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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]
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.
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
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 (%) :
"ominal 0 2 4 10 13 15 20 25 30 40 41 47 50
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
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 (◇).
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)
300 400 5000
10
20
30
40
50
0
200
400
600
800
E, %
T test
, K
Ef , Unirrdiated
Euni
, Unirrdiated
Ef , Irradiated
Euni
, Irradiated
YS
, M
Pa
YS, Unirrdiated YS, Irradiated
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.
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>
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)
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>
0 10 20 30 400
10
20
Unirradiated, SR Unirradiated, Rec Irradiated, SR Irradiated, Rec
Elo
nga
tion, %
Re content, %
T=300K
Fig. 7. Fracture elongation tested at 300 K in Mo-Re welds. (Open mark=Recrystallized (Rec), Closed mark=Stress-relieved (SR). : unirradiated, ----- : irradiated).
10
0 10 20 30 400
10
20
30
Elo
nga
tion, %
Re content, %
Unirradiated, SR Unirradiated, Rec Irradiated, SR Irradiated, Rec
T=1023K/1073K
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).
11
(a)
0 10 20 30 40 500
200
400
600
800
1000
1200
1400
1600
Te
nsile
str
en
gth
, M
Pa
Re content, %
Unirradiated, SR
Unirradiated, Rec
Irradiated, SR
Irradiated, Rec
(b)
0 10 20 30 40 500
200
400
600
800
1000
1200
1400
1600
Ten
sile
str
eng
th, M
Pa
Re content, %
Unirradiated, SR
Unirradiated, Rec
Irradiated, SR
Irradiated, Rec
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).
0 10 20 30 40 500
200
400
600
800
1000
1200
1400
Mic
roh
ard
ne
ss
Re content, %
Unirrad., SR, BM
Unirrad., Rec, BM
Irrad., SR, WM
Irrad., SR, HAZ
Irrad., SR, BM
Irrad., Rec, WM
Irrad., Rec, HAZ
Irrad., Rec, BM
Fig. 10. Hardness at 300 K in Mo-Re welds. (Open mark=Recrystallized (Rec), Closed mark=Stress-relieved (SR). : unirradiated, ----- : irradiated).
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)).
13
30Re 50Re
BM
(a) (d)
HAZ (b) (e)
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
40µ
m
14
(a) (b) (c)
(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.
40 µm
15
a b
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
4 µm 4 µm
4 µm 4 µm
Particles of
second phase
16
a
b
c
Fig. 16. Metallography of Mo-20Re weld; 1673 K, 1h: a) BM b) HAZ c) WM
40 µm
40 µm
40 µm
8 µm
8 µm
8 µm
17
a
b
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
90 nm
90 nm
150 nm
18
0 10 20 30 40 50
0
200
400
600
800
1000
Hv
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
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.
20
IV. Irradiated Mo-Re welds at LT 6, 7 <LT irradiation>
0 10 20 30 40
0
10
20
30
0
10
20
30
E f, %
YR
XT B C R D
Legend
Line
0 10 20 30 400
200
400
600
800
1000
YR
XT
B C R D
Text
σY
S , M
Pa YL
Re content, % XB
Arrow
Arrow1
Arrow2 Arrow3
Irradiated Text1
Unirradiated Text2
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.
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.
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
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.
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.
1015
1016
1017
4
6
8
10
12
1111
2222
3333 4444
5555
Loop Size / nm
Loop Density / cm-3
30Re30Re30Re30Re 20Re20Re20Re20Re
41Re41Re41Re41Re
BMBMBMBM WMWMWMWM BMBMBMBM
BMBMBMBM
WMWMWMWM
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. χχχχ-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 χχχχ-phase along grain boundaries [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
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.
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.
28
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