1. to be submitted at the 2011 International Symposium on Rhenium held in July 4-8, 2011, Moscow, Russia
RHE IUM EFFECT I MOLYBDE UM-RHE IUM WELDS
F. Morito1, M. I. Danylenko and A. V. Krajnikov.
Institute for Problems of Materials Science (3, Krzhizhanivsky Street, 03680, Kiev, Ukraine)
1
MSF Laboratory (4-20-12, Keyakidai, Moriya, 302-0128, Japan)
Keywords: rhenium effect, Mo-Re alloys, electron-beam welds, sigma-phase, neutron
irradiation, radiarion-induced strengthening
Abstract: << MoRe_Report_Dec1_2010 by Sasha>>
This work analyses several Mo-Re alloys and welds with the Re content 0-50% in as-received
state and after electron beam welding and/or radiation treatment. The mechanical properties
and microstructure of Mo-Re welds are examined focusing on the effect of Re concentration.
Phase stability, microstructural changes and impurity redistribution are studied for better
understanding and predicting the long-term performance of Mo-Re alloys at high
temperatures and/or high neutron fluences.
In particular, a strategy of welding of Mo-Re alloys is discussed with emphasis on the
sensitivity of alloys to pre-weld heatings and on the development of post-weld treatments,
such as warm rolling and annealing, to provide optimal phase composition. Grain refinement
during directional solidification after welding, ductility improvement and fracture mode
change from intergranular to transgranular one are clearly observed with an increase of Re
content.
Effect of neutron irradiation on the strength of Mo-Re welds is studied for a wide temperature
range. Mo-Re welds exhibit a large radiation-induced strengthening. At room temperature, the
strengthening effect is rather limited and unstable because of lack of ductility. The
strengthening becomes strongly pronounced at high temperatures. Damaging effect of
neutrons at high temperatures is shown to be smaller than that at low temperatures.
Intensification of homogeneous nucleation of Re-rich sigma phases in all studied Mo-Re
alloys is observed after high temperature neutron irradiation. As a result, all parts of as-
irradiated welds display approximately same level of strength.
High-temperature annealings with different heating/cooling rates have been used to simulate
thermal conditions in different welding zones. Impurity redistribution in Mo-Re alloys has
been studied by surface analysis methods. The role of carbon and oxygen segregation as well
as formation of carbides and Re-base phases is discussed to minimise intergranular
embrittlement of welds. [14, Sasha]
1

2. Introduction
Rhenium is widely used to alloys VIA group refractory metals. In particular, the strength and
plasticity, creep resistance, and low temperature ductility of Mo are all improved with
increasing the rhenium content due to the so-called “rhenium effect” [1-5].
Fig. 1 shows change of microhardness as a function of Re content in Mo-Re alloys.
1. Solid solution type Mo-Re alloys
1-1. Solution softening and hardening in Mo-Re alloys
Fig. 1 Hardness of Mo-Re alloys annealed at 1873 K for 3.6 ks.
Fig. 2 Yield stress by bend test vs test temperature in Mo-Re welds.
Open marks denote preweld annealing at 1923 K, 1 h and closed ones postweld annealing at
1923 K, 1 h, respectively. ( ) means brittle fracture before yielding.
2

3. Fig. 2 shows yield stress by bend test vs test temperature in Mo-Re welds. Mo-41Re welds
had much higher yield stress than those of Mo-5Re and Mo-13Re, which failed in a brittle
manner at 77 K. But Mo-41Re welds showed more ductile and farctured after yielding even at
77 K.
Fig. 2 also shows a transition of solution softening by a test temperature. At 300 K, Mo-5
wt.% Re exhibited more solution softening than Mo-13 wt.% Re. However minimum yield
stress moved to Mo-13 wt.% Re from Mo-5 wt.% Re with a decrease of test temperature.
2. Previous report on Re effect
In the VI A group metals, the conditions of a resonance covalent bond are fulfilled in the best
way. A deep minimum corresponds to these metals in the plot of state density at Fermi level
(N(EF)). This kind of electron structure causes the peculisrities of structure and mechanical
properties specific for these metals [1]. (Milman et al)
Rhenium Effect by Korotaev et al [2-4]
1. Significant enhancement of low temperature plasticity
2. Reduction of the temperature of the viscous-brittle transition, Tv.b.
3. Suppression of brittle sliding fracture
4. Upgrading of the weldability of high rhenium alloys
Very interestiing characteristic changes in strength properties such as
5. Solution softening in low-rhenium alloys
6. Plastification and reduction of the temperature of the viscous-brittle transition temperature.
Tv.b.
7. Achievement of a remarkably high strength (σ0.1~ 8-9 x 103 MPa ) after deep deformation
(ε> 99 %) by rolling
8. Extraordinaly pronounced effects of solution hardening in the region T > 0.2 Tmelt
A a result, the phenomenology of the enhancement of the strength and plastic properties and
of the features of the deformation and hardening of transition metal-Re alloys. However,
Physical nature of Rhenium Effect has not been elucidated unambiguously up to the present.
(The subject matter of this work is a critical review of the physical nature of Rhenium Effect
and to elucidate the most probable mechanism for these phenomena.)
1. Y. V. Milman and G. G. Kurdyumova, Rhenium effect on the improving of mechanical
properties in Mo, W, Cr and their alloys (review), Proc. International Symposium on
Rhenium and Rhenium Alloys, B.D. Bryskin, ed. TMS, (1997)717-728
2. A. D. Korotaev, A. N. Tyumentsev and Yu. I. Pochivalovl, The rhenium effect in W- and
Mo-base alloys: The experimental regularities and the physical nature, Proc. International
Symposium on Rhenium and Rhenium Alloys, B.D. Bryskin, ed. TMS, (1997)661-670
3

4. 3. A. D. Korotaev, A. N. Tyumentsev, V. V. Manako and Yu. P. Pinzhun, The solubility of
oxygen in rhenium-alloyed molybdenum, Proc. International Symposium on Rhenium
and Rhenium Alloys, B.D. Bryskin, ed. TMS, (1997)671-680
4. A. N. Tyumentsev, A. D. Korotaev, Yu. P. Pinzhu and V. V. Manako, Dispersion and
substructure hardening of Mo-Re-base alloys, Proc. International Symposium on
Rhenium and Rhenium Alloys, B.D. Bryskin, ed. TMS, (1997)707-716
3. Sasha : In addition to traditional applications, such as heating elements, electron tube
components, etc, Mo-Re alloys are considered as candidate materials for structural
applications for chemicals and energy facilities, including elements of fusion or fast
breeder reactors. Welds are obligatory elements of practically any complex construction
while working conditions are characterized by high temperatures (>1000 K), aggressive
medium (liquid metals) and high neutron fluence (>1021 n/cm2).
Therefore, good weldability, high radiation performance, thermal stability and corrosion
resistance are the key issues for these application fields. Tendency of Mo alloys to embrittle at
low temperatures assumes probable degradation of the mechanical properties either during
welding [5-9] or under irradiation [10-13]. In spite of the fact that significant progress has
been achieved in studying the mechanism of rhenium effect, many details of Mo-Re alloy
behaviour under extreme operating conditions are not studied yet.
Table 1 Chemical content of various Mo-Re alloys (wt.ppm).
Re (%) :
0 2 4 10 13 15 20 25 30 40 41 47 50
ominal
Re (wt%) 0 1.7 3.3 8.8 12.0 15.9 21.4 25.6 31.7 37.0 43.6 47.1 50.1
Re (at%) 0 1.0 1.9 5.1 7.1 9.6 13.2 15.2 20.6 24.8 30.2 33.2 36.0
Al 11 7 9 11 <10 <10 <10 <10 83 <5 <10
Ca 5 4 7 7 <1 <1 <1 <1 14 <1 <1
Cr 9 8 9 9 <5 <5 <5 <5 10 <5 23 <5
Cu <3 <3 <3 <3 <5 <5 <5 <5 <3 <1 <5
Fe 20 50 30 50 60 10 10 10 10 150 38 5 10
Mg 2 2 3 3 <1 <1 <1 1 2 <1 <1
Mn <3 <3 <3 <3 <3 <1 <1 <1 <1 <3 <1 <1
i 5 11 10 13 11 <5 <5 <5 <5 20 <5 <3 <5
Pb <3 <3 <3 <3 <10 <10 <10 <10 <3 <10
Si 30 30 30 60 <10 <10 <10 <10 40 <10 <20 <10
Sn <3 <3 <3 <3 <10 <10 <10 <10 <3 <10 <10
a 1 1 1 1 <1 <1 <1 <1 1 <1
K 1 1 1 1 <1 <1 <1 <1 1 <1
C 5 12 7 4 4 10 10 10 10 12 10 10 <10
6 4 2 1 1 <1 <1 <1 <1 1 <10 6
O 18 23 15 14 20 11 10 7 8 66 <10 12
4

5. 4. Irradiation effect on Mo alloys and welds
4-1. Welds of Mo, TZM and Mo-0.56%Nb irradiated to 1017 cm–2 ~ 1020 cm–2 (E> 1 MeV) at
348-1073 K6 (irradiation by JRR-2 and JRR-4 at JAERI).[5]
I. Tensile properties
Mo and TZM, irradiated to 1.2 x 1024 n/m2 at 1073 K (1.2 × 1020 cm–2 (E> 1 MeV))
Fig.3. Tensile properties as a function of test temperature. (a) Tensile stress and (b) Total
elongation in as-welded PM Mo (▽), postweld annealed PM Mo (○), postweld carburized
PM Mo (□), postweld annealed TZM (△) and postweld annealed Mo-0.56 wt.% Nb (◇).
5

6. Fig. 4. Yield stress (△), tensile stress (○) and total elongation (□) of (a) as-welded PM Mo
and (b) postweld annealed TZM irradiated to 1.2 x 1024 m-2 at 1073 K. Closed marks denote
to post-irradiation annealing at 1273 K for 1 h.
For comparison,
PM Mo recrystallized at 1523 K for 1 h and irradiated to 1.3 x 1022 n/cm2 at LT
(Kazaakov & Chakin, 1993)
800
600 YS, Unirrdiated
YS, Irradiated
YS, MPa
400
200
Ef , Unirrdiated
0 Euni , Unirrdiated
50 Ef , Irradiated
40
Euni , Irradiated
30
E, %
20
10
0
300 400 500
T test
, K
Fig. 5. Mechanical properties of PM Mo (recrystallized at 1523 K, 1 h) tensile tested at 300 K,
Blue marks for unirradiated and red marks for irradiated to 1.3 x 10 22 n/cm2 at LT.
6

7. Mo-0.56%Nb, irradiated to 6.0 x 1023 n/m2 at 1073 K (6.0 × 1019 cm–2 (E> 1 MeV))
Total elongation <Tensile test at LT, RT & HT>
Mo-0.56%Nb: BM annealed at 1423K, 1h, 181K
Mo-0.56%Nb: as-welded <Ef~>2% limited to WM and HAZ rather ductile> 235K
Mo-0.56%Nb: postweld annealed at 1673K, 1h, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
<Ef~>20% ductile>
Mo-0.56%Nb: as-welded & irradiated, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
Mo-0.56%Nb: postweld annealed & irradiated, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
Embrittlement due to irradiation was not so significant at RT in the case of Mo-0.56%Nb !!
Fig. 6. Total elongation <Tensile test at LT, RT & HT>
(△) Mo-0.56%Nb: BM annealed at 1423K, 1h, 181K
(△) Mo-0.56%Nb: as-welded <Ef~>2% limited to WM and HAZ rather ductile>
235K
(▽) Mo-0.56%Nb: postweld annealed at 1673K, 1h, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
<Ef~>20% ductile>
7

8. (□) Mo-0.56%Nb: as-welded & irradiated, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
(○) Mo-0.56%Nb: postweld annealed & irradiated, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV)
443K
Embrittlement due to irradiation was not so significant at RT in the case of Mo-0.56%Nb !!
Mo-0.56%Nb
Neutron irradiated to 6.0 x 1023 n/m2 at 1073 K (6.0 × 1019 cm–2 (E> 1 MeV) at 1073 K).
Summary:
We examined mechanical properties of electron-beam welds of Mo and its alloys, TZM and
Mo-0.56wt.% Nb, for nuclear applications. The main results are as follows.
(1) It was shown that mechanical properties of Mo and its alloys were generally degraded by
electron-beam welding and further by neutron irradiation.
(2) Addition of carbon was effective to suppress intergranular embrittlement which was
caused by electron-beam welding and neutron irradiation. It is considered that segregation of
carbon and precipitation of carbides enhanced the intergranular cohesion so that considerable
strength and ductility was maintained.
(3) TZM showed higher strength at high temperatures, but it seemed unavoidable to improve
lower bend ductility especially in the case of neutron irradiation to 1.2 x 1024 n/m2 at 1073 K.
(4) The embrittlement of Mo-0.56wt.% Nb was not so significant at room temperature under
the irradiation to 6.0 x 1023 n/m2 at 1073 K. It is considered that good mechanical properties
of the unirradiated weld of Mo-0.56wt.% Nb were maintained in this case.
Welds of PM-Mo and TZM, irradiated to 1.2 x 1024 n/m2 at 1073 K (1.2 × 1020 n/cm–2 (E>
1 MeV)), F. Morito and K. Shiraishi, (1991)[5]
PM Mo recrystallized at 1523 K for 1 h and irradiated to 1.3 x 1026 n/m2 at LT (1.3 x 1022
n/cm2 at LT), (Kazaakov & Chakin, 1993)
Welds of Mo-0.56wt.% Nb, irradiated to 6.0 x 1023 n/m2 at 1073 K (6.0 × 1019 n/cm–2 (E>
1 MeV)), F. Morito and K. Shiraishi, (1991) )[5]
******************** Attention between Fluences *******************************
1. RIAR : Fig. X Mechanical properties of PM Mo (recrystallized at 1523 K, 1 h) tensile
tested at 300 K, Blue marks for unirradiated and red marks for irradiated to 1.3 x 10 26 m-2
at LT, 1.3 x 10 22 n/cm2 at LT.
2. Mo and TZM, irradiated to 1.2 x 1024 m-2 at 1073 K (1.2 × 1020 cm–2 (E> 1 MeV))
3. Mo-0.56%Nb: postweld annealed at 1673K, 1h, 6.0 x 1023 m-2 at 1073 K 1073K, 6.0 ×
1019 cm–2 (E> 1 MeV)
8

9. 2. RIAR (SM reactor)
2-a LT irradiation : Mo-15Re、20Re、30Re、41Re at 393-433 K to 3.6 - 6.0 × 1025 m–2 (E>
0.1 MeV)7, 8 , 3.6 - 6.0 × 1021 cm–2 (E> 0.1 MeV)7, 8.
2-b HT irradiation : Mo-15Re、20Re、30Re、41Re、50Re at 1023-1073 K to 5.5 - 7.3 ×
1025 m–2 (E> 0.1 MeV)8, 9, 5.5 - 7.3 × 1021 cm–2 (E> 0.1 MeV)8, 9.
[5] F. Morito and K. Shiraishi, Journal of Nuclear Materials, 179, (1991) 592-595.
[6] V.P. Chakin, F. Morito, V.A. Kazakov, Yu.D. Goncharenko and Z.E. Ostrovsky, Journal
of Nuclear Materials, 258-263 (1998) 883-888.
[7] F. Morito, V.P. Chakin, H. Saito, N.I. Danylenko and A.V. Krajnikov, Proc. 17th
International Plansee Seminar, P. Rhoedhammer et al. (Eds.), Reutte/Tirol, vol. 1,
(2009) RM63.
[8] F. Morito , V. P. Chakin, M. I. Danylenko and A. V. Krajnikov, J. Nuc. Mater.,
**(2010)***
III. Irradiated Mo-Re welds at HT 7, 8 <HT irradiation>
20
T = 3 0 0 K
Unirradiated, SR
Unirradiated, Rec
Elongation, %
Irradiated, SR
Irradiated, Rec
10
0
0 10 20 30 40
Re content, %
Fig. 7. Fracture elongation tested at 300 K in Mo-Re welds. (Open mark=Recrystallized
(Rec), Closed mark=Stress-relieved (SR). : unirradiated, ----- : irradiated).
9

10. T = 1 0 2 3 K /1 0 7 3 K
30
Elongation, %
20
10 Unirradiated, SR
Unirradiated, Rec
Irradiated, SR
Irradiated, Rec
0
0 10 20 30 40
Re content, %
Fig. 8. Fracture elongation tested at 1023 K for unirradiated and 1073 K for irradiated in Mo-
Re welds. (Open mark=Recrystallized (Rec), Closed mark=Stress-relieved (SR). :
unirradiated, ----- : irradiated).
10

11. (a)
1600
Unirradiated, SR
1400
Unirradiated, Rec
Tensile strength, MPa
1200 Irradiated, SR
Irradiated, Rec
1000
800
600
400
200
0
0 10 20 30 40 50
Re content, %
(b)
1600
Unirradiated, SR
1400 Unirradiated, Rec
Irradiated, SR
Tensile strength, MPa
1200 Irradiated, Rec
1000
800
600
400
200
0
0 10 20 30 40 50
Re content, %
Fig. 9. Tensile strength tested (a) at 300 K and (b) at 1023 K for unirradiated and 1073 K for
irradiated in Mo-Re welds. (Open mark=Recrystallized (Rec), Closed mark=Stress-relieved
(SR). : unirradiated, ----- : irradiated).
1400 Unirrad., SR, BM
Unirrad., Rec, BM
Irrad., SR, WM
1200 Irrad., SR, HAZ
Irrad., SR, BM
Irrad., Rec, WM
Microhardness
1000 Irrad., Rec, HAZ
Irrad., Rec, BM
800
600
400
200
0
0 10 20 30 40 50
Re content, %
Fig. 10. Hardness at 300 K in Mo-Re welds. (Open mark=Recrystallized (Rec), Closed
mark=Stress-relieved (SR). : unirradiated, ----- : irradiated).
11

12. Fig. 11. Density change (∆d/dini) in Мо-Re welds irradiated at Тirr = 1023-1073 K. ( ∆d=dirr -
dini , dini : Density before irradiation, dirr : Density after irradiation).
(a) Unirrad. BM, (b) Unirrad. WM
Fig. 12. Unirradiated microstructutre of Mo-Re welds (postweld annealed at 1673 K, 1 h) : (a)
Mo-15Re: Unirrad. BM, postweld, (b) Mo-15Re: Unirrad. WM.
In the initial state of the Mo-Re welds, for example, only dislocations, dislocation networks
and dislocation tangles with rather low density are seen (Fig. 12-11 (a ) and (b)).
12

13. 30Re 50Re
BM
(a) (d)
HAZ
(b) (e)
40µ
m WM
(c) (f)
Fig. 13. Optical micrography of Мо-Re welds irradiated at Тirr = 1023-1073 K. Mo-30Re
(1173 K, 1 h): (a) BM, (b) HAZ, (c) WM and Mo-50Re (1173 K, 1 h): (d) BM, (e) HAZ, (f)
WM
13

14. (a) (b) (c)
40 µm
(d) (e)
Fig. 14. Fracture surfaces of postweld recrystallized Mo-Re welds irradiated at Тirr = 1023-
1073 K. (a) Mo-15Re, (b) Mo-20Re, ( c) Mo-30Re, (d) Mo-41Re, (e) Mo-50Re.
14

15. 4 µm 4 µm
a b
Particles of
second
phase
4 µm 4 µm
c d
Fig. 15. Fracture surface of Мо-Re welds irradiated at Тirr = 1023-1073 K after tensile test at
room temperature:
a) Mo-15Re; 1673 K, 1h; BM b) Mo-20Re; 1673 K, 1h; WM
c) Mo-30Re; 1673 K, 1h; BM d) Mo-41Re; 1673 K, 1h; BM
15

16. 8 µm
40 µm
a
8 µm
40 µm
b
8 µm
40 µm
c
Fig. 16. Metallography of Mo-20Re weld; 1673 K, 1h:
a) BM
b) HAZ
c) WM
16

17. 90 nm
a
90 nm
b
150 nm
c
Fig. 17. TEM microstructure of Мо-Re welds irradiated at Тirr = 1023-1073 K.
a) particles of second phase, Mo-15Re; 1673 K, 1h;
b) particles of second phase, Mo-41Re; 1673 K, 1h;
c) grain boundary, Mo-41Re; 1673 K, 1h
17

18. 1000
800
600
Hv
400
200
0
0 10 20 30 40 50
Re content, %
Fig. 18. Microhardness of Mo-Re welds (postweld annealed at 1173 K, 1 h) :
Irradiation at Тirr = 1023-1073 K up to fluence F=(5.5-7.3)×1021 cm-2 (Е>0.1 MeV).
- initial state ○ - irradiated
3. Results <HT irradiation>
The tensile strength of both irradiated and unirradiated welds is shown in Fig. 1a and Fig. 1b
as a function of Re content for room temperature and high temperature tests respectively.
Almost all welds failed in a very brittle manner at room temperature. But all the welds
demonstrated adequate ductility with elongation ~20-30% at high temperatures. Although
some variations of strength data were recognized in irradiated samples, tensile strength
generally increased with Re content. After neutron irradiation, Mo-Re welds showed a large
radiation-induced strengthening. At room temperature, the strengthening effect was rather
limited and unstable because of lack of ductility. It appeared in some samples with residual
ductility but was absent in absolutely brittle samples. The strengthening became strongly
pronounced at high temperatures. In particular, tensile strength of all irradiated welds at
1073 K was sufficiently higher than that of unirradiated specimens at 1023 K. For example,
tensile strength of Mo-16Re and Mo-45Re alloys in Rec state increased from 240 to 960 MP
and from 420 to 1250 MPa respectively. The radiation-induced strengthening at high
temperature was estimated to be about 700-800 MPa and rather independent on the Re content.
Microhardness of irradiated and unirradiated welds is shown in Fig. 2. The hardness of all
zones of welds, such as weld metal (WM), heat-affected zone (HAZ) and base metal (BM),
usually increased with the Re content. For any irradiated specimen, the difference in hardness
18

19. between different welded zones was very small, if any. Each zone of all irradiated welds
exhibited a severe radiation-induced hardening compared to that of unirradiated samples. An
increase of the microhardness resulted from high temperature irradiation was approximately
300 MPa for SR welds and varied between 400 and 600 MPa for Rec samples depending on
the Re content. Fig. 3 shows an increase of density (∆d/ dini) in Mo-Re welds after high
temperature irradiation. Density of irradiated welds (dirr = dini+∆d) increased to 3.8-9.1%
compared with their initial density (dini). The value of ∆d/dini is almost constant for Mo-Re
welds with 16-21% Re. Further increase of Re content monotonously weakens the observed
effect of density growth. As shown in Fig. 4, massive secondary phases were recognized as
the form of thin plane particles with 70-160 nm length in all the Mo-Re welds. It indicates
that microstructure reconstruction of irradiated Mo-Re welds is connected with the formation
of secondary phases of much higher density compared with that of unirradiated specimen.
19

20. IV. Irradiated Mo-Re welds at LT 6, 7 <LT irradiation>
XT Legend
B
30 30 C
R
D
20 20
YR
E , %
f
10 10
0 Line 0
0 10 20 30 40
XT
1000
Text
B
C
800 R
D
YL
600 U n ir r a d ia te d
σYS , MPa
Text2
YR
400
Arrow1
Arrow
Ir r a d ia te d
200 Text1
Arrow2 Arrow3
0
0 10 20 30 40
XB
Re content, %
Fig. 19. Result of Mo-Re welds tensile tested at 293 K and 673 K. (a) Total elongation (δtot),
(b) Yield stress (σYS). △：Unirrad. postweld annealed at 1173 K, Ttest = 293 K, ▲：Irrad.
postweld annealed at 1673 K, Ttest = 293 K, ○：Unirrad. postweld annealed at 1673 K, Ttest =
293 K, ●：Irrad. postweld annealed at 1673 K, Ttest = 293 K, ◇：Unirrad. postweld annealed
at 1173 K, Ttest =673 K, 1173 K, ◆：Irrad. postweld annealed at 1173 K, Ttest = 673 K, □：
Unirrad. postweld annealed at 1673 K, Ttest = 673 K, ■：Irrad. postweld annealed at 1673 K,
Ttest = 673 K, br: brittle rupture.
20

21. (a) (b) (c)
Fig. 20. Macrostructure of tensile specimen fractured at WM. (a) Mo-15Re fractured during
disassembling of capsule, (b) Mo-20Re fractured during disassembling of capsule, (c) Mo-
30Re fractured by tensile test at RT.
21

22. (a) (b)
(c) (d)
(e) (f)
Fig. 21. Microstructure of postweld annealed Mo-Re welds irradiated at LT.
(a) Mo-20Re: Irrad. BM, (b) Mo-20Re: Irrad. WM (postweld annealed at 1673 K, 1 h)
(c) Mo-30Re: Irrad. WM, (d) Mo-30Re: Irrad. WM (postweld annealed at 1173 K, 1 h)
(e) Mo-41Re: Irrad. BM, (f) Mo-41Re: Irrad. WM (postweld annealed at 1173 K, 1 h)
In recryatallized Mo-Re welds, only dislocations, dislocation networks and dislocation tangles
with rather low density are seen (Fig. 12-11 (a ) and (b)). Dislocation density in WM was less
than in the BM. So, the dislocation density in BM was 2.9 x 1010 cm-2 and that in WM was 1.9
x 1010 cm-2 for the Mo-15Re weld.
The results of TEM investigations of irradiated Mo-Re welds are presented in Fig. 12-13 (c) -
(f) and in Table 3. Iirradiation leads to the formation of the dislocation loops typical for such
low-temperature irradiation. Their average size is 7.5 ~ 10 nm. Density of dislocation loops
22

23. varied in the range from 4.5 ´ 1015 to 2.6 ´ 1016 cm-3 and decreased with an increase of
rhenium content.
The structure of all irradiated Mo-Re alloys is retained as one phase. The second phase
precipitations, which were remarkably observed in Mo-Re welds irradiated at higher
temperatures [1], were not detected.
Dislocation loop distribution in BM of Mo-41Re welds was extremely irregular, but
dislocation loops were almost absent in WM as in Fig. 12-13 (e), (f). The reason is to be
solved why these phenomena are mainly due to Re effect.
12
10 WM 2
Loop Size / nm
30Re 20Re
1
8 BM 3 4
WM BM
5
BM
6 41Re
4 15 16 17
10 10 10
Loop Density / cm-3
Fig. 22. Dislocation loop density vs size in Mo-Re welds irradiated at LT.
5. Conclusions
1. Irradiation of the tensile specimens and TEM disksof the welds of Mo±alloys with 15%,
20%, 30% and 41%rhenium contents at 120±160°C to the neutron ¯uence of 6.0 x 1021 n/cm2
(E > 0.1 MeV) led to the strong radiation embrittlement. The fracture took place only over the
specimens centre through the weld-fusion zone. With increasing rhenium content the fracture
type changed from the brittle intergranular type to the transgranular one.
23

24. 2. It was shown that with increasing rhenium contentthe dislocation loop density with an
average size of 7.5~10 nm was reduced 4~6 times and in the fusion zone of Mo-41Re welds
were quite absent.
V. Summary <HT irradiation>& <LT irradiation>
Radiation-Induced Segregation, Precipitation, Hardening, Embrittlement, Transmutation
11. RIP
Nucleation sites of σ-phase are considered to be subgrain boundaries consisting of tangled
dislocations, because the number of density of subgrains is changed with thermal treatment.
Consequently it is considered that the number of density of σ-phase precipitations can be
controlled by thermal treatment.
On the other hand, the mean length and the number of density of χ-phase precipitates were not
changed with thermal treatment. This suggests that the nucleation site of χ-phase are
dislocations and dislocation loops formed at the beginning of irradiation ( < 1 dpa ) **.
**. B. N. Singh, J. H. Evans, A. Horsewell, P. Toft and G. V. Muller, J. Nuc. Mater., 258-
263(1998)865
After nucleation of these σ-phase and χ-phase precipitates, the surface of the precipitates acted
as a sink for under size elements such as Re atoms, leading to the growth of the precipitates.
Nelson et al124 : satulated radius and number of density of sphere shaped RIP gamma prime
phase precipitates in Ni-Al alloy -> Fig. 11
12. RIH and RIE by Orowan model126-128 -> Fig.12: Re content dependence of radiation
hardening calculated from results of the microstructural observation. Calculated Hv agreed
with the measurement in the specimens irradiated at 1072 K. In the specimens irradiated at
681 or 874 K, the scatter of the calculation became larger. It is considered because there
would be invisible small defects at lower temperature irradiation.
12*. RIE
For Mo-41Re irradiated at 874 K or below, there were cracks observed around the indentation
after Hv measurements. This embrittlement is thought to be caused by large and hard σ-phase
precipitates. σ-phase has very high HV value of about 1500 (MPa), which is much larger than
that of the matrix, and hence the surface or inside of such hard precipitates would be the
initiation sites of cracks. Thus the formation of large σ-phase ppts led to drastic embrittlement
in the Mo-41Re.
2. χ-phases which are usually close to spheroids by ion irradiation [20]
3. Ageing at 1098 and 1248 K of two-phase Mo-47.5 wt% Re with αMo + σ structure was
recently shown to form χ-phase along grain boundaries [24].
24

25. [24] K.J. Leonard, J.T. Busby and S.J. Zinkle, Journal of Nuclear Materials, 366, (2007) 369-
387
Microstructural and mechanical property changes with aging of Mo–41Re and Mo–
47.5Re alloys
The changes in microstructure and mechanical properties of Mo–41Re and Mo–47.5Re alloys
were investigated following 1100 h thermal aging at 1098, 1248 and 1398 K. The electrical
resistivity, hardness and tensile properties of the alloys were measured both before and after
aging, along with the alloy microstructures though investigation by optical and electron
microscopy techniques.
(1) The Mo–41Re alloy retained a single-phase solid solution microstructure following
1100 h aging at all temperatures, exhibiting no signs of precipitation, despite measurable
changes in resistivity and hardness in the 1098 K aged material.
(2) Annealing Mo–47.5Re for 1 h at 1773 K resulted in a two-phase αMo + σ structure,
(3) with subsequent aging at 1398 K producing a further precipitation of the σ phase along
the grain boundaries. This resulted in increases in resistivity, hardness and tensile strength
with a corresponding reduction in ductility.
(4) Aging Mo–47.5Re at 1098 and 1248 K led to the development of the χ phase along grain
boundaries, resulting in decreased resistivity and increased hardness and tensile strength
while showing no loss in ductility relative to the as-annealed material.
RIT : [25] E.J. Edwards, F.A. Garner and D.S. Gelles, Journal of Nuclear Materials, 375,
(2008) 370-381.
4-1. Mo–41 wt% Re irradiated in the fast flux test facility (FFTF) experienced significant and
non-monotonic changes in density due to radiation-induced segregation, leading to non-
equilibrium phase separation, and progressive transmutation of Re to Os [25].
4-2 irradiation of Mo–41 wt% Re over a range of temperatures (743-1003 K) to 28–96 dpa
produced a high density of thin platelets of a hexagonal close-packed (hcp) phase identified
as a solid solution of Re, Os and possibly a small amount of Mo.
4-2* < Grain boundaries are also enriched with Re to form the hcp phase, but the precipitates
are much bigger and more equiaxed in shape.>
4-3. Although not formed at a lower dose, continued irradiation at 1003 K leads to the co-
formation of late-forming χ-phase, an equilibrium phase that then competes with the pre-
existing hcp phase for rhenium.
4. Discussion <ICFRM-14=JNM2010> 9
In accordance with the phase diagrams [18, 19], formation of σ-phases and χ-phases takes
place at high temperatures in Mo-Re alloys with a high Re content. Formation of secondary
phases under irradiation in solid solution is possible due to radiation-induced segregation of
Re atoms followed by Re enrichment of internal sinks up to solubility limit or even exceeding
it. We already reported that σ-phase was recognized in Mo-50Re alloys and welds [15, 16].
We also showed that thermo-mechanical treatment was much effective to improve mechanical
25

26. properties of Mo-50Re welds by controlling dispersion and size of σ-phase along grain
boundaries and in the matrix [16, 17]. The results obtained here also correspond well with the
previous ones [20-22]. As was shown here, formation of secondary phases is significant and
intensive in Mo-Re welds. Character and rate of phase formation are almost independent on
the Re content as well as on the location within the welded zone. Secondary phases look like
thin plane particles located along some crystallographic planes as shown by TEM. Based on
the particle shape, one may conclude that they are σ-phases rather than χ-phases which are
usually close to spheroids [20]. High intensity of secondary phase formation, homogeneous
distribution of the particles over the bulk, and absence of depleted zones along grain
boundaries indicate that the phases were formed by means of mechanism of homogeneous
nucleation which is based on the radiation-induced separation of atoms in alloys [23]. In this
case, formation of stable nuclei can occur in defect-free areas of welds. In other words,
nucleation occurs homogeneously over the grain bulk and it does not preferentially correlate
with structure defects such as voids, dislocations, grain boundaries and so on. Susceptibility
of Mo and Re atoms to separation under irradiation is high and stability of the formed nuclei
is sufficient to continue growth of particles as a consequence of radiation-induced segregation.
Independence of phase formation rate on the Re content supports our assumption about a very
high intensity of radiation-induced separation of Mo and Re atoms in Mo-Re welds under
neutron irradiation. This eliminates effect of alloy composition on the final microstructure and
leads to formation of practically the same number of secondary phases in the welds in spite of
Re content. Ageing at 1098 and 1248 K of two-phase Mo-47.5 wt% Re with αMo + σ
structure was recently shown to form χ-phase along grain boundaries [24]. Mo–41 wt% Re
irradiated in the fast flux test facility (FFTF) experienced significant and non-monotonic
changes in density due to radiation-induced segregation, leading to non-equilibrium phase
separation, and progressive transmutation of Re to Os [25]. Beginning as a single-phase solid
solution of Re and Mo, irradiation of Mo–41 wt% Re over a range of temperatures (743-
1003 K) to 28–96 dpa produced a high density of thin platelets of a hexagonal close-packed
(hcp) phase identified as a solid solution of Re, Os and possibly a small amount of Mo. These
hcp precipitates are thought to form in the alloy matrix as a consequence of strong radiation-
induced segregation to Frank loops. Grain boundaries are also enriched with Re to form the
hcp phase, but the precipitates are much bigger and more equiaxed in shape. Although not
formed at a lower dose, continued irradiation at 1003 K leads to the co-formation of late-
forming χ-phase, an equilibrium phase that then competes with the pre-existing hcp phase for
rhenium.
After low temperature irradiation, fracture surface is mainly located in WM, because this zone
usually has the lowest strength among other welded zones [7-9].
Mo-Re welds have a tendency to exhibit intergranular embrittlement in alloys with a lower Re
content and under low temperature irradiation.
In contrast to low temperature irradiation, all zones of welds after high temperature irradiation
are characterized by comparable strength level. Thus, fracture at room temperature occurred
at any place of the welds by brittle transgranular cleavage.
Numerous secondary phases were observed at cleavage lines and within matrix. In addition
the number and morphology of these particles were almost independent on the Re content.
The reason of this effect is the same as discussed above.
26

27. Intensive phase formation in Mo-Re welds eliminated any difference in mechanical properties
between WM, HAZ and BM. Mo-Re welds with such a characteristic microstructure still kept
sufficient plasticity at high temperatures. Therefore it is clear that high temperature neutron
irradiation of Mo-Re welds have lesser damaging effect than low temperature irradiation.
As a result, we obtained a very important practical conclusion that weld-fabricated
constructions operating under high temperature radiation do not limit their life because
mechanical properties of welds with WM and HAZ are not worse than those of BM.
Concluding remarks
1. The presence of σ-phases is effective to suppress embrittlement of two-phase precipitated
Mo-Re alloys and welds, controlling their population such as distribution, size and density
combined with thermo-mechanical treatments (TMT). As a result, it is possible to enhance
Re effect keeping lower temperature ductility and higher temperature strength.
2. Weldability was much improved among Mo-Re alloys with higher Re contents, which is
exhibited as significant role by Re effect both in unirradiated and irradiated conditions.
3. Due to irradiation, Mo-Re alloys in solid solution signicantly formed σ-phases. Mo-Re
alloys of two-phase type further enhanced precipitation and growth of σ-phases. With an
increase of irradiation, σ-phases precipitation much increased by promoting radiation-
induced segregation and radiation-induced precipitation.
4. Radiation-induced segregation, radiation-induced precipitation, radiation-induced
hardening was recognized during neutron irradiation of Mo-Re alloys and welds. Among
Mo-Re with higher Re contents, radiation-induced strengthening was considerably
promoted at higher temperature irradiation. But radiation-induced embrittlement at RT
remains to be overcome in the case of lower temperature irradiation less than 400 K.
5. Problems on density increase of Mo-Re welds after irradiation are desirable and raise a
good discussion about Re effect, I think.
27

28. References-x
1. R. EcK, Proc.11th International Plansee Seminar, 1(1985)39
2. G. A. Geach and J. R. Hughes, Proc. 2nd Plansee Seminar, Pergamon Press, London,
(1956)245
3. Von. I. Glass and G. Bohm, Planseeberichte. 10(1962)144
4. A. Lawlay and R. Maddin, Trans. Met. Soc. AIME, 224(1962)573
5. D. L. Davidson and F. R. Brotzen, Acta Met. 18(1970)463
6. W. D. Klopp and W. R. Witzke, Met.Trans. 4 (1973)2006
7. L. B. Lundberg, E. K. Ohriner, S. M. Tuominen, E. P. Whelan and J. A. Shields, Jr.
Physical Metallurgy and Technology of Molybdenum and Its Alloys, Eds. K. H. Miska et al.
AMAX, Ann Arbor, MI, (1985)71
8. J. Wadsworth, C. M. Packer, P. M. Chewey and W. C. Coons, Mater. Sci. Eng.
59(1983)257
9. F. Morito, J. Less-Comm. Met. 146(1989)337
10. F. Morito, Proc. 12th Intern. Plansee Seminar, 1(1989)417
11. F. Morito, J. Phys. IV, 1(1993)553
12. F. Morito, JOM, 45(1993)54
13. F. Morito, Proc. 13th Int. Plansee Seminar, 1(1993)585
14. F. Morito, High Temp. & High Press. 26(1994)101
15. T.B. Massalski, Eds. Alloy Phase Diagrams, vol. 2, ASM International, Metals ParK, OH,
(1986)
16. J. Wadsworth, T. G. Nieh and J. J. Stephens, Scripta Met. 20(1986)637
17. R. EcK, Planseeberichte. 22(1974)165
18. T. Takebe, T. Akiyama, K. Shimatani and M. Endoh, Physical Metallurgy and
Technology of Molybdenum and Its Alloys, Eds. K.H.Miska et al. AMAX,Ann Arbor,MI,
(1985)41
19. C. M. McNally, W. H. Kao and T. G. Nieh, Scripta Met. 22(1988)1847
20. J. Schrank and M. Kuhlein, Refractory and Hard Metals, 6(1987)106
21. F. Morito and K. Shiraishi, J. Nuc. Mat. 179-181(1991)592
22. S. A. Fabritsiev, V. A. Gosudarenkova, V. A. Potapova, V. V. Rybin, L. S. Kosachev, V.
P. Chakin, A. S. Pokrovsky and V. R. Barabash, J.Nuc.Mat. 191-194(1992)426
23. T. Suzuki and I. Mutoh, J. Nuc. Mat. 184(1991)81
28

29. 24. T. Suzuki and I. Mutoh, Proc. Workshop on High Temperature Corrosion of Advanced
Materials and Protective Coatings," Eds. Y. Saito et al. Elsevier Sci. Pub. B. V. ToKyo,
(1992)227
25. T. Flament and J. Sannier, Proc. 4th Intern. Conf. on Liquid Metal Engineering and
Technology, Societe Francaise d'Energie Nucleaire and Commissariat a l'Energie Atomique,
Avingon, France, 2(1988)520/1
26. J. A. Shields Jr, Proc. 1992 Powder Metallurgy World Congress, Metal Powder
Federation, Princeton, NJ, 8(1992)199
27. J. A. Shields Jr, Adv. Mat. Proc. 1421(1992)28
28. A. J. Bryhan, Adv. Mat. Proc. 1431(1993)50
29. F. Morito, J. Nuc. Mat. 165(1989)142
30. F. Morito, J. Mat. Sci. 24(1989)3403
31. F. Morito, Proc.12th International Plansee Seminar'89, 1(1989)313
32. F. Morito, Vacuum, 41(1990)1743
33. F. Morito, Colloque de Physique, C1-51(1990)281
34. F. Morito, Surf. Interf. Analysis, 15(1990)427
35. N. Igata, A. Kohyama and K. Itadani, J. Nuc. Mat. 85-86(1979)895
36. A. Kohyama and N. Igata, J. Nuc. Mat. 122(1984)767
29

30. References-y
1. J. W. Blossom, “Annual Report 1990”, U. S. Department of the Interior, Bureau of Mines,
April 1992
2. J. Wadsworth and J. P. Wittenauer, LLNL Report, 1993
3. 五十嵐廉、まてりあ、41(2002)362
4. S. Primig, H. Leitner, A. Rodriguez-Chavez, H. Clemens, A. Lorich, W. Knabl, R. Stickler,
Proc. 17th International Plansee Seminar, G. Kneringer et al. Eds. Reutte/Tirol, 1(2009)RM15
5. M. Endo, K. Kimura, T. Udagawa and H. Seto, Proc. 12th International Plansee Seminar, G.
Kneringer et al. Eds. Reutte/Tirol, 1(1989)37
6. C. Stallybrasse, G. Leichtfried and A. Kneissl, Proc. 15th International Plansee Seminar, G.
Kneringer et al. Eds. Reutte/Tirol, 1(2001)267
7. 瀧田朋広、五十嵐廉、土肥義治、粉体および粉末冶金、45(1998)969
8. 瀧田朋広、五十嵐廉、馬淵守、中村守、土肥義治、長柄毅、粉体および粉末冶金、
46(1999)1031
9. 瀧田朋広、五十嵐廉、斎藤尚文、中村守、粉体および粉末冶金、46(1999)1261
10. 五十嵐廉、粉体および粉末冶金、49(2002)163
11. Y. Kitsunai, H. Kurishita, M. Narui, H. Kayano, Y. Hiraoka, J. Nuc. Mat. 239(1996)253
12. H. Kurishita, Y. Kitsunai, T, Kuwabara, M. Hasegawa, Y. Hirao Ka, T. Takida and T.
Igarashi, J. Nuc. Mat. 75(1999)594
13. Y. Kitsunai, H. Kurishita, T. Kuwabara, M. Narui, M. Hasegawa, T. Takida and K.
Takebe, J. Nuc. Mat. 346(2005)233
14. H. Kurishita, S. Satou, H. Arakawa, T. Hirai, J. Linke, M. Kawai and N. Yoshida,
Advanced Materials Research, 59(2009)18
15. 藤井忠行、北海道大学学位論文、1984
16. T. Fujii, R. Watanabe , Y. Hiraoka and M. Okada, J. Less-Common Met. 96(1984)297
17. T. Fujii, R. Watanabe , Y. Hiraoka and M. Okada, J. Less-Common Met. 97(1984)163
18. K. Tsuya and N. Aritomi, J. Less-Common Met. , 15(1968)245
19. J. P.Touboul, L. Minel and J. P. Langeron, J. Less-Common Met. 30(1973)279
20. Y. Hiraoka, F, Morito, M. Okada and R. Watanabe, J. Nucl. Mat. 78(1978)192
21. S. Suzuki, H. Matsui and H. Kimura, Mat. Sci. Eng. 47(1981)209
22. D. S. Wilkinson, K. Abiko, N. Thyagarajan and O. P. Pope, Met. Trans. 11A(1980)1827
23. 木村宏, 日本金属学会報, 22(1983)85
30

31. 24. M. P. Seah, Acta Met. 28(1980)955
25. W. Losch, Acta Met. 27(1979)1885
26. C. L. Briant and R. P. Messmer, Phil. Mag. B42(1980)569
27. 小川孝好, 石田洋一, 日本金属学会誌, 43(1979)1048
28. 安彦兼次, 木村宏, 鈴木茂, 鉄と鋼, 69(1983)625
29. M. Oku, S. Suzuki H. Kurishita and H. Yoshinaga, Appl. Surf. Sci. 26(1986)42
30. A. Lawlay, J. Van den Sype and R. Maddin, J. Inst. Metals, 23(1962-63)23
31. D. J. Capp, H. W. Evans and B. L. Eyre, J. Less-Common Met. , 40(1975)9
32. 及川秀男, 材料科学, 20(1984)36
33. L. E. Olds and G. W. P. Rengstorff, J. Metals, 8(1956)150
34. J. B. Brosse, R. Fillet and M. Biscondi, Scripta. Met. 15(1981)619
35. A. Kumar and B. L. Eyre, Proc. Roy. Soc. Lond. A370(1980)431
36. 野田哲二, 貝沼紀夫, 岡田雅年, 日本金属学会誌, 48(1984)604
37. H. Kurishita, H. Yoshinaga, K. Abi Ko, S. Suzuki and H. Kimura, Proc. 4th Int. Symp. on
Grain Boundary Structure and Related Phenomena, Supplement to Trans. JIM, 27(1986)739
38. G. Lorang, These Doctorat d'Etat, Orsay, Dec. 1980
39. G. Lorang, P. Ailloud, L. Debove, J. C. Rouchaud, M. Fedoroff and J. P. Langeron, Z.
Metallkde. 74(1983)458
40. A. Kobylanski et C. Goux, C. R. Acad. Sci. Paris, 272C(1971)1937
41. J. M. Jardin, A. Kobylanski et C. Goux, C. R. Acad. Sci. Paris, 280C(1975)717
42. 栗下裕明, 久芳俊一, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)539
43. 栗下裕明, 大石朗, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)546
44. J. M. Vitek and M. Ruhle, Acta Met. 34(1986)2085
45. A. Kumar and B. L. Eyre, Acta Met. 26(1978)569
46. T. W. Haas and J. T. Grant: Appl. Phys. Let. 16(1970)172
47. J. T. Grant and T. W. Haas, Surf. Sci. 24(1971)332
48. E. Gillet, J. Less-Common Met. , 71(1980)277
49. M. Frorjancic, M. Ruhle and S. L. Saas, Proc. 10th Int. Congr. Electron Microscopy,
2(1982)359
50. J. M. Penisson, M. Bacia and M. Biscondi, Phil. Mag. A, 73(1996)859
51. I. F. Fergson, J. B. Ainscough, D. Morse and A. W. Miller, Nature, 02(1964)1327
52. H. Lux and A. Ignatowicz, Chem. Ber. 101(1968)809
31

32. 53. V. I. Arkharov, V. N. Koney and A. F. Gerasimov, Fiz. Metal. Metalloved. 9(1960)695
54. P. Ettmayer, Monatsh. Chem. 101(1970)127
55. K. Kunimori, T. Kawai, T. Kondow. T. Onishi and K. Tamaru, Surf. Sci. 54(1976)525
56. T. Kawai, K. Kunimori, T. Kondow, T. Onishi and K. Tamaru, Phy. Rev. Let.
33(1974)533
57. H. Mi Ki, K. Kato, A. Kawana, T. Kioka, S. Sugai and K. Kawasaki, Surf. Sci.
161(1985)446
58. A. Ignatiev, F. Jona, D. W. Jepsen and P. M. Marcus, Surf. Sci. 49(1975)189
59. H. Kurishita, H. Yoshinaga, K. Abi Ko, S.Suzuki and H. Kimura, Trans. JIM,
27(1986)739
60. H. Kurishita, O.Tokunaga and H.Yoshinaga, Mater. Trans. JIM, 31(1990)190
61. M. Taheri, Mat. Sci. Eng. 32(1978)221
62. T. Noda and M. Okada, Trans. JIM. 28(1987)517
63. M. K. Miller, E. A. Keni K, M. S. Mousa, K. F. Russel and A. J. Bryhan, Scripta Met.
46(2002)299
64. 斎藤秀雄, 森藤文雄, 鐵と鋼, 89(2003)750
65. F. Morito, N.I. Danylenko, H. Saito and A.V. Krajnikov, Metallic Materials with High
Structural Efficiency, O. N. Senkov, D. B. Miracle and S. A. Firstov, eds. Dordrecht,
Netherlands, Kluwer Academic Publishers, (2004)347
66. 斎藤秀雄, 森藤文雄, 材料工学領域における最新オートラジオグラフィーの解析と
その応用, 編著:斎藤秀雄, うらべ書房, 千葉, (2008)101
67. F. Morito, V. P. Chakin, H. Saito, A. V. Krajnikov and M. I. Danylenko, Proc. 17th
Plansee Seminar, G. Kneringer et al. Eds. Reutte, Tirol, Austria, 1(2009)RM63
32

33. References-a
1. W. N. Platte, Weld. J. 35(1956)369s
2. W. N. Platte, Weld. J. 36(1957)301s
3. M. M. Nerodenko, E. P. Polishchuk and T. D. Galinzovskays, Autom. Weld. 32(1979)39
4. M. M. Nerodenko, E. P. Polishchuk and T. D. Galinzovskays, Autom. Weld. 32(1979)42
5. A. J. Bryhan, Welding Research Council, Bulletin-312, 1986
6. J. L. Jellison, Sandia National Laboratories Report SAND79-0193, (1979)
7. F. Matsuda, M. Ushio, K. Nakata and Y. Edo, Trans. JWRI. 8(1979)217
8. R. EcK, Proc. 11th Intern. Plansee Seminar, 1(1985)39
9. R. Kishore and A. Kumar, J. Nucl. Mat.. 101(1981)16
10. Y. Hiraoka, M. Okada and H. Irie, J. Nucl. Mat. 1555-157(1988) 381
11. Y. Hiraoka, Mater.Trans. JIM, 31(1990) 861
12. C. V. Robino, PhD dissertation, Lehigh Univ. 1988
13. N. Igata, A. Kohyama and K. Itadani, J. Nuc. Mat. 85-86(1979)895
14. A. Kohyama and N. Igata, J. Nuc. Mat. 122(1984)767
15. J. Wadsworth, G. R. Morse and P. M. Chewey, Mater. Sci. Eng. 59(1983)257
16. J. B. Brosse, R. Fillet and M. Biscondi, Scripta. Met. 15(1981)619
17. A. Kumar and B. L. Eyre, Proc. Roy. Soc. Lond. A370(1980)431
18. B. Tabering and N. Reheis, Intern. J. Refractory Metals and Hard Materials, 28(2010)728
19. F. Morito, J. Less-Common Met. 146(1989)337
20. F. Morito, J. Phys. IV, 1(1993)553
21. F. Morito, J. Mat. Sci. 24(1989)3403
22. F. Morito, J. Nuc. Mat.165(1989)142
23. F. Morito, Proc. 13th Intern. Plansee Seminar, Eds. H. Bildstein and R. EcK, 1(1993)585
24. F. Morito, JOM, 45(1993)54
25. F. Morito, Colloque de Physique, C1-51(1990)281
26. B. V. Zatolokin, Weld. Prod. 26(1979)7
27. Y. V. Milman, Auto. Weld. 33(1980)68
28. A. S. Wronski, A. C. Chilton and E. M. Capron, Acta Met. 17(1969)751
33

34. References-b
1. F. Morito, J. Mat. Sci, 24(1989)3403
2. F. Morito, J. Nuc. Mat. 165(1989)142
3. F. Morito, J. Less Common Met. 146(1989)337
4. F. Morito, Processing and Applications of High Purity Refractory Metals and Alloys, P.
Kumar, H. A. Jehn and M. Uz, eds. Pittsburgh, Pennsylvania, TMS, (1993), p.197-208
5. L. E. Davis, N. C. MacDonald, P. W. Palmberg, G. E. Riech and R. E. Weber, Handbook of
Auger Electron Spectroscopy, 2nd Ed. Physical Electronics Industries Inc. Eden Plairie,
Minnesota, 1976
6. F. Morito, Proc. 13th Intern. Plansee Seminar, H. Bildstein and R. EcK, eds. Reutte/Tirol,
1(1993)585
7. 栗下裕明, 久芳俊一, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)539
8. 栗下裕明, 大石朗, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)546
9. F. Morito, Proc. 12th Intern. Plansee Seminar, Eds. H. Bildstein and H. M. Ortner,
1(1989)313
10. F. Morito, Colloque de Physique, C1-51(1990)281
11. F. Morito, Vacuum, 41(1990)1743
12. F. Morito, JOM, 45(1993)54
13. Y. Hiraoka, F. Morito, M. Okada and R. Watanabe, J. Nuc. Mat. 78(1978)192
14. F. Morito, Proc. 12th International Plansee Seminar, H. Bildstein and H. M. Ortner, eds.
Reutte, Tirol, Austria, 1(1989)417
15. F. Morito, JOM, 45(1993)54
16. F. Morito, High Temperatures and High Pressures, 26(1994)101
17. F. Morito, J. Phys. IV, 1(1993)553
18. N. Igata, A. Kohyama and K. Itadani, J. Nuc. Mat. 85-86(1979)895
19. J. Wadsworth, C. M. Packer, P. M. Chewey and W. C. Coons, Mater. Sci. Eng.
59(1983)257
20. A. S. Wronski, A. C. Chilton and E. M. Capron, Acta Met. 17(1969)751
21. 栗下裕明, 久芳俊一, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)539
22. 栗下裕明, 大石朗, 久保晴義, 吉永日出男, 日本金属学会誌, 7(1983)546
23. A. V. Krajnikov, F. Morito and V. N. Slyunyaev, Intern. J. Refractory Metals and Hard
Materials, 15(1997)325
34

35. 24. H. Kurishita, H. Yoshinaga, K. Abi Ko, S. SuzuKi and H. Kimura, Trans. JIM,
27(1986)739)
25. H. Kurishita, O. To Kunaga and H. Yoshinaga, Mater. Trans. JIM, 31(1990)190
35

36. References-c
1. T.B. Massalski, Eds. Alloy Phase Diagrams, vol. 2, ASM International, Metals ParK, OH,
(1986)
2. G. A. Geach and J. R. Hughes, Proc. 2nd Plansee Seminar, Pergamon Press, London,
(1956)245
3. Von. I. Glass and G. Bohm, Planseeberichte. 10(1962)144
4. M. S. Duesbery, Dislocation in Solids, Elsevier Science Publishers, Netherlands, (1989)69
5. A. D. Korotaev, A. N. Tyumentsev and Yu. I. Pochivalovl, Proc. International Symposium
on Rhenium and Rhenium Alloys, B.D. Bryskin, ed. TMS, (1997)661
6. J. A. Shields Jr, Proc. 1992 Powder Metallurgy World Congress, Metal Powder Federation,
Princeton, NJ, 8(1992)199
7. J. A. Shields Jr, Adv. Mat. Proc. 1421(1992)28
8. R. L. Mannhreim and J. L. Garin, J. Materials Processing Techrnology, 143-144(2003)533
9. F. Morito, Proc.12th Plansee Seminar, Reutte, Tirol, Austria, 1(1989)417
10. F. Morito, Colloque de Physique, C1-51(1990)281
11. F. Morito, Surface and Interface Analysis, 15(1990)427
12. F. Morito, JOM, 45(1993)54
13. F. Morito, Proc. 13th Plansee Seminar, Reutte, Tirol, Austria, 1(1993)585
14. F. Morito, Proc. International Symposium on Rhenium and Rhenium Alloys, B.D.
Bryskin, ed. TMS, (1997)559
15. F. Morito, Proc. 14th Plansee Seminar, Reutte, Tirol, Austria, 1(1997)103
16. J. C. Carlen and B. D. Bryskin, J. Materials Engineering and Performance, 3(1994)282
17. I. Wesemann, A. Hoffmann, T. Mrotzek and U. Martin, Int. J. Ref. Met. Hard Mat.
28(2010)709
18. R. L. Fleischer, Acta Metall. 9 (1963) 203-209
19. H. Suzuki, Sci. Rep. Res. Inst. Tohoku Univ. A4 (1952) 455-463
20. R. A. Labusch, Phys. Status Solidi, B-41 (1970) 659-669
21. J. Xu, E. A. Kenik and T. Zhai, Phil. Mag. 88(2008)1543
22. J. Xu, E. A. Kenik and T. Zhai, Mater. Sci. Eng, A(2008)76
36

37. References-d
1. N. N. Morgunova, B. A. Klypin et al. Molybdenum Alloys, Moscow, Metallurgy, 1975
2. T. B. Massalski, Eds. Alloy Phase Diagrams, vol. 2, ASM International, Metals Park, OH,
1986
5. 3. R. A. Erck and L. E. Rehn, Phil. Mag. A62(1990)29
6. 4. A. Hasegawa, K. Ueda, M. Satou and K. Abe, J. Nuc. Mat. 258-263(1998)902
7. 5. Y. Nemoto, A. Hasegawa, M. Satou, K. Abe and Y. Hiraoka, J. Nuc. Mat. 324(2004)62
8. 6. Sh. Sh. Ibragimov, V. V. Kirsanov and Yu. S. Pjatiletov, Radiation Damage of Metals
and Alloys, Moscow, Energoatomizdat, (1995)210
9. 7. K. J. Leonard, J. T. Busby and S. J. Zinkle, J. Nuc. Mat. 366(2007)336
10. 8. K. J. Leonard, J. T. Busby and S. J. Zinkle, J. Nuc. Mat. 366(2007)369
11. 9. J. T. Busby, K. J. Leonard and S. J. Zinkle, J. Nuc. Mat. 366(2007)388
12. 10. E. J. Edwards, F. A. Garner and D. S. Gelles, J. Nuc. Mat. 375(2008)370
13. 11. F. Morito and K. Shiraishi, J. Nuc. Mat. 179-181(1991)592
14. 12. V. P. Chakin, F. Morito, V. A. Kaza Kov, Yu. D. Goncharenko and Z. E. Ostrovsky, J.
Nuc. Mat. 258-263(1998)883
15. 13. F. Morito, V. P. Chakin, Yu. D. Goncharenko, A. O. Posevin and L. A. Yevseev,
Microstructure evolution during high-temperature annealing in molybdenum-rhenium
alloys welds irradiated at low temperature, presented at ICFRM-11, Kyoto, Dec. 2003
16. 14. F. Morito, V. P. Chakin, H. Saito, N. I. Danylenko and A. V. Krajnikov, Proc. 17th
International Plansee Seminar, P. Rhoedhammer et al. Eds. Reutte/Tirol, vol. 1, (2009),
RM63
17. 15. F. Morito, V. P. Chakin, M. I. Danylenko and A. V. Krajnikov, J. Nuc. Mat. (2011),
doi;10.1016/j.jnucmat.2010.12.183
18. 16. F. Morito, Proc. 12th International Plansee Seminar, G. Kneringer et al. Eds.
Reutte/Tirol, 1(1989)417
19. 17. F. Morito, Proc. 13th International Plansee Seminar, G. Kneringer et al. Eds.
Reutte/Tirol, 1 (1993)585
20. 18. F. Morito, Proc. International Symposium on Rhenium and Rhenium Alloys, TMS,
(1997)559
21. 19. F. Morito, Proc. 14th International Plansee Seminar, G. Kneringer et al. Eds.
Reutte/Tirol, 1(1997)1037
37

38. 22. 20. V. A. Kazakov, A. S. Pokrovsky, A. V. Smirnov and L. I. Smirnova, Physics of
Metals and Metal Science, 42(1976)357
23. 21. V. A. Kazakov, A. S. Pokrovsky and A. V. Smirnov, Nuclear Technology,
53(1981)392
38

39. References <JNM-2010> 8
[1] S.R. Agnew and T. Leonhardt, JOM, 55(10), (2003) 25-29.
[2] R.W. Buckman, Jr., Proc. Nuclear Space Power Systems Materials Requirements, AIP,
699, (2004) 816-821.
[3] V.P. Chakin, Yu.A. Vlasov, V.A. Kazakov, Yu.D. Goncharenko and A.O. Posevin,
Prospective Materials, No.5, (2005) 50-56.
[4] M.S. El-Genk and J.-M. Tournier, Journal of Nuclear Materials, 340, (2005) 93-112.
[5] D.J. Olson, B. Mishra and D.W. Wenman, Mineral Processing and Extractive Metallurgy
Review, 22, (2001) 1-23.
[6] F. Morito and K. Shiraishi, Journal of Nuclear Materials, 179, (1991) 592-595.
[7] V.P. Chakin, F. Morito, V.A. Kazakov, Yu.D. Goncharenko and Z.E. Ostrovsky, Journal
of Nuclear Materials, 258-263 (1998) 883-888.
[8] F. Morito, V.P. Chakin, H. Saito, N.I. Danylenko and A.V. Krajnikov, Proc. 17th
International Plansee Seminar, P. Rhoedhammer et al. (Eds.), Reutte/Tirol, vol. 1, (2009)
RM63.
[9] F. Morito, Proc. 12th International Plansee Seminar, G. Kneringer et al. (Eds.),
Reutte/Tirol, vol. 1, (1989) 417-431.
[10] F. Morito, JOM, 45(6), (1993) 54-58.
[11] F. Morito, Proc. 13th International Plansee Seminar, G. Kneringer et al. (Eds.),
Reutte/Tirol, vol. 1, (1993) 585–598.
[12] A.V. Krajnikov, F. Morito and V.N. Slyunyaev, International Journal of Refractory
Metals and Hard Materials, 15, (1997) 305-339.
[13] F. Morito, T. Noda and A.V. Krajnikov, Proceedings International Conference on
Science for Materials in the Frontier of the Centuries: Advantages and Challenges, IPMS
NASU, Kiev, VI (2002) 140.
[14] H. Saito and F. Morito, Basics of Autoradiography and Its Application in Materials
Science and Engineering, H. Saito (Eds.), Urabe Publishing, Chiba, (2008) pp. 101-112.
[15] F. Morito, Journal de Physique, IV, 1, (1993) 553-556.
[16] F. Morito, Proc. 14th International Plansee Seminar, G. Kneringer et al. (Eds.),
Reutte/Tirol, vol. 1, (1997) 1037-1049.
[17] F. Morito, Proc. International Symposium on Rhenium and Rhenium Alloys, (1997) 559-
568.
[18] N.N. Morgunova, B.A. Klypin, et al., Molybdenum Alloys, Moscow, Metallurgy, 1975.
[19] T.B. Massalski, Eds., Alloy Phase Diagrams, vol. 2, ASM International, Metals Park,
OH, 1986.
[20] R.A. Erck and L.E. Rehn, Philosophical Magazine A, 62, (1990) 29-51.
[21] A. Hasegawa, K. Ueda, M. Satou and K. Abe, Journal of Nuclear Materials, 258-263,
(1998) 902-906.
[22] Y. Nemoto, A. Hasegawa, M. Satou and K. Abe, Journal of Nuclear Materials, 324,
39

40. (2004) 62-70.
[23] Sh.Sh. Ibragimov, V.V. Kirsanov and Yu.S. Pjatiletov, Radiation Damage of Metals and
Alloys, Moscow, Energoatomizdat, (1995) 210-213.
[24] K.J. Leonard, J.T. Busby and S.J. Zinkle, Journal of Nuclear Materials, 366, (2007) 369-
387
[25] E.J. Edwards, F.A. Garner and D.S. Gelles, Journal of Nuclear Materials, 375, (2008)
370-381.
40