defects,•

Thermal response of structures to applied
heating or boundary temperatures,•

Chemical reactions within the enclosures, and•User specified
enclosuresconvective

flow

betweenBecause the equations used to simulate these processes
are highly nonlinear, a Newton-Krylov type solver is_______________________________________________________________________________
aBrad.Merrill@inl.gov

71©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

B.J. Merrill et al., Simulating Tritium Retention in Tungsten with a Multiple Trap Model in the TMAP Codeused to update these equations at each point in time
during a calculation.

where μ is the lognormal mean and ı is the lognormal
standard deviation. Given this distribution, the total
concentration of trap sites (cto) at any location in the bulk
material equals:

The basic equation for movement of these species inthe bulk material is the one-dimensional Fick’s Second
Law of diffusion written as follows:nwC iw § wC m ·¸  Sm  ¦ t¨D
wx © wx ¹
i 1 wtwC m
wt

x

D is specie diffusivity (m2-s-1)x

Sm is a specie source term (m-3-s-1)

x

Cti is the trapped specie concentration in the“ith” trap site (m-3)

xn is the number of traps allowed by TMAP

C t ( x, t )wCt
wtĮt is the trapping rate coefficient (s-1)

x

ft is the probability of landing in a trap site (-)

xĮr is the release rate coefficient for exiting outof that trap site (s-1)

O

; ft

cto  Ct
§ E ·; D r Q o exp¨  t ¸N© kT ¹(3)

In order to allow the use of more than the three traps
presently allowed by TMAP7, it was a trivial matter ofmerely changing the TMAP code to accept more thantraps, which was also done for the purpose of this study.
However, when viewing the Thermal Desorption System(TDS) data from neutron irradiated tungsten it became
apparent that perhaps a more accurate approach would
be to modify the TMAP code trapping model to accept
any distribution of trap site energies; in particular for thisstudy, a distribution based on previous work performedby the INL Fusion Safety Program (FSP) in developing
aerosol resuspension models, which is a lognormal orbimodal lognormal distribution [5].

­ >ln (E)  ȝ @ 2 ½
exp
®
¾
1/ 2
2ı 2
2ʌ ı E¯¿§ Įt C m³ ¨© Nf·c ot M(E)  Ct (E)  Į r E Ct (E) ¸ dE¹

(7)

3. Experiments Conducted at the INL
Under the US/Japan TITAN collaborative agreement,
tungsten samples that are 6 mm in diameter and 0.2 mm
thick were irradiated to three different damage levels(0.025, 0.3 and 2.4 dpa) in the HFIR reactor at ORNL ata temperature of 50ºC.
Following an appropriate cool down period, thesesamples were exposed to the plasma of the Tritium
Plasma Experiment (TPE) at INL’s STAR facility. The
TPE is a linear plasma device [6] capable of exposing
target material to low-energy (100 eV) deuterons at a
plasma ion flux of up to (5–7)×1021 mí2 sí1. For the

The density function for a lognormal trap distribution
in terms of trap energy is written as:

1f

While Eq. 7 could be used directly in Eq. 1, theTMAP code has always been meant to be a relatively
fast running computer code capable of performing quick
safety scoping and data reduction analysis tasks for the
INL FSP. To accomplish this for the present trapping
model, we again adopted the same approximation madefor our dust resuspension model. By noting that highenergy traps will be the last traps to release their trappedinventory, we assume that the traps of this distribution
will fill from high energy to low energy in that order,that is from i = N to 1 as illustrated in Fig. 1. The
advantage of this assumption is the TMAP is now only
tracking the total trapped inventory in this distribution.
Since the fraction of traps filled at any point in timeequals the trapped inventory divided by the trap density(e.g., F = Ct/cto), finding the trap site that is changing inconcentration for this approximation is a simple matterof interpolating with respect to the Cumulative
Distribution Function (CDF) for a lognormal
distribution. This trap is assigned the net transient
trapping and release rates for the entire distribution as its
concentration is changing in time.

and Ȝ is the jump distance or lattice constant (m), cto is
the trap site concentration, N is the bulk material number
density (atoms-m-3), Ȟo is the Debye frequency (s-1), Et is
the trap binding energy (eV) and k is Boltzmann’sconstant (eV/K).

M E:ȝ, V

(6)

§Į c
·
# ¦ ¨ t m c ot ǻM i  Cti  Į r i Ct i ¸¹
i 1© NThe definition of these terms is:
2

t

N

(2)

x

D

³ C ( x, t , E ) dE

Differentiating Eq. 6 with respect to time andsubstituting from Eqs. 2 and 3, the result for the rate ofchange of the total trapped concentration of a given
specie becomes:

where:

Dtf

f

The rate at which the mobile species are entering trap
defects in the bulk material is:t(5)

and the trapped concentration of a given specie Ct at thatlocation, and at a given point in time equals:

Cm is the mobile specie concentration (m-3)

wt

o
t

(1)

x

D t f t i Cm  D r C i³ c M E dE

fwhere:

wC ti

f

c ot

(4)



72

B.J. Merrill et al., Simulating Tritium Retention in Tungsten with a Multiple Trap Model in the TMAP Codeorder to verify TMAP’s trapping models against TPE
results, property data for tungsten diffusivity, solubility,
and surface recombination are required.

Fig. 1. Schematic of trap filling approximation for alognormal trap density distribution.
Fig. 2. TDS data of the deuterium retained by tungsten
irradiated in HFIR to a damage level of 0.025 dpa after
exposures to TPE plasma while maintaining the tungsten
temperature at 200ºC.

irradiated tungsten samples the plasma was maintained
until an ion fluence of (5–7)×1025 mí2 was achieved forsample temperatures of 100, 200 and 500ºC.
After the deuterium plasma exposure, the samples
were removed from TPE and were transferred to the TPE
Thermal Desorption System (TDS) in air, where thesamples were heated to produce a controlled linear
temperature change from 25 to 800ºC. The off-gas fromthe samples was passed through the Residual Gas
Analyzer (RGA) of the TDS that was specially calibrated
to measure the absolute concentration of off-gassing D2.
Only the peaks associated with HD (mass 3) and D2(mass 4) were considered in this analysis. Fig. 2
presents a TDS spectrum obtained for a 0.025 dpa
sample used in this study.
Several of the exposed samples were sent to the Ion
Beam Laboratory at University of Wisconsin, Madison
for Nuclear Reaction Analysis (NRA) [7] prior toundergoing assay in the STAR facility’s TDS. At thisfacility, a high energy 20 mA helium ion beam is
directed at the target to create a D(3He, p)4He nuclear
reaction with the deuterium implanted by a TPE plasma.
The high energy protons emanating from the sample areused to estimate the depth profile of the deuterium
retained in the sample. Given the relatively immobile
nature of deuterium in tungsten at room temperature, thisdepth profile can also be used to infer the hydrogen trapsite concentrations for those traps that saturate under
TPE irradiation conditions. Fig. 3 contains a sample of
NRA data as used in this study.

Fig. 3. Deuterium depth profiles for tungsten exposed toa TPE plasma at 200ºC as measured by NRA.
As was the case for Ref. [3], in this study we adopted
the hydrogen diffusivity (corrected for deuterium) andsolubility formulas proposed by Frauenfelder [8], where(D = 2.9×10-7 exp(-0.39 eV/kT) (m2/s) and Ks =1.79×1024 exp(-1.04 eV/kT) (D/m3-Pa1/2). In literature,surface recombination coefficient formulas can vary by
several orders of magnitude [9]. However at this time,
the only formula based on experimental data for tungsten
is that by Anderl of Kr = 3.2×10-15 exp(-1.16 eV/kT)

4. Application of Trapping Models to IrradiatedTungsten Deuterium Retention Data
A key component of any examination ofexperimental data is the validation of the physics models
being solved by the computer code being applied. In73

B.J. Merrill et al., Simulating Tritium Retention in Tungsten with a Multiple Trap Model in the TMAP Code(m4/s) was adopted for this study as a representativevalue for a plasma vacuum environment [10].

the mobile deuterium concentration in the tungsten
target, which is several orders in magnitude less than thetrapped concentration, is too low to fill these traps once
the TPE plasma has been terminated, the TPE plasma
ramp down had to be extended in time beyond what had
actually occurred. The result marked “no overfill” inFig. 4 illustrates the predicted TDS spectrum when theplasma was ramped down as reported. While additional
validation of this new TMAP capability is still needed toverify the accuracy of this result, it does suggest that thedeuterium being released in the TDS spectrum below
200ºC had nothing to do with trapping on dislocations
but is more likely associated with the deuterium that
resides in surface blisters known to develop in tungsten
at the implantation fluence experienced by this TPE
target (note discussion for Fig. 4 of Ref. [2]).

Physical data of trap density, D+ implantation depth,
plasma surface flux intensity and target temperatures arealso required. The NRA measured D/W atom fraction
presented in Fig. 3 was used as a starting value for the
total trap site density in neutron irradiated tungsten to adamage level of 0.025 dpa. An implantation depth of
3 nm, as estimated by Wampler [2], was used in ourstudy. In addition, in order to understand if these new
TMAP models accurately address the physics of plasma
driven retention in tungsten, the actual TPE plasma
operating history for the TDS data being examined must
be used. For the data presented in Fig. 2, the TPE
plasma flux (ramped up to 5×1021 D+/m2-s and held for7200 s) and target temperature (under plasma heating to
200ºC) histories for the TPE run that resulted in the TDS
data of Fig. 2 were follow in the TMAP model asreported by Shimada [3].
Our first goal is to verify the TMAP results obtainedby Shimada [3], for this same TDS case. A goodcomparison between TMAP7 and TDS data was
obtained by these authors by assuming six traps.
However, because that version of TMAP only allowed amaximum of three user specified trap sites, the results of
two separate TMAP simulations had to be manually
added to obtain the result presented. The trap energies
and total density for these traps were reported as 0.9, 1.1,
1.3, 1.5, 1.75 and 2.0 eV, and 1.3 at% D/W, respectively.
As noted by these authors, the trap density required tomatch the TDS was much higher than the NRA D/W
atom fraction measured for the 0.025 dpa sample that
underwent the NRA in Fig. 3. There are a number ofpossible explanations for this difference, including losses
of deuterium by diffusion and surface release between
the time of TPE irradiation and NRA examination,
uncertainty in the NRA results, or differences in TPE
operating history for the TDS sample of Fig. 2 and NRA
sample of Fig. 3, etc. However, because we do not havethe NRA depth profile for the sample used in the TDS
measurement of Fig. 3, it is equally plausible that thetrap density is 1.3 at% for the sample analyzed by the
TDS.

Fig. 4. Comparison of TMAP simulation with multiple
traps and the TDS spectrum for neutron irradiatedtungsten exposed to a TPE plasma at 200ºC.
Fig. 5 presents the results obtained for this same casewith the new TMAP lognormal trap model. An adequate
representation of the TPE TDS spectrum was obtainedwith a truncated (at 1.6 eV) bimodal lognormal
distribution, which is a distribution that is the sum of two
lognormal populations with proportions p and q, namely:

Fig. 4 contains the results of using six traps used inthis study to simulate the TDS of Fig. 3 using the newmulti-trap site TMAP model. Unlike the simulations ofRef. [3], where the decision to arbitrarily assign thetrapped deuterium concentration profile as the initialcondition for performing their TDS simulations, thisresult was obtained by simulating the implantation
process of the TPE plasma, with the goal of determining
if the adopted trap sites would actually saturate, asassumed by Ref. 3, at the irradiation temperature (200ºC)
experienced by the target. The six traps required tosimulate the TDS spectrum had energies of 0.9, 1.0, 1.1,
1.2, 1.4 and 1.6 eV, and a total trap density of 0.35 at%.

M E

pM1 E:ȝ1 ,V 1  qM 2 E:ȝ2 ,V 2

(8)

For the comparison in Fig. 5, p and q were both set to0.5, and the mean and standard deviations of thelognormal distribution ĳ1 were set at 0.95 eV and
0.18 eV and for ĳ2 set at 1.45 eV and 0.21 eV. The total
trap density for this case was 0.47 at%. As can be seen,
the simulation represents the data well. As was the casefor the six trap model, the lower energy traps could only
be filled by ramping TPE down more slowly that what
actually occurred.

One aspect of this simulation that was not true to
TPE plasma operation details deals with the filling of thelower energy trap sites. TMAP predicts that at 200ºC the
0.9 eV traps will not saturate. Since TMAP predicts that
74

B.J. Merrill et al., Simulating Tritium Retention in Tungsten with a Multiple Trap Model in the TMAP Coderesult, the results appear to suggest that the deuterium
being released in the TDS spectrum below a temperatureof 200ºC had nothing to do with trapping on dislocations
but is more likely associated with the deuterium that
resides in surface blisters known to develop in tungsten
at implantation fluence experienced by this TPE target.
Future work in this area with a broader set ofexperimental data must be undertaken to confirm thisconclusion.

Acknowledgments
This work was prepared under the auspices of the US
Department of Energy, Office of Fusion Energy Science,
under the DOE Idaho Field Office contract number DEAC07-05ID14517.

References[1]

N. TaylorǡǤǡǤǡǤǡ ǤǦǡǤ ǡ S. Reyes, L. Rodriguez-Rodrigo, S. Rosanvallon,
L. Topilski, “Updated safety analysis of ITER,” Fusion
Engineering and Design, 86, 619-622 (2011).
[2] W. R. Wampler, R. P. Doerner, “The influence of
displacement damage on deuterium retention in tungsten
exposed to plasma,” Nuclear Fusion, 49, 11502 (2009).
[3] M. Shimada, G. Cao, Y. Hatano, T. Oda, Y. Oya, M. Hara,and P Calderoni, “The deuterium depth profile inneutron-irradiated tungsten exposed to plasma,” Physica
Scripta, T145, 014051 (2011).
[4] G. R. Longhurst, “TMAP7: Tritium Migration Analysis
Program,” User Manual, Idaho National Laboratory,
INEEL/EXT-04–02352, Rev. 2 (2008).
[5] B. J. Merrill, P. W. Humrickhouse, J. P. Sharpe, ”Anaerosol resuspension model for MELCOR for fusion,”
Fusion Engineering and Design, 86, 2686–2689 (2011).
[6] M. Shimada, R. D. Kolasinski, J. P. Sharpe, and R. A.
Causey, “Tritium plasma experiment: Parameters andpotentials for fusion plasma-wall interaction studies,”
Review of Scientific Instruments, 82, 083503 (2011).
[7] M. Mayer, E. Gauthier, K. Sugiyama, U. von Toussaint,“Quantitative depth profiling of deuterium up to very large
depths,” Nuclear Instruments and Methods in Physics
Research Section B, 267, 506 (2009).
[8] R. Frauenfelder, “Solution and Diffusion of Hydrogen in
Tungsten,” J. Vac. Sci. Technol. 6 (3), 388 (1969).
[9] B. Lipshultz, et al., “An Assessment of the Current Data
Affecting Tritium Retention and its Use to Project Towards
T Retention in ITER,” Massachusetts Institute ofTechnology Report, PSFC/RR-10-4, April (2010).
[10] R. A. Anderl, D. F. Holland, G. R. Longhurst, R. J.
Pawelko, C. L. Trybus and C. H. Sellers, "Deuterium
Transport and Trapping in Polycrystalline Tungsten,"
Fusion Technology, 21, 745 (1992).

Fig. 5. Comparison of TMAP simulation with a bimodal
lognormal trap distribution and the TDS spectrum forneutron irradiated tungsten exposed to a TPE plasma at
200ºC.

5. Summary and Conclusions
Accurately estimating the quantity of tritium retainedin or permeating through PFCs is a key safety
requirement for licensing future fusion power reactors.
The magnitude of these processes is greatly affected bydamage created in the structural materials of these PFCs
by the 14 MeV neutrons generated by the fusion process.
This damage acts as ideal trap sites for tritium retention
within the PFCs.
In this paper, we have introduced two extensions to
the TMAP computer code trap site model. These
extensions allow the TMAP user to specify as many trapsites or a distribution of trap sites as needed to simulate
the material damage being modeled. These models were
benchmarked against data generated under the US/JapanTITAN collaborative research agreement in the INL’s
TPE device for retention of deuterium in tungsten that
had been irradiated to a damage level of 0.025 dpa in the
ORNL HFIR.
We found that the multi-trap extension to TMAP
could reproduce the off-gassing data from a tungsten
sample exposed to a TPE plasma by employing six traps
of energies of 0.9, 1.0, 1.1, 1.2, 1.4 and 1.6 eV, and a
total trap density of 0.35 at%. Alternatively, the data
could also be matched if a bimodal lognormal trap site
distribution is used that has proportions p and q equal
0.5, and mean and standard deviations of the lognormal
distribution ĳ1 of 0.95 eV and 0.18 eV and for ĳ2 of 1.45
eV and 0.21 eV, respectively. The total trap density forthis case was 0.47 at%.
While additional validation of these new TMAP
capability is still needed to verify the accuracy of this75

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Implantation Energy Dependence on Deuterium Retention Behaviors for the
Carbon Implanted Tungsten
Yasuhisa Oya1)*, Makoto Kobayashi1), Naoaki Yoshida2), Naoko Ashikawa3), Akio Sagara3),
Yuji Hatano4), and Kenji Okuno1)
1)

Radioscience Research Laboratory, Faculty of Science, Shizuoka University, 836, Ohya, Suruga-ku, Shizuoka 422-8529,
Japan.
2)

Institute for Applied Mechanics, Kyushu University, Fukuoka 816-8580, Japan.
3)4)National Institute for Fusion Science, Gifu 509-5292, Japan.

Hydrogen Isotope Research Center, University of Toyama, 3190, Gofuku, Toyama 930-8555, Japan.
(Received: 9 May 2012 / Accepted: 8 August 2012)

Effects of energetic carbon implantation on deuterium retention behavior in tungsten were studied. The
dislocation loops and vacancies were introduced in tungsten as irradiation damages by energetic carbon
implantation. The density of irradiation damages was almost saturated by C+ implantation below 1 dpa according to
TEM observations. The retentions of deuterium trapped by dislocation loops and vacancies were observed in TDS
measurements, which were increased as increasing the carbon implantation energy. The deuterium desorption at
higher temperature above 600 K was also found, corresponding to the desorption of deuterium retained in the
intrinsic cavity in bulk region. However, this trapping was refrained by energetic C+ implantation, concluding thatthe retained carbon would play a role as the diffusion barrier of deuterium, which would prevent deuterium
penetration through bulk.
Keywords: Tungsten, TDS, TEM, Carbon, Irradiation damage, Diffusion, Simultaneous implantation

introduction of irradiation damages was clearly observed
because 10 keV C+ implantation induced higher
displacement damage in tungsten than that by deuterium[6]. On the other hand, it is reported that the carbon
retained in tungsten plays a role of diffusion barrier forhydrogen isotopes toward the bulk, resulting in thedecrease of hydrogen isotope retention in tungsten [7].
Recent ion-driven permeation experiments under carbon deuterium simultaneous implantation indicate that
deuterium permeation is promoted by simultaneous
carbon implanted tungsten [8, 9]. These results indicate
the estimation of hydrogen isotope retention in tungsten
is complicated under fusion environment. Therefore,
deuterium retention in tungsten implanted with carbon
was investigated in this study to understand the effect ofcarbon implantation on deuterium retention behavior intungsten. For this objective, two types of implantation
were carried out. One was the simultaneous implantation(C+ - D2+ sim. imp.) where carbon and deuterium were
implanted into tungsten simultaneously. In thesimultaneous implantation circumstance, deuterium wasretained in tungsten with formations of irradiation defects
and W-C mixed layer. The other implantation was thesequential implantation (C+ - D2+ seq. imp.), meaning that

1. Introduction
Tungsten is a candidate for plasma facing materials(PFMs) in fusion devices due to its low hydrogensolubility, low sputtering yield and high melting point [1].
The usage of carbon fiber composite (CFC) is also
considered in the high heat load region of divertor infusion reactor, because of its high melting point [2].
These materials will be exposed to energetic particles
from plasma during reactor operation. It is well known
that carbon materials are easily eroded by energetichydrogen isotopes by chemical or physical sputtering
processes. Then, sputtered carbon will be implanted into
tungsten with hydrogen isotopes during plasma operation,
which will form the tungsten-carbon (W-C) mixed layeron the surface region of tungsten [3-5]. Hydrogenisotopes including tritium will be implanted
simultaneously with formation of W-C mixed layer, andwill dynamically interact with carbon in tungsten. In theview point of the fuel recycling, it is important toelucidate the retention behavior of hydrogen isotopes
retained on tungsten materials in the fusion environment.
In our previous studies, deuterium retention behavior intungsten with simultaneous 10 keV C+ implantation was
evaluated. Enhancement of deuterium retention by theauthor’s e-mail: syoya@ipc.shizuoka.ac.jp

76©2013 by The Japan Society of Plasma
Science and Nuclear Fusion ResearchY. Oya et al., Implantation Energy Dependence on Deuterium Retention Behaviors for the Carbon Implanted Tungsten

carbon was implanted into tungsten prior to deuterium. In
the sequential implantation, deuterium was implanted into
tungsten after formations of irradiation defects and W-C
mixed layer. In addition, implantation energy of carbon
was changed to investigate dependences of the
distribution of carbon and damage on deuterium retention
behavior. The deuterium retentions were evaluated bymeans of Thermal Desorption Spectroscopy (TDS)
measurements. The irradiation damages were also
observed by Transmission Electron Microscopy (TEM) at
Kyushu University.

2. Experimental
A disk-type polycrystalline tungsten with stress-relieved
conditions (heated at 1173 K), which was supplied byAllied Material Co. Ltd was used. The sample size was
10 mm diameter and 0.5 mm thickness. The Sample waspolished mechanically to the roughness of less than 1 μm
by SiC abrasive papers and mono-crystalline diamondsuspensions, and preheated at 1173 K for 30 minutes
under ultrahigh vacuum to remove the impurities anddamages introduced during the polishing processes. Then,simultaneous or sequential carbon ion (C+) and deuteriumion (D2+) implantations were performed in the C+-D2+
simultaneous implantation apparatus at room temperature.
In this apparatus, the sample can be irradiatedindependently by two different ion guns, namely C+ andD2+ guns. In the C+ gun, CO2 was used as the C+ sourcegas to get rid of hydrogen impurities. The E × B mass
separator was equipped at the head of the C+ gun to
extract pure C+. The incident angles for the C+ and D2+guns were 15 degrees from the sample surface normal,
forming an angle of 30 degrees between the two beams.
The energy of C+ was changed from 5 to 10 keV and thatof D2+ was set to be 2 - 3 keV. The C+ and D+ flux were
0.1 × 1018 C+ m-2 s-1 and 1.0 × 1018 D+ m-2 s-1,
respectively. Therefore, for C+-D2+ simultaneous
implantation, the ion flux ratio was set to be C+/D+ = 0.1.
The C+ fleunce was constant to be 0.1 × 1022 C+ m-2 in
C+-D2+ sequential implantation. D+ fluence was changedin the range of (0.1 - 1.8) × 1022 D+ m-2. The only D2+
implantation was also performed to compare thedeuterium retention in C+ implanted tungsten. The
chemical states of carbon and tungsten were studied by
X-ray Photoelectron Spectroscopy (XPS). The TDS
measurements were performed from room temperature to1173 K with a heating rate of 0.5 K s-1. The TEM
observations were also performed to investigate theirradiation damages formed by the ion implantations. The
depth profile of deuterium was analyzed by Glow
Discharge Optical Emission Spectroscopy (GD-OES).

Fig. 1 TEM images of tungsten implanted with 3 keV
D2+ or 10 keV C+ at various ion fluences.
keV D2+ with the fluence of 1.0 × 1022 D+ m-2 andtungsten with 10 keV C+ with different C+ fluences in therange of (0.001-0.1) × 1022 C+ m-2. For all the samples,lots of dislocation loops were introduced in tungsten byion implantations, which was quite different from puretungsten where its density was low. In addition, since theirradiation vacancy and dislocation loop are pair defects,
the formation of dislocation loops indicated theintroduction of vacancies. It was also found that evenhigher fluence of D+, density of irradiation defects intungsten with 10 keV C+ implantation was higher thanthat with 3 keV D2+ implantation, indicating that
irradiation damages were easily introduced in tungsten by
10 keV C+ implantation compared to that by 3 keV D2+
implantation due to higher implantation energy, higher
mass and larger atomic radius. It was also found thatdensity of irradiation damages in tungsten induced by 10
keV C+ implantation was increased in the C+ fluence
range of (0.001-0.01) × 1022 C+ m-2, but its density wasseen to be the same with C+ fluence larger than 0.01 ×

1.5

 C-C bond (XPS)

D (GD-OES)

1.0

30
25
20
15

0.5
0.0

10
5

20
30
40
50
depth / nm
Fig. 2 Depth profiles of D observed by GD-OES andW-C and C-C bonds by XPS. The TEM image wasalso shown in this figure.

3. Results and discussion
Fig. 1 show TEM images of tungsten implanted by 3

77

0

10



 W-C bond (XPS)XPS Intensity / arb. unit

GD-OES Intensity/ arb. unit20nm

Y. Oya et al., Implantation Energy Dependence on Deuterium Retention Behaviors for the Carbon Implanted Tungsten

Peak 2

Peak 1
4.0

Peak 322

D fluence (x 10 )+-2
0.3 D m+-2
1.8 D m-2
17+

Only D2 imp.
+

Desorption rate / 10 D2 m s-11022 C+ m-2, indicating that density of irradiation defects
were saturated in the C+ fluence of 0.01 × 1022 C+ m-2. By
the SRIM (Stopping range of Ions in Matter) code
calculation [10], irradiation defects in tungsten implantedby 10 keV C+ with the C+ fluence of 0.01 × 1022 C+ m-2
were distributed within 30 nm beneath the surface, withthe peak at around 5 nm (1.3 dpa), indicating that thedensity of irradiation defects would be saturated below
1.3 dpa. The depth profiles of D observed by GD-OES at
University of Toyama and, W-C and C-C bonds
evaluated by XPS were shown in Fig. 2. The TEM image
was also overlapped in this figure. It was found that mostof D was retained in W-C mixed layers and major
chemical state of C in this layer was found to be W-C
bond.
Fig. 3 shows D2 TDS spectra for only 3 keV D2+
implanted tungsten and 10 keV C+ - 3 keV D2+
sequentially implanted one with various D+ fluences.
These spectra were consisted of three desorption stages ataround 370, 480 and 620 K which were named as Peaks 1,
2 and 3, respectively. In lower D+ fluence of 0.1 ~ 0.3 x1022 D+ m-2, the deuterium retentions as Peaks 1 and 2
were clearly increased by C+ implantation, indicating thatthese desorption stages would be related to the deuterium
interacted with the irradiation damages. The density ofirradiation damages induced by 10 keV C+ implantation
was higher than that by 3 keV D2+ implantation as shownin Fig. 1, resulting in the higher deuterium retention for
the 10 keV C+ implanted tungsten. In general, trapping
energy of deuterium by vacancy would be higher thanthat by dislocation loop [11,12], indicating that Peaks 1
and 2 could be assigned to the desorption stages of
deuterium trapped by dislocation loops and vacancies,
respectively. In higher fluence of 1.8 × 1022 D+ m-2, thedeuterium retentions as Peaks 1 and 2 for C+ implanted
W were almost the same as that without C+ implantation,
indicating that deuterium trappings by irradiation
damages should be almost saturated. However, thedeuterium retention as Peak 3 for only D2+ implanted
tungsten was higher than that for 10 keV C+ - 3 keV D2+
sequentially implanted one. The amount of trapping site
as Peak 3 was not induced by ion implantation because
the retention of deuterium as Peak 3 was not increased by
10 keV C+ implantation. Therefore, the promisingtrapping site for Peak 3 would be intrinsic sites such as
cavities, grain boundary and/or oxygen impurity [13,14].
Lower deuterium retention as Peak 3 for 10 keV C+ - 3
keV D2+ sequentially implanted tungsten would indicatethat deuterium diffusion in tungsten was refrained by 10
keV C+ implantation.
The retentions of deuterium retained as Peaks 1, 2 and 3
for tungsten with 10 keV C+ - 3 keV D2+ sequential
implantation or simultaneous implantations werecompared in Fig. 4. The deuterium retentions for the
samples with the D+ fluence of 1.8 × 1022 D+ m-2 were

3.0

+

+

C -D2 seq. imp.

2.0

+

22

D fluence (x 10 )+-2
0.1 D m+-2
1.8 D m

1.0
0.0
300

500
700
Temperature / K

900Fig. 3 D2 TDS spectra of 3 keV D2+ implanted
tungsten and 10 keV C+- 3 keV D2+ implanted one.

1.5

Peak1
Peak2
Peak3

20

Deuterium retention / 10 D m-2compared. In the case of C+- D2+ simultaneous
implantation, the retention of deuterium retained as Peak
3 was decreased by C+ implantation, which was almost
the same with both of sequential and simultaneous
implantations. However, the deuterium retentions as
Peaks 1 and 2 for the simultaneous implantation wereincreased compared to those for the sequential
implantation. The deuterium trappings as Peaks 1 and 2
for C+- D2+ simultaneously implanted tungsten werehigher than that for sequentially implanted tungsten,
indicating that simultaneous implantation effects were
contributed to the increase of deuterium retentions inthese peaks.
To clarify the simultaneous implantation effects on
deuterium retention behavior in tungsten, incident energy
of C+ was changed for both of sequential or simultaneous
implantations. The implantation energy of D2+ was fixed
to be 2 keV. Figs. 5 (a) and (b) show the D2 TDS spectra
for C+ - D2+ simultaneously implanted tungsten and C+ D2+ sequentially implanted one as a function of C+
incident energy, respectively. Both of D+ and C+ fluenceswere set to be 1.0 x 1022 D+ m-2 and 0.1 x 1022 C+ m-2,
respectively. The D2 TDS spectra were slightly changed
in simultaneous implantation case although little changes

1.0

0.5

0.0

+

Only D2++

C -D2 seq.

++

C -D2 sim.

Fig. 4 Comparison of deuterium retentions as Peaks 1, 2
and 3 for tungsten implanted with various implantation
procedure.

78

6.0

+

+

+

+

+

-2

5.0

+

7 keV C -2 keV D2

17

5 keV C -2 keV D2

4.0

+

3.0

only 2 keV D22.0(a) C+-D2+ sim. imp

1.0
0.0
300

500
700
Temperature / K

6.0

-1

10 keV C -2 keV D2Desorption rate / 10 D2 m s

17-2

Desorption rate / 10 D2 m s-1

Y. Oya et al., Implantation Energy Dependence on Deuterium Retention Behaviors for the Carbon Implanted Tungsten

900

+

+

10 keV C -2 keV D2

5.0

+

+

+

+

7 keV C -2 keV D2

4.0

5 keV C -2 keV D2

3.0+

only 2 keV D22.0(b) C+-D2+ seq. imp

1.0
0.0
300

500
700
Temperature / K

900Fig. 5 D2 TDS spectra of tungsten with (a) C+- 2 keV D2+simultaneous and(b) C+- 2 keV D2+ sequential implantation with various C+ incident energies

2.0

-2

(a) C+-2 keV D2+ sim. imp

20

2.5

Peak1
Peak2
Peak3

1.5
1.0
0.5
0.0

the simultaneously implanted tungsten was resulted from
distribution of irradiation damages. Deuterium trappings
by irradiation defects in deeper region (20 - 30 nm
beneath surface) would be contribution to these increasesof deuterium retention as Peaks 1 and 2. However in thecase of C+ - D2+ sequential implantation, deuterium
retention was not increased with increasing C+ incident
energy, considering that retained carbon would prevent
deuterium penetration into bulk region. In C+ - D2+
sequential implantation, the 2 keV D2+ was implanted
into tungsten after C+ implantation. Therefore, implanted
deuterium should be affected by retained carbon intungsten. On the other hand in C+- D2+ simultaneous
implantation, several portion of deuterium was implanted
into tungsten before retaining carbon enough to prevent
deuterium diffusion into bulk.
In summary, the retention of deuterium trapped byirradiation damages was enhanced by energetic carbon
implantation. The deuterium retention by intrinsic cavity
in bulk region was also found and was decreased byenergetic carbon implantation. Therefore, it wasDeuterium retention / 10 D m

20

Deuterium retention / 10 D m-2

were observed in that for C+- D2+ sequentially implanted
tungsten. Fig. 6 summarizes the retentions of deuterium
retained as Peaks 1, 2 and 3. The deuterium retention as
Peak 1 was clearly increased for C+- D2+ simultaneously
implanted tungsten. It was found that the no difference of
deuterium retention for C+- D2+ sequentially implanted
tungsten was observed nonetheless C+ incident energy
was changed from 5 to 10 keV. By changing the C+
incident energies from 5 to 10 keV C+, the density ofirradiation damages was estimated to be above 10 dpa,
indicating that the density of irradiation defects was
saturated in all the samples. However, distributions ofirradiation defects induced by C+ implantation werechanged with incident energy from SRIM code
calculation. With increasing C+ incident energy,
distribution of irradiation damages was expanded toward
the depth. The irradiation damages induced by 10 keV C+
were extended from surface to 30 nm of depth althoughthose by 5 keV C+ were limited with 20 nm from thesurface, indicating that the increase of deuterium
retention with increasing C+ incident energy observed for

0

2
4
6
8
+
C incident energy / keV

10

2.5
2.0

(b) C+-2 keV D2+ seq. imp.
Peak1
Peak2
Peak3

1.5
1.0
0.5
0.0

0

2
4
6
8
+
C incident energy / keV

Fig. 6 Deuterium retention for each peak of tungsten with (a) C+- D2+simultaneous and(b) C+- D2+ sequential implantation with various C+ incident energies.

79

10

Y. Oya et al., Implantation Energy Dependence on Deuterium Retention Behaviors for the Carbon Implanted Tungsten

considered that implanted carbon would play a role as thedeuterium diffusion barrier leading the reduction of
deuterium penetration toward the bulk. However, several
experiments showed deuterium ion-driven permeation
was promoted by carbon simultaneously implanted into
tungsten. These facts can be explained by the project
range of carbon and deuterium ions. In our experiments,
the project range of 5 - 10 keV C+ were almost the sameas that of 2 - 3 keV D2+. As the project range of carbon
reported in ref. [8,9] was shorter than that of deuterium,
higher portion of deuterium was implanted deeper regionof tungsten than carbon. The implanted deuterium waseasily permeated through tungsten because of thedecreased diffusivity in the implanted surface by carbon.

[14]

4. Conclusions
The effects of energetic carbon implantation on
deuterium retention behavior in tungsten were
investigated in this study. Dislocation loops andvacancies were formed in tungsten as irradiation defects
by energetic carbon implantation. The density ofirradiation damage was saturated below about 1 dpa by
TEM observations. Deuterium retentions by dislocation
loops and vacancies were observed in TDS measurements.
Deuterium retentions by irradiation defects wereincreased by energetic carbon implantation. Deuterium
retention by intrinsic cavity in bulk region was also found
and was decreased by energetic carbon implantation. It
was concluded that retained carbon would play a role as
the diffusion barrier of deuterium, preventing deuterium
to penetrating into bulk.
Acknowledgement
This study was supported by JSPS Kakenhi No.
22360389, NIFS collaboration program NIFS11KEMF022,
University of Toyama collaboration program and theCenter for Instrumental Analysis at Shizuoka University.

References[1]

R. P. Doerner et al., J. Nucl. Mater. 363-365, 32-40(2007).
[2] S. Pestchanyi et al., Fusion Eng. Des. 66-68, 271-276(2003).
[3] I. Bizyukov, et al., J. Nucl. Mater., 363-365, 1184-1189(2007).
[4] V. Kh. Alimov, Phys. Scr., T108, 46-56 (2004).
[5] H.T. Lee, et al., J. Nucl. Mater. 390-391, 971-974 (2009).
[6] M. Kobayashi, et al., Phys. Scr. T138, 014050 (2009).
[7] V. Kh. Alimov, et al., J. Nucl. Mater., 375, 192 (2008).
[8]
Y. Ueda et al., J. Nucl. Mater., 2012 in press.
[9]
H. Y. Peng et al., Phys. Scr. T145, 014046 (2011).
[10] J.F. Ziegler, 2006 SRIM code.
[11] M. Poon et al., J. Nucl. Mater., 374, 390 (2008)[12] J. Roth and K. Schmid Phys. Scr. T145, 014031 (2011).
[13] V.Kh. Alimov et al., J. Nucl. Mater. 337–339, 619–623

80(2005).
T. Ahlgren et al., J. Nucl. Mater. In press.

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Tritium distribution on the tungsten surface exposed to deuterium plasma
and then to tritium gas
Kanetsugu ISOBE1)*, Vladimir Kh. ALIMOV1.2), Akira TAGUCHI2), Makiko SAITO1,3),
Yuji TORIKAI2), Yuji HATANO2), Toshihiko YAMANISHI1)

1) Tritium technology group, Japan Atomic Energy Agency, shirakatashirane 2-4, Tokai-mura, Ibaraki 319-1195, Japan.
2) Hydrogen Isotope Research Center, University of Toyama, Gofuku 3190, Toyama-shi, Toyama 930-8555, Japan.
3)ITER project Management group, Japan Atomic Energy Agency, mukooyama801-1, Naka-shi, Ibaraki 319-0193, Japan.
(Received: 12 May 2012 / Accepted: 5 October 2012)

Tungsten (W) samples were exposed to Low-energy (38 eV/D), high flux (1022 D+/m2/s) deuterium (D) plasma
at 495, 545 and 550 K to a fluence of 1026 D/m2. After the plasma exposure, tritium was introduced into the samples
by exposure to deuterium-tritium gas mixture at 473K. Tritium distribution on the W surface was examined by thetechniques of imaging plate and autoradiography. The results of the imaging plate revealed that tritium wasconcentrated mainly within the area exposed to the D plasma, and the concentration of tritium is slightly varied
within the area. Autoradiography of the W surface showed that tritium was concentrated on the grain boundary andblisters.
Keywords: Tungsten, Plasma-surface interaction, Tritium, Imaging plate, Tritium autoradiography.

10 u 10 u 2 mm3 and polishing. The grain size was
20-200 ȝm.
The linear plasma generator used for delivering aplasma beam is described elsewhere [4]. A plasma beam
highly enriched with a species of D2+ to over 80% was
obtained. The bias voltage of -80 V was applied to the W
sample resulting in a incident of 38 eV/D, taking into
account the plasma potential of about -4 V measured by aLangmuir probe. The incident deuterium ion flux was
fixed at 1022 D+/m2/s, and the samples were exposed toion fluence of 1026 D/m2. The sample was passively
heated by the plasma itself and exposure temperature wasset by the thermal contact between the sample and thecooled holder. D plasma exposure to recrystallized W was
carried out at around 495, 545 and 550K. At the fluenceof 1026 D/m2, the deuterium inventory of tungsten
exposed at temperature over the range of 500 - 550 K
become higher than that at other temperature range [15].
The surface area of the W samples exposed to the plasma
was 9 mm in diameter. For a comparison, as-received
recrystallized W sample was also used for tritium loading.
Note that the as-received sample was not exposed to the
D plasma.
4 months after the D plasma exposure, the samples
were loaded with tritium from gaseous phase. Fig. 1
shows a schematic view of a tritium exposure apparatus.
The tritium exposure apparatus consisted of a vacuum
pumping system, a quartz tube for gas exposure equippedwith an external heater, quadrupole mass spectrometer

1. Introduction
Since tungsten (W) is one of the candidate materials
for plasma-facing components of fusion reactor, many
investigations relating to the plasma-W interaction havebeen carried out. From a view point of safety, the tritium
inventory in plasma-facing materials is a one of key
issues due to limitation of tritium inventory in thevacuum vessel. Most of the results on hydrogen isotope
retention and recycling in W materials have beenreviewed by Causey and Venhaus [1,2], and Skinner et al[3]. It has been reported that during exposure to D
plasmas, blisters with various shapes and sizes dependingon the exposure temperature are formed on the W surface[4-14]. Blister formation is accompanied by deformationof the near-surface layers [12, 13], and this deformation isthought to create additional traps for hydrogen isotopes.
There is also a possibility that amount of hydrogenisotope which could be solved or trapped in the deformed
W near-surface layer increases. In this study, W samples
with D-plasma-induced blisters on the surface were
exposed to D2-T2 gas mixture, and the distribution oftritium on the W surface was examined by the techniquesof imaging plate and autoradiography.

2. Experimental
The W specimens exposed to D plasma wereprepared from polycrystalline tungsten (A.L.M.T. Corp.,
Japan) with a purity of 99.99 wt%. The samples were
fully recrystallized at 2070 K for 1 hour after cutting into
author’s e-mail: isobe.kanetsugu@jaea.go.jp

81©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

K. Isobe et al., Tritium Distribution on the Tungsten Surface Exposed to Deuterium Plasma and then to Tritium Gas

to expose the emulsion to the tritium beta electrons. Then,
specimens were dipped in a developer solution and thenin a fixer solution. Under exposure to beta electrons, theAgBr grains were transformed into Ag grains. The
distribution of Ag grains was observed by the scanningelectron microscope (SEM).

(QMS), and a tritium source. The W samples were placedin the quartz tube, in doing so pieces of glass wool wereinserted between samples to avoid a contact between
them. The quartz tube with samples was heated in
vacuum at 473K for degassing until termination of gas
release that was controlled by QMS. At this temperature,
we have already confirmed that deuterium in tungsten is
not released to the vacuum by thermal desorption
spectroscopy [15]. When high vacuum was obtained,
tritium diluted with deuterium was introduced into thequartz tube from the tritium source. W samples were
exposed at 473 K to D (92.2%)-T (7.8%) gas mixture at a
total gas pressure of 1 kPa for 5 hours.

QMS

IG

3. Results and Discussion
Surface morphology of the recrystallized W after D
plasma exposure at 495K is shown in Fig. 2. Grainboundary could be clearly observed in Fig. 2. Largeblisters with sizes of a few tens of microns and small
blisters with size of less than a few microns are observed.
The large blisters demonstrate a multi-layered structure
like steps. The small blisters are formed on the flat
surface of crystal grains.

CM
Chamber

495K
TMP

SIP

RP
Tritium source Specimen
Fig. 1 Schematic view of tritium exposure apparatus.
QMS:Quadropole mass spectrometer,
IG:Ionization gauge, CM:Capacitance manometer
SIP:Suptter ion pump, RP:Rotary pump
TMP:Turbo molecular pump

20μm
Fig. 2. SEM image of recrystallized W exposed at 495 K
to low-energy high flux D plasma with fluence of 1026
D/m2. The surface was tilted at an angle of 45ºto theelectron beam.

The amount of tritium retained in the W near-surface
layers was evaluated by the imaging plate technique.
Imaging plate (IP) is a 2-dimensional radiation detector
with high sensitivity. IP used in our experiments was
BAS-TR2025 (Fuji Photo Film Co. Ltd.) which consistsof 50 ȝm thickness photostimulable phosphor mounted
on a plastic support. The IP surface was contacted to the
W surface and exposed to tritium beta electrons in 1 week.
After the exposure to tritium beta electrons, the IP wasprocessed by an IP reader FUJIFILM BAS2500 (Fuji
Photo Film Co. Ltd.) in order to obtain the digital image.
Tritium activity is expressed as the photo-stimulate
luminescence (PSL) intensity, which is proportional to theabsorbed radiation energy.
Tritium autoradiography technique was applied totwo specimens; recrystallized W exposed to the D plasma
at 495K and as-received, non-exposed recrystallized W.
At first, to avoid direct contact of the W samples with
AgBr grains in radiographic emulsion, thin collodion film
was deposited on the W surface by dipping the samples incollodion diluted by ethanol. The monolayer of theradiographic emulsion (ILFORD L4) was placed onto thecollodion film by a wire-loop method. The specimens
were kept in a light-tight box at around 273 K for 12 days

Fig. 3 shows the results of IP measurements. IP
measurement can easily observe tritium distribution of
the whole of specimen as shown in Fig. 3. A level ofdarkness corresponds to a concentration of tritium innear-surface layer. As seen Fig. 3 (a, b, c), tritium isconcentrated on the surface area exposed to the D plasma.
It is known that under exposure to high flux D plasmas,
vacancy-type defects, voids and cracks are generates in W,
and these defects are responsible for accumulation
deuterium [11-14, 16]. Obviously, subsequent exposure totritium leads to isotope exchange in the near-surface layer[17] and accumulation of tritium at the defects createdunder previous D plasma exposure.
The average tritium concentration of the W surface
exposed at 550 K is significantly higher than that for the
W surface exposed at 545 K, although the difference inthe exposure temperature is only 5 K. Obviously, a small
variation in the exposure temperature within thistemperature range has a significant influence on a82

K. Isobe et al., Tritium Distribution on the Tungsten Surface Exposed to Deuterium Plasma and then to Tritium Gas

concentration of the D-plasma-induced traps generated innear-surface layer. This suggestion is supported by sharp
temperature dependence of the D retention in therecrystallized W exposed to low-energy, high-flux D
plasma [16].
(a)

(b)

(1)

(b)

100μm

(4)
(d)

(6)

Grain
boundary

(3)

(2)
(c)

(a)

(b)

(5)

BlisterFig. 3. IP images for recrystallized W exposed to
low-energy, high flux D plasma at 495 K (a), 545 K (b),550 K (c), and for as-received recrystallized W (d).
Dashed circles (a-c) indicate the surface areas exposed to
D plasma. Tritium was loaded into all samples at 473 K,
1 kPa D-T gas mixture for 5 hours.
However, within the exposed surface area the tritium
distribution is not uniform. Intensities of PSL at points((1)-(6)) in Fig. 3 are shown in Table I. The ratio between
the maximum and minimum PSL intensities on the same
exposed sample surface ((2)/(1), (4)/(3), (6)/(5)) varied
from 4:1 to 6:1. It was reported recently that hydrogenisotopes can be retained by tungsten oxide layers [18].
Probably, the surfaces of the W samples werenonuniformly oxidized before tritium gas exposure.

10μm

(c)

Table I Intensity of photo-stimulate luminescence atsome points on the plasma-exposed W surface
Point
PSL/mm2

(1)
320

(2)
1800

(3)
500

(4)
2100

(5)
560

20μm

(6)
3200

Figure 4 demonstrates autoradiographic patterns, inwhich white spots correspond to Ag grains tracing thetritium beta electrons. The Ag grains, as white lines, wereobserved along grain boundary (Fig. 4 (a)). The Ag grainsdecorate also blisters and are distributed uniformly on thegrain surface (Fig. 4 (b, c)). However, the Ag grains are notobserved on the surface of as-received recrystallized W(Fig. 4 (d)). As was mentioned above, various types oftrapping sites for hydrogen isotopes are generated in W
under exposure to high-flux D plasma. Under subsequent
exposure to tritium gas at elevated temperature, tritium isthought to be occupied these trapping sites. It should be
noted that techniques of imaging plate and autoradiography
allow detecting tritium retention in a shallow, near-surface
layer because the escape depth of tritium beta electrons in

(d)

30μm
Fig. 4. Autoradiographic patterns of recrystallized W
exposed with D plasma at 495K (a, b, c) and as-received
recrystallized W (d). Tritium was loaded into thesesamples at 473 K, 1 kPa D-T gas mixture for 5 hours.

83

K. Isobe et al., Tritium Distribution on the Tungsten Surface Exposed to Deuterium Plasma and then to Tritium Gas

W does not exceed one micrometer [19, 20]. Thus, resultsof this study shows that under exposure to the D plasma atelevated temperatures, traps in the near-surface layer are
generated mainly along grain boundaries and in layers
subjected to tensile stress due to material migration above
the surface during formation of blisters [12].

report of Hydrogen Isotope Research Center, Toyama
University, 29-30(2010)71V.
[16] V.Kh. Alimov, W.M. Shu, J. Roth et al., J. Nucl. Mater.,
417(2011)572.
[17] Y. Hatano, M. Shimada, V.Kh. Alimov et al., Proc. 20th
Intern. Conf. on Plasma Surface Interactions, May 2012,
Achen(submitted to J. Nucl. Mater.)[18] V.Kh. Alimov, B. Tyburska, M. Balden et al., J. Nucl.
Mater., 409(2011)27.K. Miyasaka, T. Tanabe, G. Mank andet al., J. Nucl. Mater., 290-293(2011)448[19] T. Hoshihira, T. Otsuka and T. Tanabe, J. Nucl. Mater.,
386-388(2009)776[20] T. Hoshihira, T. Otsuka and T. Tanabe, J. Nucl. Mater.,
390-391(2009)1029

4. Conclusion
Techniquesofimaging
plate
(IP)
and
autoradiography were applied to observe a distribution oftritium on the W surface preliminary exposed to
low-energy, high-flux D plasma at elevated temperatures(495-550 K) and then to tritium gas at 473 K. According
to IP images, tritium was highly concentrated within
D-plasma-exposed surface area; however, the tritium
concentration was varied within this area. With the use of
autoradiography, it was found that tritium wasconcentrated on the grain boundary and along/on blisters
formed under D plasma exposure. These results suggestthat D plasma exposure generates traps for hydrogenisotopes in the W near-surface layer.

Acknowledgements
A part of this work is performed with the supportand under the auspices of NIFS (NIFS11KUHR012) anda KAKENHI Grant-in-Aid for Scientific Research (B)(Project No. 21360458, No. 24360380). Authors reallywish to acknowledge for these supports.

5. References[1]

R. A. Causey and T. J. Venhaus, Phys. Scripta T94 (2001)
9.
[2] R. A. Causey, J. Nucl. Mater. 300 (2002) 91.
[3] C. H. Skinner, A. A. Haasz, V. Kh. Alimov, et al., Fusion
Sci. Technol. 54 (2008) 891.
[4] G.-N. Luo, W.M. Shu, H. Nakamura et al., Rev. Sci.
Instrum. 75 (2004)4374.
[5] G. -N. Luo, W. M. Shu, M. Nishi, J. Nucl. Mater. 347(2005) 111.
[6] G. -N. Luo, W. M. Shu, M. Nishi, Fusion Eng. Design 81(2006) 957.
[7] W. M. Shu, A. Kawasuso, Y. Miwa et al., , Physica
Scripta T128 (2007) 96.
[8] V. Kh. Alimov, J. Roth, Physica Scripta T128 (2007) 6.
[9] W. M. Shu, G. -N. Luo, T. Yamanishi, J. Nucl. Mater.
367-370 (2007) 1463.
[10] W. M. Shu, E. Wakai, T. Yamanishi, Nucl. Fusion 47(2007) 201.
[11] V. Kh. Alimov, W. M. Shu, J. Roth and et al., Physica
Scripta T138 (2009)014048.
[12] S. Lindig, M. Balden, V.Kh. Alimov et al., Phys. Scripta
T138(2009)014040.
[13] S. Lindig, M. Balden, V.Kh. Alimov et al., Phys. Scripta
T145(2011)014039.
[14] V. Kh. Alimov, B. Tyburska-Püschel, S. Lindig and et al., J.
Nucl. Mater., 420(2012)519[15] V.Kh. Alimov, S. Lindig, M. Balden and et al., Annual84

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Study on plasma – tungsten surface interactions using the new experimental
device EXPRESS
Ryo Miura1)*, Tetsuo Fujishima1), Yuto Miyahara1), Tomohisa Taguchi1),
Masahiro Tanaka2), Akio Sagara2),Yasuhisa Oya 1), and Kenji Okuno1)
1) Radioscience Research Laboratory, Faculty of Science, Shizuoka University
2) National Institute for Fusion Science(Received: 25 May 2012 / Accepted: 5 December 2012)

An experimental apparatus named as EXPRESS (Effluence and Exchange Probe for Recycling Estimation at
NIFS and Shizuoka University) has been designed and established in Shizuoka University to acquire laboratory data(well-conditioned data) needed for the design of an actual reactor. The behaviors of deuterium retention and thesurface chemical state after EXPRESS plasma operation were evaluated and compared with that for LHD (Large
Helical Device, NIFS) hydrogen plasma exposed sample. It was found that a thin impurity deposition layer with the
thickness less than 0.3 nm was formed on the surface, which was one thirteenth thinner than that on LHD plasma
exposed tungsten. The TDS experiment showed that major deuterium desorption stage for EXPRESS plasma
exposed tungsten was located at the lower temperature less than 550 K, although that for LHD plasma exposed one
was found at the temperature between 550 K – 1000 K. These facts suggest that the impurity concentration in thedeposition layers have a large influence on the deuterium retention behavior in tungsten.
Keywords: deuterium plasma exposure, TDS, XPS, impurity deposition layer, tungsten(Large Helical Device, NIFS) plasma exposed tungsten
and only D2+ implanted tungsten for the evaluation of fuelrecycling during plasma-surface interactions in
EXPRESS.

1. Introduction

Tungsten will be used as a plasma facing material in
future fusion devices. Many studies have been devoted toelucidate fundamental tritium retention and transport
behaviors in tungsten under well-controlled condition [1-3].
However, it is quite difficult to apply these data obtained
under well-controlled condition to actual and larger-scale
plasma devices or fusion devices where conditions aremuch more complex. For example, a wall temperature
gradient will arise when the plasma is controlled. Therefore,
the key parameters should be acquired to complete
laboratory data (well-controlled condition data) needed forthe design of an actual reactor design. The design of a new
plasma experimental device, namely EXPRESS (Effluenceand Exchange Probe for Recycling Estimation at NIFS and
Shizuoka University) was designed to clarify the hydrogenisotope recycling dynamics. This devise has thesefollowing features: i) many samples can be placed on thesample stage for plasma exposure and the sample
temperature gradient can be controlled, and ii) 1 GBq / dayof tritium can be handled.
As a preliminary experiment, the deuterium plasma
exposure to pure tungsten was performed in this study tocharacterize and evaluate the behaviors of the deuterium
retention and the surface chemical state after EXPRESS
plasma operation at R.T. These experimental results
obtained at EXPRESS were compared with those in LHD

2. EXPRESS Apparatus

The recently designed plasma experimental device toinvestigate hydrogen isotope recycling behavior andplasma surface interaction, EXPRESS consists of threeparts: i) a plasma generator and plasma exposure system,ii) a fuel storage and purification system, and iii) a fuelrecovery-supply systems. Figure 1 shows the schematic
view and picture of the plasma generator and plasma
exposure system. The plasma generator and plasma

Fig. 1 Schematic view and picture of the plasma
generator and plasma exposure system in
EXPRESS.

e-mail: r0232027@ipc.shizuoka.ac.jp85©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

R. Miura et al., Study on Plasma-Tungsten Surface Interactions Using the New Experimental Device EXPRESS

exposure system can be evacuated to 1 × 10-5 Pa by a scroll
pump (SP) and a turbo-molecular pump (TMP)
combination. The pressure in the system is measured by
Baratron and ionization gauge (I.G.). The source gas issupplied through a mass flow controller (M.F.C.) at a flowrate of 0.1-5.0 sccm into the plasma generation tube madeby quartz with a size of 150 mm in length and 42 mm indiameter. The excitation frequency is fixed at 13.56 MHz
by a RF matching box. The bias voltage can be changed
from 0- 200 V between the chamber wall and the sample
stage to control implantation energy similarly larger-scale
plasma devices. The size of sample stage is 100 mm indiameter, which is made of tungsten. A heating unit isinstalled beneath the center of sample stage to generate atemperature gradient up to 200 K between the center andthe edge of the sample stage as shown in Fig. 2. Thisapparatus is used to investigate the study ofplasma-surface interactions on the first wall of actual
reactor. The heating unit can heat up to 873 K at the centerof stage. The ion flux can be measured by a faraday cup
installed in the plasma exposure system.
㻌900

at maximum heating

850
800

㻌

Temperature / K

energy for the stable plasma operation. The ion flux was
3.4 × 1019 D+ m-2 s-1and the fluence was 1.0 × 1024 D+ m-2
at R.T. This sample was named as EXPRESS sample.
After plasma exposure, TDS (Thermal Desorption
Spectroscopy) were performed from R.T. to 1173 K with aheating rate of 0.5 K s-1 to investigate the deuterium
desorption and retention behaviors. The chemical states
and the thickness of the impurity layer on tungsten were
evaluated by XPS (X-ray Photoelectron Spectroscopy)(ESCA1600 system, ULVAC-PHI Inc.) using a Mg - KD
X-ray source (1253.6 eV) and 1 keV Ar+ sputtering
technique [4]. The thickness of impurity deposition layer
was defined as the depth where impurity concentration wasreached to be half compared to that on the surface. The
sputtering rate by Ar+ implantation was estimated by AFM(Atomic force microscope).
Two types of tungsten samples, namely D2+ imp.
sample and LHD sample, were prepared to compare thedeuterium retention behavior. In the D2+ imp. sample, 3.0
keV D2+ implantation was performed for tungsten withflux of 1.0 × 1018 D+ m-2 s-1 and fluence of 1.0 × 1022 D+
m-2 at R.T. In the case of LHD sample, the sample was
exposed to 134 shots (exposure time 271 seconds) ofhydrogen plasma in LHD, with the ion energy of about
10-100 eV and the temperature of less than 373 K, at afirst wall position in the plasma operation campaign of
2011.

750

4. Results and discussion

Figure 3 shows the TDS spectra for EXPRESS and
LHD and D2+ imp. samples. The TDS spectrum for the D2+
imp. sample consists of two desorption stages at 400 K and
550 K, attributing to the desorption stages of deuterium
adsorbed on the surface and/or trapped by dislocation loops(Peak 1), [5,6] and deuterium trapped by vacancies (Peak2) [7] respectively. In the D2 desorption stages for the
EXPRESS sample, Peak 1 was the dominant stage which issimilar for the D2+ imp. sample, but the ratio of Peak 1 to

700
650
0

5 10 15 20 25 30 35 40 45 50

Distance from center / mm

Fig. 2 The maximum temperature profile on the sample
stage in EXPRESS.

3. Experimental

Peak 1
-1
-2
17

1.5

EXPRESS sample
LHD sample+
D2 imp. sample

3.0
2.5
2.0

Peak 3

1.5

1.0

1.0
0.5

Peak 4

0.5

-2

500

700
900
Temperature / K

1100

0.0

-1

0.0
300

Fig. 3 TDS spectra for the EXPRESS, LHD
and D2+ imp. samples

86

18

Desorption rate / 10 D m s

2.0

Peak 2

Desorption rate / 10 H m s

A disk-type polycrystalline tungsten withstress-relieved conditions (heated at 1173 K for 1 hour),which was purchased by Allied Material Co. Ltd, wasused for all experiments. The sample size was 10 mmdiameter and 0.5 mm thickness. The samples were
polished mechanically to the roughness of less than 1 Pm
by SiC abrasive papers and 1-3 Pm diamond suspensions,and preheated at 1173 K for 30 minutes under ultrahighvacuum (< 10-6 Pa) to remove the impurities and damages
introduced during the polishing processes.
After preheating, the sample was introduced into theplasma exposure system of EXPRESS. The deuterium gas
was supplied at a flow rate of 3.0 sccm and the pressure of
3.5 Pa. The D plasma was exposed to the samples with theD2+ energy of 63 eV to implant at the lowest possible

R. Miura et al., Study on Plasma-Tungsten Surface Interactions Using the New Experimental Device EXPRESS

surface and thickness of impurity deposition layers on the
EXPRESS and LHD samples. Large amount of oxygenand carbon were accumulated on the surface for both
samples and its impurity concentrations were almost thesame. However, the thickness of the impurity deposition
layer for the EXPRESS sample was about one thirteenththinner than that for the LHD sample. It can be said thatthe thin impurity deposition layer was achieved by
EXPRESS device and the formation of impurity deposition
layer hardly affected on deuterium retention as O-D bonds
for the EXPRESS sample, resulting in decrease of Peak 3
in the TDS spectrum of the EXPRESS sample compared to
the LHD sample.
Figure 6 shows C-1s XPS spectra of EXPRESS,
LHD and D2+ imp. samples. Three chemical states were
found at the binding energies of 286.1 eV, 284.8 eV and
283.3 eV, attributing to C-O, C-C and C-W bonds,
respectively [13-15]. C-O and C-C bonds were major
chemical states in the LHD sample, although C-C bondwas formed by plasma exposure for the EXPRESS sample.
However, implantation depth of carbon was limited within
the surface region and the total amount of carbon for the
EXPRESS sample was quite small compared to that for the
LHD sample indicating that carbon would hardly affect the
behavior of hydrogen isotope retention in the EXPRESS
sample.
It was concluded that the EXPRESS sample had thethinner impurity deposition layer compared with LHD
sample and hydrogen isotope retention is affected bycarbon and oxygen chemical states and the amounts ofimpurity in deposition layer. In the EXPRESS sample,
most of the impurities were adsorbed on the top surface

Peak 2 was quite different among these samples, It can besaid that the plasma exposure hardly produces theirradiation defects in tungsten due to lower D+ implantationenergy according to SRIM code calculation [8], indicatingthat the formation of irradiation defects was limited for
EXPRESS sample, and most of deuterium would beaccumulated on the sample surface. In the LHD sample,
additional desorption stages were observed at around 750
K (Peak 3) and 900 K (Peak 4). It was reported that theW-O bond in WO3 was one of the stable deuterium
trapping sites with a desorption temperature of around 750
K [9]. In our previous study, the deuterium implanted into
HOPG was trapped as C-D bond after dissociation of C-C
bond, which was desorbed at 900 K [10]. Therefore, thedesorption stages of deuterium at these higher temperaturesin the LHD sample may be assigned to that bound tooxygen as O-H (Peak 3) and carbon as C-H bonds (Peak 4),
respectively.
Figure 4 shows W-4f XPS spectra for the top
surfaces of the EXPRESS, LHD and D2+ imp. samples.
The W- 4f7/2 XPS spectra consisted of two chemical states
located at 30.8 eV for metallic W and 35.8 eV for WO3.
[11,12]. The abundance of WO3 for total W-4f XPS peakarea was 23% for the EXPRESS sample, and 63% for the
LHD sample as shown in Table 1. In the EXPRESS sample,
the major chemical state of tungsten was the metallic state
which was different from the LHD sample, indicating thatmost of deuterium for the EXPRESS sample would beinteracted with metallic W rather than WO3. Figure 5
shows the depth profiles of W, WC, WO3 and C for the
EXPRESS sample. The impurities were accumulated
within the depth of ~ 0.3 nm for the EXPRESS sample.
Table 2 summarizes atomic concentrations on the top

WO3 㻌

1.0

4f5/2

0.8

4f7/2

4f5/2 4f7/2

0.6

LHD sample
EXPRESS sample
D imp. sample

0.4㻌

Intensity / -Metalic WW
mtallic

0.2
0.0

40

35

30

Binding energy / eV

Fig. 4 W-4f XPS spectra for the EXPRESS,
LHD and D2+ imp. samples

Fig. 5 Depth profiles of W and C for the
EXPRESS sample

Table 1 Ratios of W and WO3 on the top surfaceof W

Table 2 Atomic concentrations for the surface of the
EXPRESS and LHD samples and thethickness of the impurity deposition layers.

87R. Miura et al., Study on Plasma-Tungsten Surface Interactions Using the New Experimental Device EXPRESS

1.0

㻌
C-O C-C C-W
LHD sample
EXPRESS sample
D imp. sample

References

P.S. Krstic et al,.Nucl. Inst. and Meth. B 267,704(2009).
K. Ohya et al., Fusion Eng. Des.85,1167(2010).
Y. Oya et al,, J. Nucl. Mater. 390, 622 (2009).
K. Okuno et al.,Fusion Eng. Des. 85, 2328(2010).
H. Eleveld et al., J. Nucl. Mater, 191 433(1992).
H. Iwakiri et al., J. Nucl Mater., 307 ,135(2002).
M. Poon et al., J. Nucl. Mater., 374, 390(2008).
Ziegler J F 2006 SRIM code.
N. Matsunami et al., J. Nucl. Mater. 390 , 693(2009)
T. Suda et al., Fusion Eng. Des. 82 , 1762(2007)
B.F. Dzhurinskii et al., Russian J. Inorganic Chem., 20,2307(1975).
[12] D.R. Chopra et al.,J. Vac. Sci. Technol. B, 10, 1218(1992).
[13] B. Rousseau et al .,J. Phys. Chem., 53, 419(1992).
[14] A. Fahlman et al.,J. Phys. Chem. Solids, 30, 1835(1969).
[15] CH. Cardinaud et al., Applied Surface Science, 68,
595(1993).
[1]
[2]
[3]
[4]
[5]
[6]
[7]
[8]
[9]
[10]
[11]

0.6
㻌

Intensity / -

0.8

0.4
0.2
0.0
288

286

284

282

280

278

Binding energy / eV

Fig. 6 C-1s XPS spectra for the EXPRESS,
LHD and D2+ imp. samples.
after deuterium plasma exposure. However, the amounts ofimpurities beneath the surface were too small to affect thedeuterium retention behavior. In addition, the deuterium
was hardly trapped by irradiation defects because theformation of defects was limited due to low energy ion
irradiation. Therefore, the deuterium was only retained bysurface adsorption during deuterium plasma exposure in
EXPRESS.
In further experiments, longer time (higher fluence)
plasma exposure and oxygen and carbon contaminatedplasma exposure will be needed to make impurity
deposition layers that have a variety of thickness andchemical state in order to clarify the influence of impurity
deposition layer on deuterium retention behavior.

5. Conclusions

A new equipment to acquire laboratory data(well-conditioned data) needed for the design of an actual
reactor named EXPRESS has been designed and isestablishing. The deuterium plasma exposure was
performed to characterize and evaluate the behaviors of
deuterium retention and surface chemical states after the
EXPRESS plasma operation. The deuterium retention
behaviors for tungsten after EXPRESS plasma exposure
was studied and these results were compared to those forthe LHD and D2+ imp. samples. It was found that most of
deuterium was desorbed at the lower temperature less than
550 K for the EXPRESS sample, which was almost thesame with the D2+ imp. sample. However, the TDS
spectrum for the LHD sample was quite different from thatfor the EXPRESS and D2+ imp. samples, which would becaused by the formation of thick impurity deposition layeron the surface for the LHD sample. These facts indicatethat the amounts of impurity in the deposition layer could
affect on the deuterium retention behavior in tungsten usedin actual plasma devices and fusion reactors.

Acknowledgement

This study was supported by NIFS collaboration
program, NIFS11KOBA025 and the Center for
Instrumental Analysis at Shizuoka University.

88

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Hydrogen incorporation into tungsten deposits growing
under hydrogen and argon mixed plasma
Kazunari KATAYAMA, Yasuhito OHNISHI, Satoshi FUKADA and Masabumi NISHIKAWA
Interdisciplinary Graduate School of Engineering Sciences, Kyushu University,
6-10-1, Hakozaki, Higashi-ku, Fukuoka, 812-8581, Japan(Received: 10 May 2012 / Accepted: 9 October 2012)

Tungsten deposits were produced by exposing a tungsten tube to hydrogen and argon mixed plasma in different
deposition conditions. The erosion rate of the tungsten tube decreased with increasing hydrogen concentration in
the working gas. This indicates that the erosion was mainly caused by argon ion incident. In order to releasehydrogen incorporated into the tungsten deposit, the tungsten deposit has to be heated up to 600 oC. While thetungsten deposition rate was promoted by the increase of negative target bias supplied to the tungsten tube, theinfluence of the target bias on hydrogen retention in the tungsten deposit was different in the deposition condition.
However, it was found that hydrogen retention increases with the ratio of the energetic hydrogen flux reflected on
the tungsten tube to the tungsten deposition flux. The flux ratio of hydrogen to tungsten toward the growing surfaceof the tungsten deposit is an important parameter to evaluate the amount of hydrogen incorporated into tungsten
deposits growing under the plasma containing hydrogen ions.
Keywords: tungsten, deposition, hydrogen retention, sputteringhydrogen and argon mixed plasma was investigated.

1. Introduction

It is important to understand the accumulation
phenomena of hydrogen isotopes in a plasma confinement
vessel of a fusion reactor from viewpoints of tritium
economy and safety. Plasma-facing wall is eroded byenergetic particle incident and sputtered atoms would form
deposition layers. It has been known that carbon deposition
layer can contain a large amount of hydrogen isotopes.
Tungsten is a candidate material for plasma-facing wall
because of low solubility and low sputtering yield forhydrogen isotopes. In a long time operation of a fusion
reactor, a certain amount of tungsten deposits would beformed in the plasma confinement vessel. Erosion anddeposition will be promoted if a noble gas such as argon is
introduced around plasma-facing wall for a reduction ofheat load. It has been reported by the present authors that alarge amount of hydrogen isotopes are incorporated into
tungsten deposition layer formed by hydrogen isotope
plasma sputtering [1,2]. The empirical equation has been
proposed from a systematic study of the influence of thedeposition conditions on the deuterium retention by using
magnetron sputtering and linear plasma devices [3]. The
influence of oxygen on hydrogen isotope retention intungsten deposition layers has been investigated [4,5].
Nevertheless, experimental data for hydrogen isotope
retention in tungsten deposits are still insufficient. In thisstudy, the influence of the deposition condition onhydrogen incorporation in tungsten deposits growing under

2. Experimental

A schematic diagram of the experimental apparatus isshown in Fig.1. A cylindrical quartz tube, 13 mm in outer
diameter, 10 mm in inner diameter and 400 mm in length,
which has a horizontal branch tubule, 6 mm in outer
diameter, 4 mm in inner diameter and 30 mm in length,
was used as a plasma sputtering chamber. A tungsten tube,
1.4 mm in outer diameter, 0.2 mm in thickness and 99.95%
in purity (Nilaco Co.), was inserted to the branch tubule asa sputtering target. A copper tube, 3 mm in outer diameter,
was rolled on the outer surface of the quartz tube in the
upstream side as an inductive coil. The distance from theedge of the coil to the tungsten tube was set to be 35 mm.
In order to collect a part of sputtered tungsten, a quartz
substrate, 10 × 5 mm2 in size and 1 mm in thickness, wasplaced beneath the tungsten tube. The detail of the
sputtering-deposition region is illustrated in Fig.2. The
distance from the center of the quartz tube to the surface ofthe substrate was 3.3 mm. The angle T is about 113o whichis 31 % of the all circumferences. The length of the quartz
substrate is 5 mm which is 50 % of the tungsten tube
length in the longitudinal direction. Therefore theplasma-facing area of the quartz substrate views
approximately 16 % of the plasma-facing area of thetungsten tube. The mixture gas of hydrogen and argon was
introduced via a mass flow controller (SEC-E40, HORIBA

author’s e-mail: kadzu@nucl.kyushu-u.ac.jp

89©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

K. Katayama et al., Hydrogen Incorporation into Tungsten Deposits Growing under Hydrogen and Argon Mixed Plasma

supplied in the tungsten tube by a DC power supply(PWR400H, KIKUSUI Co.). The target current wasmeasured by a digital multi-meter (PC520M, SANWA
Inc.). Evacuation of the quartz tube was performed by arotary pump (GLD-201B, ULVAC Inc.). The working gas
pressure during sputtering-deposition process wasmonitored by a Pirani gauge (GP-1G, sencer:WPB-10-034,
ULVAC Inc.). After the process, the absolute pressure wasmeasured by a capacitance manometer (Baratron 626B,
MKS Ltd.) in same deposition condition. The true flowrate of the mixture gas was obtained from the absolutepressure in the quartz tube and the exhaust velocity by therotary pump. The temperature in the vicinity of the quartz
substrate during sputtering-deposition process wasmeasured by a thermocouple fixed on the outer surface ofthe quartz tube. In order to confirm the temperature inside
the quartz tube, a thermocouple was inserted from thedownstream side of the quartz tube to the position of thesubstrate. Both temperatures were almost same. The
deposition conditions were summarized in Table 1. In casesthat temperature is just 100 oC, the temperature wascontrolled by a ribbon heater. The sputtering-deposition
region was wrapped by the ribbon heater except in Ex.1
and Ex.2. The temperatures in Ex.1 and Ex.2 are lower
than that in the others because there is no effect of heatretention by the ribbon heater.
The amount of deposited tungsten was obtained from theweight change of the substrate before and after
sputtering-deposition process. The weight change wasmeasured by an electric balance with sensitivity of 0.01 mg(CP225D, Sartorius Inc.), whereas the original weight ofthe substrate was about 100 mg. The cross section
observation of the tungsten deposit was performed by thescanning electron microscopy (SEM:SU8000, HITACHICo.). The weight percent of oxygen in the tungsten deposit
was analyzed by energy dispersive X-ray equipment(EDX:Genesis2000, EDAX Inc.). SEM and EDX used inthis study were installed at the Center of Advanced

Fig.1 The schematic diagram of the experimental apparatus.

Fig.2 The illustration of the sputtering-deposition region.
(a) : a view from the gas flow direction. (b) : a view from thecross-section direction of the W tube.

Ltd.) calibrated with pure hydrogen gas. Hydrogenconcentrations in the mixture gas used in this study were 2,
20, 50 and 80 %. A packed bed of adsorbent, MS5A wasinstalled between the gas cylinder and the mass flow
controller to exclude a small amount of impurity water
vapor in the gas cylinder. Plasma was inductively
generated inside of the quartz tube by supplying the RF
power of 40 W or 100 W with 13.56 MHz (NR05N-06,
NRF Co.) to the copper coil via matching box(NM05AL-05, NRF Co.). In order to promote erosion of
the tungsten tube, a negative bias of 200 V or 300V wasTable 1 The summary of deposition conditions.

H2[%]

Flow rate[sccm]

Pressure[Pa]

RF power[W]

Targetbias [-V]

Flux to target[ions/m2/s]

Temperature[oC]

Deposition
period [h]

Ex.1

2

3.8

7.6

40

200

1.30㬍1020

58

20

Ex.2

2

3.8

7.6

40

300

1.53㬍1020

58

17

Ex.3

20

6.7

7.1

100

200

20

1.31㬍10

104

20

300

1.96㬍1020

94

21

100

14

Ex.4

20

6.7

7.1

100

Ex.5

20

3.1

4.6

100

300

1.62㬍1020

Ex.6

20

6.7

7.1

100

300

2.72㬍1020

96

36.5

Ex.7

20

3.1

4.6

40

200

1.23㬍1020

100

24

Ex.8

50

3.8

5.4

100

200

20

2.61㬍10

146

25

106

46

156

30

Ex.9

50

3.8

5.4

100

300

2.77㬍1020

Ex.10

80

2.5

4.2

100

200

3.99㬍1020

90

K. Katayama et al., Hydrogen Incorporation into Tungsten Deposits Growing under Hydrogen and Argon Mixed Plasma

0.35

-300V

0.30

6

1500

1000

-200V

4

2

500

Deposition rate [mg/h]

+

8

+

Erosion yield [mg/h/A]

2000

0.40

10

RF 40W RF 100W

Ion density ratio, H2 /Ar [-]

2500

0.25

H2/Ar ratio
2%
20%
50%
80%

0.20
0.15

Deposition rate
= 0.12
Erosion rate

0.10
0.05

0

0

20

40
60
H2/(Ar+H2) [%]

80

0

0.00
0.0

0.5

1.0

1.5 2.0 2.5 3.0
Erosion rata [mg/h]

3.5

4.0

Fig.3 The dependence of hydrogen concentration in theFig.4 The deposition rate on the substrate with the erosion

working gas on the erosion yield of the tungsten tube and on

rate of the tungsten tube.

the ratio of hydrogen ions to argon ions.

increasing hydrogen concentration. Additionally, theweight loss of the tungsten tube was undetectable inseveral-hour exposure under pure hydrogen plasma. These
results indicate that the erosion of tungsten was caused byargon ion incident. Impinging energy of argon ionsobtained from negative bias is mitigated by collisions withneutral molecules. Therefore the erosion yield isconsidered to decrease with decreasing the argon ion
density. The erosion yield at a negative bias of 300 V wasquite lager than that at a negative bias of 200 V in 2 % and
20 % of hydrogen concentration but these are close in
50 % of hydrogen concentration. This indicates that thenegative bias effect on erosion is large when argon ion
density is high.
The deposition rate against the erosion rate in eachdeposition condition is plotted to Fig.4. The ratio of thedeposition rate to the erosion rate was 0.12 on the average.
This value is smaller than 0.16 which is estimatedgeometrically. This difference is due to the erosion of thetungsten tube in the branch tubule, the slip of the substratefrom the designed position and the scattering of sputteredtungsten atom by collision with neutral molecules.
SEM image of the cross-section of the tungsten deposit
formed under hydrogen and argon mixed plasma, 20 % inhydrogen concentration, 7.1 Pa in working gas pressure,
100 W in RF power, 300 V in negative bias, 96 oC in
temperature and 36.5 h in deposition period (Ex.6) isshown in Fig.5. The quartz substrate was broken in half
and the surface was observed. It was found that thetungsten deposit has a columnar structure. The weightpercent of oxygen in the tungsten deposit formed underhydrogen and argon mixed plasma, 20 % in hydrogenconcentration, 4.8 Pa in working gas pressure, 40 W in RF
power, 200 V in negative bias and 100 oC in temperature(Ex.7), is 2.0 wt%.
An example of hydrogen release behavior from tungsten

Instrumental Analysis, Kyushu University.
Ion densities at the vicinity of the substrate under purehydrogen plasma and pure argon plasma were measured by
the double probe method. The probes were inserted fromthe branch tubule to plasma without the target tube.
Additionally, the emission intensities from H2+ and Ar+ in
plasma at the vicinity of the substrate were measured forpure hydrogen, pure argon and hydrogen-argon mixed
plasma by a plasma process monitor (PlasCalc-2000,
Mikropack). Ion densities of hydrogen and argon in themixed plasma were estimated from the double probe
measurement and the emission intensity measurement
assuming the emission intensity is proportional to the ion
density and H2+ is dominant in hydrogen ions.
After sputtering-deposition process, the substrates wereplaced into a separate quartz tube connected to a gas
chromatograph (GC:GC-8A, SHIMAZU Co.). The quartz
tube was filled with argon gas and heated from 200 oC to
800 oC in 200 oC step by an electric furnace. Argon gas
was introduced into the reaction tube every 20 minutes inorder to transport the released hydrogen into GC. Since theamount of hydrogen released at 800 oC was comparable tothe background level, the hydrogen released by 600 oC was
regarded as the hydrogen incorporated in the tungsten
deposit.

3. Results and discussions

Fig.3 shows the dependence of hydrogen concentrationin the working gas on erosion yield of the tungsten tube
which was derived from dividing the erosion rate (mg/h)by the target current (A). The right axis indicates the ratioof hydrogen ion density and argon ion density under eachdeposition condition. The hydrogen ion ratio increased
approximately proportionally to hydrogen concentration. It
was observed that the erosion yield decreases with91

K. Katayama et al., Hydrogen Incorporation into Tungsten Deposits Growing under Hydrogen and Argon Mixed Plasma
㩷

㩷

1.5

3.0

3Pm

800

-200V (Total:1.22×10 H2-mol/g)

700

-5

-300V (Total:4.99×10 H2-mol/g)

8
7

600
Temperature

500

5

400

4

300

3

200

2

0

20

40

60

-4

0.0

6.9 u10 3 f H f W 0.42 ,(1)

where 7.3 ≤ fH/fW ≤ 4780, 65 ≤ En ≤ 652 eV, 331 ≤Temperature ≤ 429 K. En is the average energy of thereflected particles on the target. The value of En was
estimated using the following formula [6]:

100

1
0

H D
W

o

6

300

between the target and the substrate. Therefore, the bias of
the target is considered to be one parameter that affects onhydrogen retention in tungsten deposits forming by
sputtering method. However, the result shown in Fig.6
suggests that other parameter also contributes to hydrogen
retention. The present authors considered that the ratio ofhydrogen flux to tungsten flux from the target to thegrowing surface, fH/fW contributes to hydrogen retention.
The hydrogen flux reflected from the tungsten target wasestimated from the hydrogen ion flux and the fraction ofbackscattered particles, RN [6]. The hydrogen ion flux was
estimated from the target current and the ratio of hydrogenion density to argon ion density. The tungsten flux was
estimated from the deposition rate. The relationshipbetween the ratio of hydrogen flux to tungsten flux andhydrogen retention as H/W is shown in Fig.8. The amountof tungsten deposited on the substrate was obtained fromthe deposition weight. Data from the parallel plate RF
sputtering device [2,7] are also indicated. Although thetungsten deposit reported in [2] contained not only
deuterium but also hydrogen, the reflected particle isassumed to be only deuterium because the working gas is
pure deuterium. The expression obtained from fitting all of
data is as follows:

Temperature [ C]-5

H2 release rate [10 mol/g/(20min)]9300 200
Target voltage [-V]

200

0.5

and deposition rate for three deposition conditions.

deposits is shown in Fig.6. For all samples, the release
behavior of hydrogen was similar. The hydrogen releasebegan at 200 oC and continued until 600 oC. The highest
release rate was observed at 400 oC. It was found that ahigh temperature of 600 oC is necessary to removehydrogen incorporated in the tungsten deposit.
The dependence of negative bias on hydrogen retention(mol-H2/g) and deposition rate (g/hour) are shown in Fig.7
for three deposition conditions. It was observed that thedeposition rate at 300 V is obviously higher than that at
200 V. This means that the impinging energy of ions to thetungsten tube at 300 V is higher than that at 200 V. On theother hand, the change of hydrogen retention with negative
bias was different in the deposition condition. A part ofhydrogen and argon impinging to the target is considered
to be reflected and implanted into the growing surface ofthe tungsten deposit. These reflected particles are
considered to have relatively-high energy as comparedwith ions which impinge to the growing surface of thetungsten deposit with the sheath potential although its
energy is reduced by the collision with neutral molecules-4

2%H2/Ar, RF 40W, 7.6Pa
20%H2/Ar, RF100W, 7.1Pa
50%H2/Ar, RF100W, 5.4Pa

Fig.7 The dependence of target bias on hydrogen retention

formed in Ex.6.

2%H2, RF40W, 7.6Pa

1.0

0.5

0.0

Fig.5 SEM image of the cross-section of the tungsten deposit

10

1.5

Deposition rate [×10 g/h]

2.0

1.0

-4

H2 retention [×10 mol-H2/g]

2.5

E n ( E0 )080 100 120 140 160 180
Time [min]E0

RE ( E0 )R N ( E0 ),(2)where E0 is the energy of the ions impinging on thetungsten target, RE is the energy reflection coefficient. The

Fig.6 Hydrogen release behavior from the tungsten deposit
formed in Ex.1 and Ex.2.

92

K. Katayama et al., Hydrogen Incorporation into Tungsten Deposits Growing under Hydrogen and Argon Mixed Plasma

4. Conclusions-3

The influence of deposition condition on hydrogen
incorporation into tungsten deposits growing underhydrogen and argon mixed plasma was investigated byexposing a tungsten tube to the plasma. The erosion of thetungsten tube is mainly caused by the incident of argon
ions. The tungsten deposit has a columnar structure. In
order to remove the hydrogen incorporated into thetungsten deposit, a high temperature of 600 oC is necessary.
While the tungsten deposition rate is promoted by theincrease of negative target bias from 200 V to 300 V, theinfluence of the target bias on hydrogen retention isdifferent in the deposition condition. Hydrogenconcentration in the tungsten deposit increases with theratio of the flux of the hydrogen reflecting from thetungsten target to tungsten deposition flux. The flux ratio ofhydrogen and tungsten toward the growing surface of thetungsten deposit is an important parameter to evaluatehydrogen accumulation amounts.

0.42(H + D)/W = 6.9×10 (fH/fw)(fitting all of data)(H+D) / W [-]10652eV/D [2]-1580eV/D [2]480eV/H [7]

10

10

-2

95eV 65eV

-3

1

10

: 7.6Pa, 2%H2
: 7.1Pa, 20%H2: 4.6Pa, 20%H2: 5.4Pa, 50%H2: 4.2Pa, 80%H2o: 10Pa, 100%H2, 68 Co
: 2Pa, 100%D2, 71 Co
: 10Pa, 100%H2, 111 C

100
1000
H flux / W flux [-]10000

Fig.8 The dependence of the ratio of hydrogen flux to tungsten
flux on hydrogen isotope retention.

Acknowledgments

This work was supported by a Grant-in Aid for Young
Scientists (B) from the Japan Society for the Promotionof Science (KAKENHI21760691).

obtained value, 0.42 is close to the value of 0.4±0.1
reported by De Temmerman and Doerner [3], where 2 ≤(fH/fW) ≤ 500, 60 ≤ En ≤ 280eV, 293 ≤ Temperature ≤ 600 K.
The dependence of En on the retention was not observed inthis work while they suggest in [3] that the retention is
proportional to En1.88. Strictly speaking, the impinging
energy of reflected hydrogen to the growing surface of thetungsten deposit is abated by collisions with neutral
molecules between the target to the substrate. That is to say,hydrogen retention depends on the working gas pressure as
has been reported in [2]. Especially, the influence of thepressure would be large in the parallel plate RF sputtering
device [7] because the distance from the target to thesubstrate is long. The impinging energy of hydrogenisotope to the growing surface in the parallel plate RF
sputtering device might be comparable with that in thedevice used in this work. Hydrogen retention depends on
temperature as has been reported in [3,4,8]. For the datasummarized in Fig.8, these effects would be able to ignore
because the temperatures are relatively low and the range
is narrow.

5. References
[1]
[2]
[3]
[4]
[5]
[6]

[7]
[8]

93

K. Katayama, K. Imaoka et al., Fusion Sci. Technol. 54,
549 (2008).
K. Katayama, S. Kasahara et al., Fusion Eng. Des. 86,
1702 (2011).
G. De Temmerman and R.P. Doerner, J. Nucl. Mater. 389,479 (2009).
V.Kh. Alimov, J. Roth et al., J. Nucl. Mater. 399, 225(2010).
T. Fujiki, K. Katayama et al., Fusion Eng. Des. 85, 1094
(2010).
W. Eckstein, Calculated Sputtering, Reflection and Range
Values, Report of theMax-Planck-Institute fr Plasmaphysik,
IPP-Report 9/132, Garching, Germany (2002).
S. Ishikawa, K. Katayama et al., Fusion Eng. Des. 87, 1390
(2012).
K. Katayama, Y. Uchida et al., J. Nucl. Mater. 390-391,
689 (2009).

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Hydrogen Storage Properties of Nanocrystalline Mg2Ni
Based Alloys Prepared by Ball-Milling
Yifu XIONG, Jingwen BA, Wuwen QINGˈWenyong JING
China Academy of Engineering Physics, Mianyang, 621900, China(Received: 15 May 2012 / Accepted: 26 November 2012)

Nanocrystalline hydrides are a new class of material in which outstanding hydrogen sorption may be obtainedby proper engineering of the microstructure and surface. The nanocrystalline Mg2Ni alloy is a promising hydrogenstorage material. In the present work, nanocrystalline Mg2Ni alloy powders with grain size of about 50 nm wereprepared by high-energy ball-milling, and its phase, crystal structure and hydrogen storage properties were
investigated by X-ray diffraction analysis, transmission electron microscopy and pressure-composition isotherms,
respectively. The hydrogen storage characteristics of Mg2Ni are also presented. Nanocrystalline Mg2Ni can readily
absorb hydrogen at temperature lower than 523 K. The reversible hydrogen capacity is up to 3.5 wt.%.
Keywords: Mg2Ni, nanocrystalline, mechanical alloying, ball-milling, hydrogen storage

elemental Mg and Ni powders under hydrogen
atmosphere. Mg2Ni alloy powder has been followed by
X-ray diffraction (XRD) and transmission electron
microscopy (TEM), and the values of enthalpy andentropy changes for hydrogen absorption and desorption
were evaluated from the pressure-composition isotherms
on van’t Hoff equation.

1. Introduction
Many metallic materials are known to form hydrides
reversibly. Intermetallic Mg2Ni with its high hydrogencapacity (up to 3.6 wt.%) is the prime candidates amonghydrogen storage systems [1–5]. However, Mg2Ni cannot
absorb hydrogen under normal conditions (i.e. room
temperature and atmospheric pressure), and the ability ofhydrogenation appears in temperature range from
523–623 K with hydrogen pressure of 1.5–5 MPa.
Moreover, even if Mg2Ni is heated more than 523 K, it
also needs to be activated before hydrogenation. The
activation process usually involves a heating process at ahigh temperature (598 K) under a high hydrogen pressure(2 MPa). The activation process has to be carried out
several times to obtain reproducible hydrogen
absorption/desorption characteristics. Improvement of thehydrogenation conditions of Mg2Ni is therefore essential
to produce a suitable material for practical hydrogenstorage [6–8].
The mechanical alloying (MA) may give various
homogeneous composite particles. Since 1970, the MA
process via the ball-milling (BM) technique has beenused to prepare several dispersion-strengthened alloy
powders. The application of the MA process has been
expanded recently to produce several hydrogen storage
metals using a method called the reactive ball-milling.
Metal hydrides are usually prepared by passing a flowof reactive hydrogen over the metallic material underhigh pressure and at temperatures usually above room
temperature. In our study, the single-phase Mg2NiH4
alloy powders with grain size about 50nm has beensynthesized by milling an equiatomic mixture of

2. Experimental
In order to produce Mg2Ni, a mixture of magnesium(99.95%, 150 mesh) and nickel (99.95%, 250 mesh),purified hydrogen and argon gas (H2O and O2, less than
20 ppm) has been used. The mixed powders were charged
and sealed in a cylindrical stainless steel shell together
with stainless steel rods in a laboratory high-energy mill
SPEX 8000 (from SPEX Industries). The rod-to-powderweight ratio was controlled to be about 20 to 1.The inletof the shell was connected to a rotary pump andevacuated for about 4 Pa. After evacuation, a flow ofhydrogen gas was passed into the rod mill through abellows pipe. Once the rod mill was filled withhydrogen at 0.25 MPa, the inlet of the vial was closed
and the reactive rod-milling (RRM) process was
carried out at ambient temperature by mounting therod mill on a rotator at the rate of 2.0 s-1. The RRM
was stopped at selected intervals and a small amountof the rod-milled powder was taken out under argon
atmosphere in a glove box.

3. Results and discussion
Mg2Ni is difficult to prepare by conventional
metallurgy because of the great difference in melting

author’s e-mail: xiongyifu@yeah.net

94©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

Y. Xiong et al., Hydrogen Storage Properties of Nanocrystalline Mg2Ni Based Alloys Prepared by Ball-Milling

temperature and vapor pressures of Mg and Ni.
Mechanical alloying by ball-milling requires no melt
process of the elements and yields the compound througha solid-state reaction. Although this ball-milling is quite acommon technique, published results on ball-milled Mg
and Ni powders have not yet been shown to produce
single-phase [2,5].
Under the ball-milling conditions (high-energy milland high ball-to-powder mass ratio) in the present work,
Mg2Ni alloy powder was formed directly. The
ball-milling parameters were as follows: The ratio of rod
and metallic powder was 20 to 1, rotational speed ofball-milling per minute was 0.2 s-1 and processing
control reagent is stearic acid.
Fig. 1 shows the X-ray diffraction patterns of Mg2Ni
after mechanically ball-milling for 2 h and 50 h underhydrogen atmosphere. Fig. 1 (a) shows an intermediatestage of mechanical alloying reaction (after 2 h ofball-milling), and this diffraction pattern includes thepeaks of the Mg2Ni compound, Mg and Ni. After millingfor 50 h, the diffraction peaks of Mg2NiH4 become
broader and none of the peaks corresponding to MgH2
and NiH2 are seen. It is worth noting that the broadness ofthese peaks can be attributed to the formation of
Mg2NiH4 solid solution powder, as shown in Fig. 1 (b).
We should emphasize that this solid solution phase has
not changed to any other phase, even after milling for as
long as 50 hours.
The size of nanocrystalline Mg2Ni alloy was calculatedby Scherrer formula as follows:
D = KO/E1/2cosT, (1)where D is the size of nanocrystalline, K is Scherrer
constant (0.89), O is the wavelength of CuKD X-rays(0.15405 nm), E1/2 is the width of half-height for
diffraction peak (rad). Those of (101) and (105) ofcommon Mg2Ni alloy are 0.3062° and 0.2879°,
respectively. In our work, those of (101) and (105) of
nanocrystalline Mg2Ni alloy were 0.3204° and 0.2950°,
respectively. The size of nanocrystalline Mg2Ni alloy wasevaluated from XRD patterns to be about 50 nm.
Detailed TEM analyses were performed in order tounderstand the change in the fine structure ofmechanically alloyed Mg2NiH4 particles during
above-mentioned stages of milling (Fig. 2). The powder
have large grains of about 50 nm diameter and theselected-area diffraction patterns (SADPs) taken at thecenter of this micrograph shows sharp rings coexisting1
with spot patterns of nanocrystalline Mg2Ni alloy powder
in Fig. 2 (c).
Figs. 3 and 4 show pressure-composition isotherms ofnanocrystalline Mg2Ni prepared with 50 h of ball-milling.
The absorption/desorption temperature was in a range
from 473 K to 623 K, and the nanocrystalline Mg2Ni-H2
system has obvious and wide plateau region. Reversiblehydrogen capacity amounts to 3.5 wt.% within the 95

Fig. 1 XRD patterns of Mg-Ni alloys prepared byball-milling for 2 h and 50 h.

1

1

1

Fig.2 TEM images of Mg-Ni alloys prepared byball-milling under different conditions(a-2h, 120 K˗b-50h, 150 K˗c-electron
diffraction pattern n (50h)).
applied hydrogenation time, and the capacity agreed withthe literature [5]. The amount of absorbed hydrogen isusually 3.0–3.45 wt.% in the reversible hydrogenation
cycles as reported in [5, 6]. The hydrogenation of thehigh-temperature phase of Mg2NiH4 is apparently mucheasier than for the low-temperature phase. The
low-temperature phase of Mg2NiH4 is very difficult toobtain by the direct hydrogenation under moderatepressure conditions. The plateau pressures difference
between absorption and desorption at the same
temperature was observed (Fig. 5). The phenomenon wasreferred to hysteresis (its represents by hysteresis factor
Hf), Hf increases with temperature. The hysteresis
originates duo to the presence of transformation strains
during both hydride formation and decomposition.

Y. Xiong et al., Hydrogen Storage Properties of Nanocrystalline Mg2Ni Based Alloys Prepared by Ball-Milling

The van’t Hoff plots of the nanocrystalline Mg2Nialloy (log(PH2) versus 1/T) were shown in Fig. 6. The
best–fitting equations for 'H 4 and 'S 4 are given byhydrogen absorption process:log(P / 0.1 MPa) = í3843/T + 7.012, (2)and by hydrogen desorption process:log(P / 0.1 MPa) = í3992/T + 7.173. (3)The values of enthalpy and entropy changes for hydrogen
absorption are í73.58 kJ·mol-1 and í134.26 J·mol-1·K-1,
and those for hydrogen desorption are 76.44 kJ·mol-1 and137.34 J·mol-1·K-1, respectively. The results agreed wellwith the literature [8].

Fig. 3 Pressure-composition isotherms of hydrogen
absorption for nanocrystalline Mg2Ni alloy.

Fig. 6 Temperature dependence of plateau fornanocrystalline Mg2Ni alloy.

Fig. 4 Pressure-composition isotherms of hydrogen
desorption for nanocrystalline Mg2Ni alloy.

hydrogen content(x)Hydrogen absorption properties for ball-milled
nanocrystalline Mg2Ni alloy were better than those forconventional Mg2Ni alloy. The as-produced powder
readily absorbs hydrogen with no activation, and the final
hydrogen content was found to be about 3.5 wt.% after
the first hydrogenation cycle at 573 K (Fig. 7). In
successive cycles, the absorption kinetics of Mg2Ni
prepared was reproducible and about four times fasterthan that of conventional Mg2Ni. The results imply thatthe nanocrystalline Mg2Ni alloys would be a promisinghydrogen storage material.

Fig. 5 The relationship between plateau pressure andtemperature for nanocrystalline Mg2Ni alloy.
According to van’t Hoff equation, the relationshipbetween  temperature and hydrogen equilibriumpressure can be expressed as 'H 4' S 4 , (1)log( P H 2 )RTR
where 'H 4 is the standard enthalpy changes (J·mol-1)
of hydriding reaction, ' S 4 is the standard entropychanges (J·mol-1·K-1) of the reaction, R is the gas
constant, T is the given measurement temperature, andPH2 is the equilibrium hydrogen pressure.

4.5
4.0
3.5
3.0
2.5
2.0
1.5
1.0
0.5
0.0
-0.5

A
B

0

200 400 600 800 1000t(min)

Fig.7 Absorption rate of hydrogen at 573K
A-nanocrystalline Mg2Ni alloy (50h)and B- conventional Mg2Ni alloy96

Y. Xiong et al., Hydrogen Storage Properties of Nanocrystalline Mg2Ni Based Alloys Prepared by Ball-Milling

4. Conclusions
The Mg2Ni alloys were prepared by ball-milling,which was composed of nanoparticles with grain size of
about 50 nm. The structures and the hydriding behaviorof Mg2Ni prepared were studied. The absorption anddesorption thermodynamic properties of Mg2Ni alloys
were improved by a ball-milling method. The reversiblehydrogen capacity was found to be to 3.5 wt.%. The
values of the enthalpy and the entropy for hydrogen
absorption are í73.58 kJ·mol-1 and í134.26 J·mol-1·K-1,
and those for hydrogen desorption are 76.44 kJ·mol-1 and137.34 J·mol-1·K-1, respectively. The nanocrystalline
Mg2Ni alloy will be a promising hydrogen storage
material for hydrogen energy and fuel cell systems.
References:[1] L. Zaluski et al., J. Alloys Compd., 217, 245 (1995).
[2] S. Orimo et al., J. Alloys Compd., 232, L16 (1996).
[3] L. E. A. Berlouis et al., J. Alloys Compd., 305, 82 (2000).
[4] L. Zaluski et al., J. Alloys Compd., 70, 253 (1997).
[5] S. Orimo et al., J. Alloys Compd., 217, 287 (1995).
[6] S. Orimo et al., J. Alloys Compd., 253, 94 (1997).
[7] H. Imamura et al., J. Alloys Compd., 386, 211 (2005).
[8] P. Selvam et. al, Int. J. Hydrogen Energy, 11,169 (1986).

97

J. Plasma Fusion Res. SERIES, Vol. 10 (2013)

Hydrogen generation property on the surface of plasma sintered Be12Ti
Kohei WADA1), Kenzo MUNAKATA1), Yusuke AKIMOTO1), Haruki TAKEDA1),
Jae-Hwan KIM2), Daisuke WAKAI2),Kazuo YONEHARA2) and Masaru Nakamichi2)
1)

Faculty of engineering and resource science, Akita university, 1-1, Tegatagakuen-cho, Akita, 010-8502, Japan

2)

Fusion Research and Development Directorate, Japan Atomic Energy Agency, 2-166, Omotedate, Obuchi, Rokkasho,
Kamikita, Aomori, 039-3212, Japan(Received: 9 May 2012 / Accepted: 5 December 2012)

Central and peripheral parts of a plasma sintered titanium beryllide disc were exposed to water vapor attemperature rising up to 1273 K. Hydrogen generation on and oxidation properties of the titanium beryllide surfaces
were investigated. According to electron probe microanalysis, the titanium beryllide parts were characterized bydifferent surface area fractions of Be, Be12Ti, Be17Ti2, and Be2Ti phases, and the central part contained lower
fraction of Be phase than the peripheral part. H2 gas generation rate in the central part was found to be lower thanthat in the peripheral part, and this difference can be attributed to larger fraction of the Be phase. Secondary
electron microscopy and X-ray diffraction revealed that the parts of the plasma sintered Be12Ti disc had higher
oxidation resistance than metallic beryllium.
Keywords: beryllide, plasma sintering, hydrogen generation, oxidation, gas chromatograhphy

1. Introduction

caused by interaction of water vapor with titanium
beryllides. Besides, oxidation resistance of the titanium
beryllides with slightly different phase compositions
was studied.

Metallic beryllium has been considered as a leading
candidate for a neutron multiplier in water- or
helium-cooled solid breeder blankets in future fusion
reactors. In these blankets, beryllium will be subjected
to a high neutron flux and located in high temperatureenvironment. However, beryllium is highly reactivewith water vapor at temperatures above 873 K,
producing hydrogen that may lead to accidents due toH2 gas explosion.
Be + H2O = BeO + H2
(1)
Titanium beryllides such as Be12Ti have many
advantages over beryllium from the perspectives ofhigher melting point, lower chemical reactivity, lower
swelling, etc. Therefore, Be12Ti material has attracted
attention as an alternative of beryllium, and could be
used as an advanced neutron multiplier in the fusion
reactor blankets [1].
Titanium beryllides have been fabricated by various
methods. Nakamichi et al. [2, 3] have developed new
synthesis process of Be12Ti material called“plasma-sintering method”. This process consists ofloading raw material powder in the punch and the dieunit, direct current pulse plasma generation to activate
the surface of powder particles, and uniaxial pressing toenhance the sinterability [2,3].
In this study, we report results on H2 generation

2. Experimental
Be and Ti powders, characterized by average grain sizesof 27 ȝm and 28 ȝm, respectively, were mixed for 60 min
by using a mortar grinder RM200 (Retsch, Germany). The
mixed ratio was adjusted to the stoichiometric compositionof Be12Ti (69.3:30.7 wt.%). The mixed powder was placedin a plasma electrical discharge sintering device KE-PAS III(manufactured by KAKEN Co. Ltd), and Be12Ti disc, 20

MS3A

CuO
bed

H2O
Ar

H2O

MS3A

Ar

M.F.

H2
G.C.

M.F.

Electric
Electricfurnace
furnaceAr

H2/Ar

M.F.͏͏͏͏Mass
M.F.͏͏͏͏Massflow
flowcontroller
controller
MS3A䈇䈇䈇Molecular
MS3A䈇䈇䈇Molecularsieve
sieve3A
3A
CuO
CuObed䈇䈇Oxidized
bed䈇䈇Oxidizedcopper
copper
G.C.䈇䈇䈇䈇Gas
G.C.䈇䈇䈇䈇GaschromatographchromatographFig. 1 Schematic flow diagram of test apparatus.

authors e-mail : kohei.wada1986@gmail.com

98©2013 by The Japan Society of Plasma
Science and Nuclear Fusion Research

K. Wada et al., Hydrogen Generation Property on the Surface of Plasma Sintered Be12Ti

Be12Ti, Be17Ti2, and Be2Ti with different fractions.
Percentages of the surface areas of the phases for each partof the disc, evaluated with the use of an image analyzer, areshown in Table 1. It is evident that the central part iscomposed mainly of Be12Ti phase with relatively large
fraction of Be17Ti2 phase. Additionally, smaller fractions ofBe and Be12Ti phases were observed. Note that phasecompositions for the central and peripheral parts wereslightly different. This observation is in good agreementwith results obtained by Kim et al. [4] proposing that thephase difference for the central and peripheral parts can be
explained by a temperature gradient maintained during thesintering process.
Experiments on generation of H2 gas on the surfaces ofthe plasma sintered titanium beryllides were carried out toinvestigate the effect of the phase composition. Fig. 3 showschanges in hydrogen concentrations in the outlet stream ofreactor with central and peripheral part of the plasma
sintered titanium beryllide disc exposed to a 10,000 ppm
H2O/Ar gas. In case of the central part, generation ofhydrogen started at a temperature of 673 K and continued
for about 27 h. The total amount of hydrogen generated for
24 h is 1.210-4 mol/cm2. For the peripheral part, however,
the hydrogen generation started at a temperature above 423
K and continued for 40 hr. The total amount of hydrogengenerated on the peripheral part for 24 h is 3.010-4
mol/cm2. As mentioned above, the samples used in thiswork included the phase of Be2Ti (Table 1) that is known to

mm in diameter and 5 mm in thickness, was synthesized bysintering for 40 min at a temperature of 1273 K and apressure of 50 MPa. Then the disc was cut into two small
samples to separate central and peripheral parts of the discfor comparative analyses. The surface area of the central
parts was 83 mm2 while that of the peripheral part was 128
mm2. The samples were mechanically polished with a#2400 SiC abrasive paper.
The surface morphology of the Be12Ti samples was
examined by secondary electron microscopy (SEM)
combined with the electron probe microanalysis (EPMA) ina SEM/EPMA instrument JXA-8230 (JEOL Co.)
Schematic flow diagram of apparatus for ameasurement of the hydrogen generation rate is shown in
Fig. 1. The Be12Ti sample was wrapped by platinum meshto avoid direct contact with the test tube, and then placed
between silica wools. 10,000 ppm H2O/Ar gas was
generated by passing 10,000 ppm H2 /Ar gas throughcopper oxide bed heated at 623 K, and then introduced tothe test tube with a flow rate of 300 cm3/min. An electric
furnace was used to heat the test tube at a heating rate of 5
K/min, and a temperature inside the tube was raised up to1273 K. This temperature was kept constant until hydrogengeneration terminated.
The concentrations of hydrogen in the outlet streamsof the reactor were measured with a gas chromatograph
GC-8A manufactured by SHIMADZU Co. The state ofoxidized sample surface was characterized by X-ray
diffraction (XRD) analysis with the Ultima IV instrument(RIGAKU Co.) and electron probe micro analysis (EPMA)
with the SEM/EPMA JXA-8230 instrument (JEPL Co.).

Table 1 Surface area fractions (in %) of phases constituent
central and peripheral parts of titanium beryllide disc.

3. Result and discussion
Figure 2 shows the SEM images using back-scatteredelectron for central and peripheral parts of the plasma
sintered titanium beryllide disc prepared in this study. The
results of EPMA analysis demonstrated that the titanium
beryllide is composed of four different phases such as Be,

Fig. 2 SEM images of central part (a) and peripheral part (b) of plasma sintered titanium beryllide disc.

99

K. Wada et al., Hydrogen Generation Property on the Surface of Plasma Sintered Be12Ti

have capability of hydrogen storage [5]. The existence of
this phase might affect the hydrogen generation of thesamples, while the fractions of the phase in the both
samples were almost the same. The Be17Ti2 phase could
also affect the hydrogen generation, but this phase isprobably stable and tolerant to oxidation as much as the
Be12Ti phase [Kim paper Numat]. Thus, it is thought thatthe sample with larger fraction of Be phase generates moreamounts of H2, since Be is known as highly reactive withoxygen [6] and water vapor [7]. As shown in the EPMA
analysis (Table 1), the edge sample was composed of larger
fraction of Be, which affect the amount of generated H2 gas.
This agrees with the report by Anderl et al. [8],demonstrating that the rate of the reaction between
beryllium surface and water vapor is associated withdifferences in the specific surface areas for the tested
materials.
The weight gain for the Be12Ti and beryllium samples
indicate that the center sample part increased by 3.6 % of
the initial weight for the central part of the titanium
beryllide and 4.7 % for the peripheral part. However, as
reported in our previous study [9], for metallic beryllium
used as a reference material, the weight gain was 50 % of
the initial weight. Data on the weight gain and amount ofgenerated hydrogen for the titanium beryllide parts indicatethat the peripheral part is more oxidized.
SEM observation and XRD analysis were performed toinvestigate the surface state of the two samples after
exposure to 10,000 ppm H2O/Ar gas. Fig. 4 shows the SEM
images of central (a) and peripheral (b) parts of the plasma
sintered titanium beryllide disc and metallic beryllium (c)after exposure to water vapor. The Be sample (shown in Fig.
4 (c)) was used in our previous work and the surface of thesample was observed using the SEM/EPMA apparatus [9].
The result of the SEM analysis clarified that after oxidation
the surfaces of the titanium beryllide parts are covered with
dense and protective oxide scales (see Fig. 4 (a) and (b)).
Gray and white areas in the SEM pictures correspond to
1400

1000
800

Hydrogenconcentration

50

600
400

25

(b)

(a)
0

1200

Temperature

75

0

5

10

15

20

Temperature (K)Hydrogen concentration (ppm)

100

BeO and TiO2, respectively. Comparison of the oxidation
behavior of Be12Ti and beryllium metal indicate that more
protective BeO scales are present on the surface of the
Be12Ti sample parts. In contrast, it could be stated thatmetallic Be possesses poor oxidation resistance due tocracked and fractured BeO scales present of the surface.
Sato et al. [10] had demonstrated that BeO formed on thesurface of metallic Be is affected by compressive stresscaused by coherency of the lattice structures between thescale and the bulk metal, since the atomic distances in the
BeO scale are smaller than those in the metallic Be
substrate.
The results of the XRD analysis (Fig. 5) demonstrated
that the intensity of peaks corresponding to BeO formed on
the peripheral part are higher than those for the central part.
This is in good agreement with the result of the weight gain

200
25

30

35

0
40

Time (h)Fig. 3 Change in concentration of hydrogen in the outlet streamof reactor with central part (a) and peripheral part (b) of theFig. 4 SEM images of central (a) and peripheral (b) parts of theplasma sintered titanium beryllide disc exposed to a 10 000 ppm

plasma sintered titanium beryllide disc and metallic beryllium

H2O/Ar gas.

[8] (c) after exposure to water vapor.

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K. Wada et al., Hydrogen Generation Property on the Surface of Plasma Sintered Be12Ti

Border Approach activities.

6. References[1] H. Kawamura, H. Takahashi, N. Yoshida, V. Shestakov, Y. Ito,
M. Uchida, H. Yamada, M. Nakamichi and E. Ishitsuka,
Fusion Eng. Des. 61-62, 391-397 (2002).
[2] M. Nakamichi, K. Yonehara and D. Wakai, Fusion Eng. Des.
86, 2262-2264 (2011).
[3] M. Nakamichi and K. Yonehara, J. Nucl. Mater. 417, 765-768(2011).
[4] J.-H. Kim, K. Yonehara and M. Nakamichi, Thermal
Stability of Plasma Sintered Beryllides, Proc. 15th Int. Conf.
on Fusion Reactor Materials, Charleston, USA, October
2011 (submitted to J. Nucl. Mater.).
[5] A. J. Maeland and G. G. Libowits, J. Less Common Met. 89,
197-200 (1983).
[6] D. W. Aylmore, S. J. Gregg and W. B. Jepson, J. Nucl. Mater.
2, No.2, 169-175 (1960).
[7] D. W. Aylmore, S. J. Gregg and W. B. Jepson, J. Nucl. Mater.
3, No.2, 190-200 (1961).
[8] R. A. Anderl, K. A. McCarty, M. A. Oates, D. A. Petti, R. J.
Pawelko and G. R. Smolik, J. Nucl. Mater. 258-263, 750-756(1998).
[9] K K. Wada, K. Munakata, M. Nakamichi, J. H. Kim,
K. Yonehara and D. Wakai, Oxidation Resistance of Be12Ti
Fabricated by Plasma-Sintering, Proc. 15th Int. Conf. on
Fusion Reactor Materials, Charleston, USA, October 2011
(submitted to J. Nucl. Mater.).
[10] Y. Sato, M. Uchida, H. Kawamura, High Temperature
Oxidation Behavior of Titanium Beryllide in Air, Proc. 6th
IEA Int. Workshop on Beryllium Technology for Fusion,
Miyazaki, Japan, December 2003, JAERI-Conf 2004-006,
p. 203.

Fig. 5 XRD spectra of central (a) and peripheral (b) parts ofthe plasma sintered titanium beryllide disc after exposure towater vapor.

and H2 gas generation. These results obviously suggest thatthe oxidation resistance for the central sample part of theplasma sintered beryllide disc is higher than that for the
peripheral part. Therefore, as mentioned above, larger
fraction of highly reactive Be in the titanium beryllide leads(i) to increase in the amount of H2 gas generated under
exposure to water vapor and (ii) to poor oxidation resistanceof the surface.

4. Conclusion

Center and peripheral parts (with different surface area
fractions of Be, Be12Ti, Be17Ti2, and Be2Ti phases) were cut
and extracted from a plasma sintered Be12Ti disk in order toinvestigate the effect of the surface area fraction difference
on not only the H2 gas generation under exposure to water
vapor but also oxidation property. It has been shown that theamount of hydrogen generated for the surface of the central
part of the titanium beryllide disc was smaller than that forthe peripheral part. This difference can be explained bylower fraction of the Be phase surface area on the surface ofthe central part in comparison with that for the peripheral
part. The oxidation property for both parts did not exhibit
big difference, while both titanium beryllides parts
demonstrated higher oxidation resistance compared tometallic beryllium.

5. Acknowledgments

This work was supported by Japan Atomic Energy
Agency under the Joint Work contract 23K131, as a part of101

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