Re effect mo re welds-2

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Re effect mo re welds-2

  1. 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 strengtheningAbstract: << MoRe_Report_Dec1_2010 by Sasha>>This work analyses several Mo-Re alloys and welds with the Re content 0-50% in as-receivedstate and after electron beam welding and/or radiation treatment. The mechanical propertiesand microstructure of Mo-Re welds are examined focusing on the effect of Re concentration.Phase stability, microstructural changes and impurity redistribution are studied for betterunderstanding and predicting the long-term performance of Mo-Re alloys at hightemperatures and/or high neutron fluences.In particular, a strategy of welding of Mo-Re alloys is discussed with emphasis on thesensitivity 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 refinementduring directional solidification after welding, ductility improvement and fracture modechange from intergranular to transgranular one are clearly observed with an increase of Recontent.Effect of neutron irradiation on the strength of Mo-Re welds is studied for a wide temperaturerange. Mo-Re welds exhibit a large radiation-induced strengthening. At room temperature, thestrengthening effect is rather limited and unstable because of lack of ductility. Thestrengthening becomes strongly pronounced at high temperatures. Damaging effect ofneutrons 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-Realloys 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 simulatethermal conditions in different welding zones. Impurity redistribution in Mo-Re alloys hasbeen studied by surface analysis methods. The role of carbon and oxygen segregation as wellas formation of carbides and Re-base phases is discussed to minimise intergranularembrittlement of welds. [14, Sasha] 1
  2. 2. IntroductionRhenium is widely used to alloys VIA group refractory metals. In particular, the strength andplasticity, creep resistance, and low temperature ductility of Mo are all improved withincreasing 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 alloys1-1. Solution softening and hardening in Mo-Re alloysFig. 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 at1923 K, 1 h, respectively. ( ) means brittle fracture before yielding. 2
  3. 3. Fig. 2 shows yield stress by bend test vs test temperature in Mo-Re welds. Mo-41Re weldshad much higher yield stress than those of Mo-5Re and Mo-13Re, which failed in a brittlemanner at 77 K. But Mo-41Re welds showed more ductile and farctured after yielding even at77 K.Fig. 2 also shows a transition of solution softening by a test temperature. At 300 K, Mo-5wt.% Re exhibited more solution softening than Mo-13 wt.% Re. However minimum yieldstress moved to Mo-13 wt.% Re from Mo-5 wt.% Re with a decrease of test temperature.2. Previous report on Re effectIn the VI A group metals, the conditions of a resonance covalent bond are fulfilled in the bestway. 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 mechanicalproperties specific for these metals [1]. (Milman et al)Rhenium Effect by Korotaev et al [2-4]1. Significant enhancement of low temperature plasticity2. Reduction of the temperature of the viscous-brittle transition, Tv.b.3. Suppression of brittle sliding fracture4. Upgrading of the weldability of high rhenium alloysVery interestiing characteristic changes in strength properties such as5. Solution softening in low-rhenium alloys6. 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 rolling8. Extraordinaly pronounced effects of solution hardening in the region T > 0.2 TmeltA a result, the phenomenology of the enhancement of the strength and plastic properties andof 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 Effectand 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)717-7282. 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. 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)671-6804. 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)707-7163. 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 corrosionresistance are the key issues for these application fields. Tendency of Mo alloys to embrittle atlow temperatures assumes probable degradation of the mechanical properties either duringwelding [5-9] or under irradiation [10-13]. In spite of the fact that significant progress hasbeen achieved in studying the mechanism of rhenium effect, many details of Mo-Re alloybehaviour 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. 5. 4. Irradiation effect on Mo alloys and welds4-1. Welds of Mo, TZM and Mo-0.56%Nb irradiated to 1017 cm–2 ~ 1020 cm–2 (E> 1 MeV) at348-1073 K6 (irradiation by JRR-2 and JRR-4 at JAERI).[5]I. Tensile propertiesMo 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) Totalelongation in as-welded PM Mo (▽), postweld annealed PM Mo (○), postweld carburizedPM Mo (□), postweld annealed TZM (△) and postweld annealed Mo-0.56 wt.% Nb (◇). 5
  6. 6. Fig. 4. Yield stress (△), tensile stress (○) and total elongation (□) of (a) as-welded PM Moand (b) postweld annealed TZM irradiated to 1.2 x 1024 m-2 at 1073 K. Closed marks denoteto 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 , KFig. 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. 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, 181KMo-0.56%Nb: as-welded <Ef~>2% limited to WM and HAZ rather ductile> 235KMo-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) 443KMo-0.56%Nb: postweld annealed & irradiated, 1073K, 6.0 × 1019 cm–2 (E> 1 MeV) 443KEmbrittlement 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. 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) 443KEmbrittlement due to irradiation was not so significant at RT in the case of Mo-0.56%Nb !!Mo-0.56%NbNeutron 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 andMo-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 byelectron-beam welding and further by neutron irradiation.(2) Addition of carbon was effective to suppress intergranular embrittlement which wascaused by electron-beam welding and neutron irradiation. It is considered that segregation ofcarbon and precipitation of carbides enhanced the intergranular cohesion so that considerablestrength and ductility was maintained.(3) TZM showed higher strength at high temperatures, but it seemed unavoidable to improvelower 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 underthe irradiation to 6.0 x 1023 n/m2 at 1073 K. It is considered that good mechanical propertiesof 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 1022n/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. 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. 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. 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 forirradiated 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), Closedmark=Stress-relieved (SR). : unirradiated, ----- : irradiated). 11
  12. 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. WMFig. 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 networksand dislocation tangles with rather low density are seen (Fig. 12-11 (a ) and (b)). 12
  13. 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. 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. 15. 4 µm 4 µm a b Particles of second phase 4 µm 4 µm c dFig. 15. Fracture surface of Мо-Re welds irradiated at Тirr = 1023-1073 K after tensile test atroom 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. 16. 8 µm 40 µm a 8 µm 40 µm b 8 µm 40 µm cFig. 16. Metallography of Mo-20Re weld; 1673 K, 1h: a) BM b) HAZ c) WM 16
  17. 17. 90 nm a 90 nm b 150 nm cFig. 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. 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 ○ - irradiated3. Results <HT irradiation>The tensile strength of both irradiated and unirradiated welds is shown in Fig. 1a and Fig. 1bas 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 weldsdemonstrated adequate ductility with elongation ~20-30% at high temperatures. Althoughsome variations of strength data were recognized in irradiated samples, tensile strengthgenerally increased with Re content. After neutron irradiation, Mo-Re welds showed a largeradiation-induced strengthening. At room temperature, the strengthening effect was ratherlimited and unstable because of lack of ductility. It appeared in some samples with residualductility but was absent in absolutely brittle samples. The strengthening became stronglypronounced at high temperatures. In particular, tensile strength of all irradiated welds at1073 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 MPand from 420 to 1250 MPa respectively. The radiation-induced strengthening at hightemperature 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 allzones 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. 19. between different welded zones was very small, if any. Each zone of all irradiated weldsexhibited a severe radiation-induced hardening compared to that of unirradiated samples. Anincrease of the microhardness resulted from high temperature irradiation was approximately300 MPa for SR welds and varied between 400 and 600 MPa for Rec samples depending onthe Re content. Fig. 3 shows an increase of density (∆d/ dini) in Mo-Re welds after hightemperature 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-Rewelds with 16-21% Re. Further increase of Re content monotonously weakens the observedeffect of density growth. As shown in Fig. 4, massive secondary phases were recognized asthe form of thin plane particles with 70-160 nm length in all the Mo-Re welds. It indicatesthat microstructure reconstruction of irradiated Mo-Re welds is connected with the formationof secondary phases of much higher density compared with that of unirradiated specimen. 19
  20. 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 annealedat 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. 21. (a) (b) (c)Fig. 20. Macrostructure of tensile specimen fractured at WM. (a) Mo-15Re fractured duringdisassembling of capsule, (b) Mo-20Re fractured during disassembling of capsule, (c) Mo-30Re fractured by tensile test at RT. 21
  22. 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 tangleswith rather low density are seen (Fig. 12-11 (a ) and (b)). Dislocation density in WM was lessthan in the BM. So, the dislocation density in BM was 2.9 x 1010 cm-2 and that in WM was 1.9x 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 suchlow-temperature irradiation. Their average size is 7.5 ~ 10 nm. Density of dislocation loops 22
  23. 23. varied in the range from 4.5 ´ 1015 to 2.6 ´ 1016 cm-3 and decreased with an increase ofrhenium content.The structure of all irradiated Mo-Re alloys is retained as one phase. The second phaseprecipitations, which were remarkably observed in Mo-Re welds irradiated at highertemperatures [1], were not detected.Dislocation loop distribution in BM of Mo-41Re welds was extremely irregular, butdislocation loops were almost absent in WM as in Fig. 12-13 (e), (f). The reason is to besolved 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-3Fig. 22. Dislocation loop density vs size in Mo-Re welds irradiated at LT.5. Conclusions1. 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 thespecimens centre through the weld-fusion zone. With increasing rhenium content the fracturetype changed from the brittle intergranular type to the transgranular one. 23
  24. 24. 2. It was shown that with increasing rhenium contentthe dislocation loop density with anaverage size of 7.5~10 nm was reduced 4~6 times and in the fusion zone of Mo-41Re weldswere quite absent.V. Summary <HT irradiation>& <LT irradiation>Radiation-Induced Segregation, Precipitation, Hardening, Embrittlement, Transmutation11. RIPNucleation sites of σ-phase are considered to be subgrain boundaries consisting of tangleddislocations, 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 becontrolled by thermal treatment.On the other hand, the mean length and the number of density of χ-phase precipitates were notchanged with thermal treatment. This suggests that the nucleation site of χ-phase aredislocations 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)865After nucleation of these σ-phase and χ-phase precipitates, the surface of the precipitates actedas 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 primephase precipitates in Ni-Al alloy -> Fig. 1112. RIH and RIE by Orowan model126-128 -> Fig.12: Re content dependence of radiationhardening calculated from results of the microstructural observation. Calculated Hv agreedwith the measurement in the specimens irradiated at 1072 K. In the specimens irradiated at681 or 874 K, the scatter of the calculation became larger. It is considered because therewould be invisible small defects at lower temperature irradiation.12*. RIEFor Mo-41Re irradiated at 874 K or below, there were cracks observed around the indentationafter Hv measurements. This embrittlement is thought to be caused by large and hard σ-phaseprecipitates. σ-phase has very high HV value of about 1500 (MPa), which is much larger thanthat of the matrix, and hence the surface or inside of such hard precipitates would be theinitiation sites of cracks. Thus the formation of large σ-phase ppts led to drastic embrittlementin 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 wasrecently shown to form χ-phase along grain boundaries [24]. 24
  25. 25. [24] K.J. Leonard, J.T. Busby and S.J. Zinkle, Journal of Nuclear Materials, 366, (2007) 369-387Microstructural and mechanical property changes with aging of Mo–41Re and Mo–47.5Re alloysThe changes in microstructure and mechanical properties of Mo–41Re and Mo–47.5Re alloyswere investigated following 1100 h thermal aging at 1098, 1248 and 1398 K. The electricalresistivity, hardness and tensile properties of the alloys were measured both before and afteraging, along with the alloy microstructures though investigation by optical and electronmicroscopy 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 alongthe grain boundaries. This resulted in increases in resistivity, hardness and tensile strengthwith a corresponding reduction in ductility.(4) Aging Mo–47.5Re at 1098 and 1248 K led to the development of the χ phase along grainboundaries, resulting in decreased resistivity and increased hardness and tensile strengthwhile 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 andnon-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 dpaproduced a high density of thin platelets of a hexagonal close-packed (hcp) phase identifiedas 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 precipitatesare 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> 9In accordance with the phase diagrams [18, 19], formation of σ-phases and χ-phases takesplace at high temperatures in Mo-Re alloys with a high Re content. Formation of secondaryphases under irradiation in solid solution is possible due to radiation-induced segregation ofRe atoms followed by Re enrichment of internal sinks up to solubility limit or even exceedingit. 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. 26. properties of Mo-50Re welds by controlling dispersion and size of σ-phase along grainboundaries and in the matrix [16, 17]. The results obtained here also correspond well with theprevious ones [20-22]. As was shown here, formation of secondary phases is significant andintensive in Mo-Re welds. Character and rate of phase formation are almost independent onthe Re content as well as on the location within the welded zone. Secondary phases look likethin plane particles located along some crystallographic planes as shown by TEM. Based onthe particle shape, one may conclude that they are σ-phases rather than χ-phases which areusually close to spheroids [20]. High intensity of secondary phase formation, homogeneousdistribution of the particles over the bulk, and absence of depleted zones along grainboundaries indicate that the phases were formed by means of mechanism of homogeneousnucleation which is based on the radiation-induced separation of atoms in alloys [23]. In thiscase, 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 correlatewith structure defects such as voids, dislocations, grain boundaries and so on. Susceptibilityof Mo and Re atoms to separation under irradiation is high and stability of the formed nucleiis 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 veryhigh intensity of radiation-induced separation of Mo and Re atoms in Mo-Re welds underneutron irradiation. This eliminates effect of alloy composition on the final microstructure andleads to formation of practically the same number of secondary phases in the welds in spite ofRe content. Ageing at 1098 and 1248 K of two-phase Mo-47.5 wt% Re with αMo + σstructure was recently shown to form χ-phase along grain boundaries [24]. Mo–41 wt% Reirradiated in the fast flux test facility (FFTF) experienced significant and non-monotonicchanges in density due to radiation-induced segregation, leading to non-equilibrium phaseseparation, and progressive transmutation of Re to Os [25]. Beginning as a single-phase solidsolution 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. Thesehcp 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 thehcp phase, but the precipitates are much bigger and more equiaxed in shape. Although notformed 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 forrhenium.After low temperature irradiation, fracture surface is mainly located in WM, because this zoneusually 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 Recontent and under low temperature irradiation.In contrast to low temperature irradiation, all zones of welds after high temperature irradiationare characterized by comparable strength level. Thus, fracture at room temperature occurredat any place of the welds by brittle transgranular cleavage.Numerous secondary phases were observed at cleavage lines and within matrix. In additionthe 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. 27. Intensive phase formation in Mo-Re welds eliminated any difference in mechanical propertiesbetween WM, HAZ and BM. Mo-Re welds with such a characteristic microstructure still keptsufficient plasticity at high temperatures. Therefore it is clear that high temperature neutronirradiation of Mo-Re welds have lesser damaging effect than low temperature irradiation.As a result, we obtained a very important practical conclusion that weld-fabricatedconstructions operating under high temperature radiation do not limit their life becausemechanical properties of welds with WM and HAZ are not worse than those of BM.Concluding remarks1. 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
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