annual report of the association euratom/cea 2000annual report of the association euratom/cea 2000...

Annual Report of the Association EURATOM/CEA

2000

Compiled by : Ph. MAGAUD and F. Le VAGUERES

ASSOCIATION CEA/EURATOM DSM/DRFC

CEA/CADARACHE 13108 Saint-Paul-Lez-Durance (France)

- I -

CONTENTS

INTRODUCTION ................................................................................................................................. 1

EFDA TECHNOLOGY PROGRAMME ........................................................................ 3

Physics Integration

Diagnostics TW0-DIAG-DEV ITER diagnostic window development ........................................................................ 5

Heating and Current Drive TW0-ICRF/ANT ICRF Antenna and vacuum transmission line development - ICRH Antenna coupling: Near field computations....................................................... 7 CEFDA 00-519 Neutral beam heating design for EDA extension ......................................................... 11 TW0-NB.DEV.1 Neutral beam development for EDA extension - EU-JA collaborative experiment on KAMABOKO source.......................................... 15 TW0-NB.DEV.3 Neutral beam development for EDA extension - 1 MeV SINGAP accelerator ......................................................................................... 19 TW0-NB.DEV.4 Neutral beam development for EDA extension - Manufacture and testing of a prototype 1 MeV bushing for SINGAP ......................... 23

Vessel/In Vessel

Plasma Facing Components CEFDA 99-501 Critical heat flux and thermal fatigue testing of CFC monoblocks - 200 kW electron beam gun test .................................................................................... 27 CNET 98-480 Thermal fatigue testing of divertor full scale prototype - 200 kW electron beam gun test .................................................................................... 29 CNET 98-485 Thermal fatigue testing of baffle full scale prototype - 200 kW electron beam gun test .................................................................................... 35 DV4.3 Optimisation and manufacture of HHF components - Study of flat tile cascade failure possibility for high heat flux components ................. 37 TW0-DV4/01 Optimisation of manufacture of HHF components ...................................................... 41 TW0-DV4/02 Optimisation of Cu-SS-Tube Joints.............................................................................. 43 TW0-T438-01 Development and testing of time resolved erosion detecting techniques ..................... 45

- II -

Vessel/Mechanical structure T216-GB8 Small scale testing of FW/BS modules ........................................................................ 49 TW0-LASER/CUT VV intersector maintenance - Further development of high power Nd YAG - Improvement of YAG laser back-plate cutting using powder injection ....................... 51 TW0-LASER/HYD YAG laser process for cutting and welding of the hydraulic module blanket - CW YAG laser process for cutting and welding of the hydraulic connections ............ 53 TW0-LASER/REWELD VV intersector maintenance - Further development of high power Nd YAG - Laser rewelding after cutting........................................................................................ 55 TW0-LASER/WELD VV intersector joining - Further development of high power Nd YAG - Laser welding with multipass filler wire ...................................................................... 59 TW0-T301/01 Ultrasonic inspection for vessel inter-sector weld - Completion of ultrasonic inspection process development - Optimisation of the ultrasonic inspection method : effect of the wave polarity and frequency on the signal to noise ratio .................................................................... 61 TW0-T420/06 Fabrication of a first wall panel with HIP’ed beryllium armor .................................... 65 TW0-T420/08 Development of HIP fabrication technique - Faisability of using HIP for the manufacturing of Primary First Wall Panel (PFWP) ........................................ 69 TW0-T427/03 Development and characterization of one-step SS/SS and CuCrZr/SS HIP joints including mock-up fabrication............................................................................ 73 TW0-T508/04 Development of Be/CuCrZr HIPping technique .......................................................... 75 TW0-T508/05 Development of Be/CuCrZr brazing techniques .......................................................... 77 TW0-T519/01 Low temperature water radiolysis ................................................................................ 79 V63 Detailed investigation of CuAl25 (IG1) - Its joints with 316LN stainless steel and joints testing procedure.......................................................................................... 81

Assembly & Maintenance T252 Radiation tolerance assessment of standard components for remote handling and process instrumentation .......................................................................... 85 T329-5 In vessel RH dexterous operations - Task extension 2 ................................................. 91 TW0-DTP/1.1 Carrier and bore tools for 4" bend pipes....................................................................... 95 T329-4 TW0-DTP/1.2 Prototypical manipulator for access through IVVS...................................................... 99 TW0-DTP/1.4

Magnet Structure CEFDA 00-517 Design of cryogenic transfer lines and ring manifolds for ITER-FEAT ...................... 103 CEFDA 00-541 PF correction coils : conceptual design and analysis, CS cooling inlets, conductor analysis tools ............................................................................................... 105 M40 Design work on magnet R&D ...................................................................................... 109

- III -

M50 Conductor R&D - Development of NbTi conductors for ITER PF coils ..................... 111 TW0-T400-1 CSMC and TFMC installation and test ........................................................................ 113

Tritium Breeding and Materials

Breeding Blanket Water Cooled Lithium Lead (WCLL) Blanket TTBA-1.1 TBM Adaptation to Next Step...................................................................................... 117 TTBA-2.1 Blanket Manufacturing Techniques - Definition of specifications for demonstrators ............................................................. 121 TTBA-2.2 Blanket Manufacturing Techniques - Solid HIP demonstrator for fabrication and coating, fabrication of double wall tubes ................................................................................... 125 TTBA-2.4 Blanket Manufacturing Techniques - Feasibility of powder HIP fabrication for TBM or TBM components ......................... 129 TTBA-2.5 Blanket Manufacturing Techniques - Double-wall tube out-of-pile testing (DIADEMO experimental program) .................. 133 TTBA-3.1 Coating qualification and irradiation tests - Permeation barrier qualification - Fabrication and characterization of CVD specimens ................................................... 137 TTBA-3.5 Coating qualification and irradiation tests - Permeation out-of-pile testing ................. 141 WPA4.2.3 TTBA-4.1 Processes and components - Tritium extraction from Pb-17Li .................................... 143 TTBA-5.3 Safety and Licensing - Modelling of the lithium/lead water interaction ...................... 147 TTBA-5.4 Safety and Licensing - Development of minor components & instrumentation........... 149 TTBA-6.2 MHD Effects - Test and modeling of natural convection MHD .................................. 153 Helium Cooled Pebble Bed (HCPB) Blanket TTBB-2.3 Blanket Manufacturing Techniques - First wall manufacturing by HIP forming technique .................................................... 157 TTBB-5.D3 Development of ceramic breeder pebble beds - Characterization of Li2TiO3 pebble beds...................................................................... 159 TTBB-5.D4 Development of ceramic breeder pebble beds - Validation of Li2TiO3 fabrication with pre-industrial means of the lab fabrication steps - Mastering/Optimizing ..................................................... 163

Structural Materials development Advanced materials TTMA-1.2 SiC/SiC ceramic composites - SiC/SiC composite development and characterization................................................. 167 TTMA-1.4 SiC/SiC ceramic composites - Compatibility of SiC/SiC composite with liquid Pb-17Li alloy................................... 169

- IV -

TTMA-1.6 SiC/SiC ceramic composites - Brazing of SiC/SiC composites with thermo-mechanical testing................................. 173 Reduced Activation Ferritic Martensitic (RAFM) Steels SM-2-3-1 Mechanical properties of F82H weldments .................................................................. 177 SM-3-3 Corrosion in water conditions - Corrosion water with additives .................................. 181 SM-3-5.1 Compatibility with hydrogen and liquid - SCC behaviour of Eurofer in aqueous environments ................................................... 185 SM-4-3 F82H weldability with filler metal ............................................................................... 187 SM-5-7 Small size specimen technology................................................................................... 191 SM-6-1 Development of high dose high temperature - Low temperature, high dose irradiation test with ISTC................................................ 193 TTMS-1.1 RAFM steels - Irradiation performance - Neutron irradiation of RAFM steels to 30-35 dpa at 325°C......................................... 197 TTMS-2.1 RAFM steels - Metallurgical and mechanical characterization - SM-2-1 Metallurgical characterization of EUROFER 97 steel ................................................. 199 TTMS-2.2 RAFM steels - Metallurgical and mechanical characterization - Mechanical characterization of Eurofer 97 : tensile and impact properties.................. 203 TTMS-2.3.1 RAFM steels - Metallurgical and mechanical characterization - Mechanical properties of diffusion bonded welds (RAFM/RAFM HIP joint) ............. 207 TTMS-2.3.2 RAFM steels - Metallurgical and mechanical characterization - Mechanical properties of weldments ............................................................................ 211 TTMS-2.4 RAFM steels - Metallurgical and mechanical characterization - Mechanical properties of powder HIP Eurofer............................................................. 215 TTMS-4.1 RAFM steels - Qualification fabrication processes - Eurofer powder HIP development and characterization............................................... 217 TTMS-4.2 RAFM steels - Qualification fabrication processes - Mechanical alloying including ODS ............................................................................ 221 TTMS-4.3 RAFM steels - Qualification fabrication processes - RAFM weldability..................... 225 TTMS-4.4 RAFM steels - Qualification fabrication processes - Dissimilar welding RAFM/316LN............................................................................... 229 TTMS-4.5 RAFM steels - Qualification fabrication processes - Solid HIP process and qualification ............................................................................. 231 TTMS-5.1 RAFM steels - Rules for design and inspection - Design code assessment and development ................................................................... 235 TTMS-5.4 RAFM steels - Rules for design and inspection - Data collection and data base maintenance .................................................................. 239 TTMS-6.3.1 RAFM steels - Qualification of high performance steels - Development of RAFM-ODS steels............................................................................. 243 Neutron Source TTMI-001 IFMIF - Accelerator facility ......................................................................................... 245

- V -

Safety and Environment CEFDA 00-539 French contribution to ITER licensing working group................................................. 249 SEA1-14 Assist JCT on safety documentation ............................................................................ 251 SEA4-4 Plant safety assessment - ITER-FEAT design CAT V accident analysis : H2 production via delayed plasma shutdown................................................ 253 SEA5-2 Validation of computer codes and models - ISAS maintenance................................... 255 SEA5-3 Validation of computer codes and models - Thermohydraulical code validate experiments ............................................................. 257 TSW-2.1 Waste and Decommissioning Strategy - Detritiation of structural material and management of tritiated waste.......................... 259 TSW-2.6 Waste and Decommissioning Strategy - Development and assessment of overall strategies....................................................... 263 TW0-SEA3.5 Hydrogen deflagration/detonation analysis .................................................................. 265 TW0-SEA4 Plant safety assessment................................................................................................. 267

System Studies

Power Plant Conceptual Studies (PPCS) TW0-TRP-1.1 Safety and environmental requirements ....................................................................... 271 TW0-TRP-2.4 Economic and operational requirements ...................................................................... 273 TW0-TRP-2.5 Supply requirements..................................................................................................... 277 TW0-TRP-3D4 Reactor integration ....................................................................................................... 279 TW0-TRP-4D5 High heat flux components........................................................................................... 283 TW0-TRP-4D7 Economics and operational analyses - Maintenance schemes...................................... 285

Socio-Economics Studies SERF2 Externalities of fusion - Accident external costs and sensitivity analysis on the site location............................ 289

UNDERLYING TECHNOLOGY PROGRAMME ................................................ 293

Vessel/In Vessel

Plasma Facing Components UT-PFC&C-HFW Transparent polycrystalline ceramics ........................................................................... 295 UT-PFC&C-HIP Mechanical behaviour of HIP joints ............................................................................. 299

- VI -

UT-PFC&C-SiC C-SiC and boron doped graphite .................................................................................. 303 UT-PFC&C-SiC/MJ Development of SiC/metal joining technology ............................................................ 307 UT-SM&C-VVI Design and analysis of vacuum vessel and internals .................................................... 309

Vessel/ Mechanical Structure UT-SM&C-BIM Dissimilar diffusion-bonded joints mechanical testing ................................................ 313

Assembly and Maintenance UT-RH1 Remote Handling Techniques - Technology and control for remote handling systems.................................................. 317 UT-RH2 Remote Handling Techniques - Graphical programming for remote handling ............................................................... 319 UT-RH3 Remote Handling Techniques - Advanced technologies for high performances actuators ............................................. 323 UT-RH4 Remote Handling Techniques - Radiation tolerance assessment of electronic components from specific industrial technologies for remote handling and process instrumentation.................... 327

Tritium Breeding and Materials

Breeding Blanket UT-SM&C-BLK Development of breeding blankets ............................................................................... 333 UT-SM&C-COAT Development of an oxide layer on Fe-Al coatings ....................................................... 337 UT-SM&C-LM/DC Diffusion coefficient measurements in liquid metals ................................................... 341 UT-SM&C-LM/MAG Liquid metal corrosion under magnetic field................................................................ 345 UT-SM&C-LM/Refrac Compatibility of refractory materials with liquid alloys .............................................. 349 UT-SM&C-LM/WET Wetting of materials by liquid metals........................................................................... 353

Materials Development Structural Materials UT-SM&C-LAM/Mic Influence of the martensite morphology on the plasticity behaviour of the Eurofer steel ....................................................................................................... 357 UT-SM&C-LAM/Weld Laser weldability of LAM steels (Eurofer97) .............................................................. 361 UT-SM&C-LAM2 Irradiated behaviour of Reduced Activation (RA) martensitic steels after neutron irradiation at 325°C................................................................................. 363 UT-SM&C-LAM3 Microstructural investigation of Reduced Activation Ferritic-Martensitic (RAFM) steels by Small Angle Neutron Scattering (SANS) ....................................... 367 UT-SM&C-ODS Development of forming and joining technologies for ODS steels .............................. 371

- VII -

Nuclear Data UT-N-NDA Nuclear data assessment ............................................................................................... 375

Fuel Cycle UT-T1 Separation of the D/T mixture from helium in fusion reactors using superpermeable membranes - Development of superpermeable membranes suitable to operate under ion bombardment .............................................. 379

Safety and Environment UT-S-BLK Blanket Safety .............................................................................................................. 383 UT-S-MITIG Evaluation and mitigation of the risk connected with air or water ingress................... 385

System Studies UT-SM&C-REL Reliability modeling & assessment .............................................................................. 389 UT-SM&C-TDC Innovative Thermo-Dynamic Cycles studies (TDC) .................................................... 393 UT-SM&C-TMNC Thermal-hydraulic models for nuclear components ..................................................... 397

INERTIAL CONFINEMENT FUSION PROGRAMME .............................. 399

ICF01 Intense laser and particle beams dynamics for ICF applications .................................. 401 ICF02 Civil applications of inertial confinement fusion - Cryogenic targets production using magnetic levitation .............................................. 405 ICF03 Laser-matter interaction at relativistic intensities and fast igniter studies .................... 411 ICF-KiT-PRC Overview on power reactor concepts ........................................................................... 415

APPENDIX 1 : Directions contribution to the fusion programme ........................ 419

APPENDIX 2 : Allocations of tasks .......................................................................................... 423

APPENDIX 3 : Reports and publications .............................................................................. 429

APPENDIX 4 : CEA tasks in alphabetical order .............................................................. 441

APPENDIX 5 : CEA sites ................................................................................................................. 445

- 167 - EFDA technology / Tritium Breeding and Materials / Structural materials development / Advanced materials

TTMA-1.2 Task Title : SiC/SiC CERAMIC COMPOSITES SiC/SiC composite development and characterization INTRODUCTION Recent years have shown increasing interest for the characterization of silicon carbide (SiC) based ceramic matrix composites (CMCs), envisaging their application in future power reactors. Interest for SiC CMCs in fusion power applications is based on their excellent thermo-mechanical properties, their low neutron activation, as well as their dimensional stability under irradiation for temperatures up to 1273 K. The most important matters to be addressed in SiC-SiC CMCs are: - the thermal conductivity, - the radiation effect on dimensional stability and

strength, - the anisotropy of properties, - the strength of joints. In this Task, the mechanical properties of 3D SiC/SiC composites are investigated before irradiation (CEA) and after irradiation (NRG). The mechanical properties are measured at room temperature and up to 950°C , for two composite grades. 2000 ACTIVITIES This year, the objectives were as following: - To compare the bend strength of both composite grades

at room temperature. - To determine the elastic properties of both grades up to

very high temperature, as this property is needed for the evaluation of irradiation effects performed at NRG.

Due to the geometry and mechanical behaviour of ceramic composites, we decided to measure the Young modulus of the composite with an acoustic resonance technique. This allows to avoid experimental difficulties arising with the determination of bending modulus (difficulties in measuring the strain in 4 point flexural testing). Room temperature strengths are determined with a crosshead speed of 0.1 mm/min. The inner and outer spans are 20 mm and 40 mm. Before testing, the average width and thickness of each specimen is measured. The specimens are not machined or grinded before testing. This grinding process is often performed to insure a precise geometry of the samples (allowing easier modulus measurements) but we think that this may have a detrimental effect on strength, due to fiber cutting and removal of the surface seal coating.

The bend tests are performed for 5 specimens of each composite grade. Testing is done in accordance to ASTM C1341-96 test method : “Standard test method for flexural properties of continuous fiber reinforced advanced ceramic composites”. Figure 1 shows the results of bend stress versus displacement curves for the two grades. The stress-displacement curves show two distinct regions : linear and then non-linear. After the elastic region, the non-linear region corresponds to the first fiber/matrix debonding or fiber cracking and sliding. The load drops suddenly after reaching a maximum value, which is representative of a brittle fracture behaviour. The average flexural strength and its standard deviation are given by: Cerasep N3-1 : σ = 674 ± 85 MPa Cerasep N4-1 : σ = 553 ± 101 MPa The standard deviations are large for both grades, but the average values indicate that CerasepN3-1 composite has the highest strength at room temperature. These large standard deviations suggest that more than 5 tests are needed at each test temperature to have a good estimation of the real scatter that can be expected for the composites.

Figure 1 : Bend strength versus crosshead displacement for both composite grades (20°C)

The specimen fractured in the middle of outer span. Fracture initiates on the tensile (lower) face and propagates towards the upper face, as shown on figure 2. From ASTM C1341 standard, this fracture behaviour is denoted as “tensile failure” in this case. Tests are considered as “valid” as no crushing is observed under the loading pins.

0

100

200

300

400

500

600

700

800

900

0 0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9

Crosshead displacement (mm)

Ben

d s

tres

s (M

Pa)

Cerasep N3-1

Cerasep N4-1


Figure 2 : Typical fracture of Cerasep composites after 4-point bending test at room temperature.

HIGH TEMPERATURE ELASTIC PROPERTIES MEASUREMENTS The elastic modulus of both Cerasep N3-1 and Cerasep N4-1 have been measured with a Grindo-Sonic Instrument implemented to work from room temperature up to 1200°C. The specimen in the furnace is supported by two cylinders at the nodes of its fundamental vibration mode. The composite bars are supported on these cylinders with the same orientation as in the bend test The instrument analyses the transient vibration of a test object following a light mechanical input. The result is a highly accurate measurement of the natural frequency. Associated with the dimensions of the bars, the calculation of the elastic constants is possible. In our case, the natural frequencies have been measured on the same composite bars as those used for the mechanical testing. No surface preparation or machining are performed, and provided the structure is not damaged by the thermal cycle experienced in air, the bars can be further used for mechanical testing. Four tests on each composite grade are planned, with the first results on all Cerasep N3-1 bars and 2 Cerasep N4-1 bars. In our case, ff is recorded from 20°C to 1150°C, with one measurement every 5°C. The furnace temperature increases at a rate of 2°C/min. Tests are performed in air, and a measure of Ef after the high temperature testing has confirmed that the modulus is not affected by the thermal treatment. As the dimensions of the bars can not be precisely defined (due to large surface roughness), the conversion from the measured frequency ff to the Flexural stiffness Ef leads to some uncertainty for the absolute value of Ef. The evolution of Ef with temperature remains however correct, as all the values are calculated with the same values for m, b, L and t (no thermal expansion effect is taken into account). The first tests results are depicted in figure 3. The measured flexural modulus decreases very slowly with temperature : less than 10% of decrease in the range 20°C-1150°C. When looking at these results, we must take into account the uncertainties associated with the measure of ff and uncertainties when calculating the Ef values (mainly from errors on the specimen dimensions). A simple calculation shows that the uncertainty on Ef is around 5.5 % at room temperature (corresponding to ∆Ef/Ef). Therefore, the scatter observed on the three “high modulus” Cerasep N3-1 samples cannot be attributed to a material heterogeneity, whereas the behaviour of the “low modulus” sample could be due to some microstructural heterogeneity.

Figure 3 : Evolution of dynamic flexural modulus E with temperature in the range 20°C-1150°C

for both Cerasep N3-1 and Cerasep N4-1 composites CONCLUSION

The work performed this year has allowed to start the experimental program, with the two main following results: - The bend strength of both grades has been measured at

room temperature, with 5 specimens per grade. - The elastic flexural modulus of both grades has been

measured from 20°C to 1150°C The results of 4-point bending tests show large scatter in the rupture strength, which is typical for ceramic materials. More tests are needed to establish a representative database of SiC-SiC composites before irradiation. In the next year, the 4-point bend tests will be performed at high temperature (650°C and 950°C were agreed with NRG) and the failure mode will be analysed using Scanning Electron Microscopy. For the mechanical testing, we will use a specific optical extensometer developed in the frame of fusion programs . This will allow to measure the exact deflection of the bar during deformation, and therefore to have a good estimate of the bend modulus for each test bar. REPORT AND PUBLICATION [R1] R. Couturier, B. Riccetti, TTMA 1.2 – SiC-SiC

composite development and characterization. Note technique DEM n° 2000/119, 12/12/2000 TASK LEADER Raphaël COUTURIER DRT/DTEN/SMP/LS2M CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 35 59 Fax : 33 4 76 88 58 91

E-mail : [email protected]

5 mm

Cerasep N4-1

Cerasep N3-1

150

200

250

300

350

0 200 400 600 800 1000 1200 Temperature (°C)

Ef-

mo

du

lus

(GP

a)


TTMA-1.4 Task Title : SiC/SiC CERAMIC COMPOSITES Compatibility of SiC/SiC composite with liquid Pb-17Li alloy INTRODUCTION OBJECTIVE The TAURO blanket developed in the European Community by CEA is a self cooled breeder blanket using the SiCf/SiC ceramic matrix composite as structural material and liquid Pb-17Li as coolant, breeder and neutron multiplier [1]. The main objective of this concept is to permit passive safety in tokamak-type fusion power reactors by minimizing the energy potentially available for breaking the confinement building. This is achieved by using low activation structural materials like the SiCf/SiC composite and a low-pressure coolant with low chemical reactivity like the Pb-17Li eutectic. The use of the SiCf/SiC composite as structural material has shown that significant R&D is still required. Among the most significant design issues of the TAuRO blanket, the compatibility between flowing Pb-17Li and SiCf/SiC (including the brazing material) has been identified. It was shown that the compatibility of sintered SiC with lithium was very poor [2], the degradation being due to the oxygen content in the material. The structure of SiC, amorphous or crystalline [3] but also the method of fabrication of the SiCf/SiC composite (e.g., fibers pre-treatment) could influence the corrosion properties. Concerning the compatibility of SiC material with Pb-17Li, a preliminary study was reported in [4]. It was concluded that SiC should be stable in the alloy. However, the compatibility of SiCf/SiC composite with molten Pb-17Li has to be investigated more in depth.

2000 ACTIVITIES In the work performed in 2000, SiCf/SiC composite specimens were exposed to static Pb-17Li up to 3000 h. The thermal analysis of the TAURO blanket indicates that the breeder maximum temperature is about 900-990 °C [5]. For this first series of tests, the temperature was chosen equal to 800 °C. MATERIAL The SiCf/SiC composite was provided by SNECMA Moteurs (SEP division). The 3 dimensional composite Cerasep® N3-1 was used. It is made of SiC fibers (Nicalon NL207) produced by Nippon Carbon and it is densified by chemical vapor infiltration process (CVI). The chemical composition (wt %) of the SiC fibers is 56.6Si-31.7C-11.7O.

The fiber content in the composite is 40 % and its porosity is about 10 %. The composite (density > 2.4 g/cm3) was delivered in the form of sheets (25x25x3 mm3). For the corrosion tests, one sheet was used as-manufactured and another one was mechanically cut to have a naked surface (in order to detect if this machining alters the corrosion behavior as the composite is not completely densified). A direct view of the as received composite is shown in Fig. 1.

Figure 1 : Surface of the as received SiCf/SiC specimen:SEM direct view of the 25x3 mm2 surface

EXPERIMENTAL PROCEDURE The Pb-17Li pieces were cleaned with a wire brush to remove the oxide scale. After cleaning of the SiCf/SiC composite specimens in acetone and ethanol, weight and geometrical measurements were made. All the operations to introduce the specimens in crucibles with molten Pb-17Li were carried out in a glove box with purified argon gas. The Pb-17Li alloy was melted in a furnace and the liquid was poured into a Mo crucible. About 800 g of Pb-17Li was filled into the Mo crucible. Two SiCf/SiC samples were mounted into a specimen holder which was immersed into the crucible with the molten alloy. After cooling, the Mo crucible with specimens was sealed by electron beam welding. The as-sealed crucible was placed into a 601 nickel alloy container which was subsequently electron beam welded. The corrosion test was performed at 800 °C for 3000 hours. After the test, the container and the crucible was opened, the Pb-17Li melted and the samples extracted from the alloy in the glove box.

porosity


In order to measure the specimen weight variations, the specimens were dipped in a chemical mixture (1/3 ethanol, 1/3 acetic acid, 1/3 hydrogen peroxide) to remove the residual lead alloy adhering to the surface. The corrosion was also evaluated by using a scanning electron microscopy (SEM) with analysis by energy dispersive spectroscopy (EDS). Chemical analyses of the Pb-17Li alloy after corrosion tests were performed by inductively coupled plasma mass spectrometry (ICP-MS). RESULTS AND DISCUSSION As the composite material was dissolved by the chemical mixture (1/3 ethanol, 1/3 acetic acid, 1/3 hydrogen peroxide) used to remove the lead alloy adhering, it was not possible to measure the weight variation of the samples after the test. Another experimental procedure must be considered to make weight measurements in the future. Analyses indicated that after the test, Pb-17Li contains about 8 ± 4 wppm of silicon. Concerning the solubility of silicon in liquid lead, data are very limited and contradictory. Molten Pb could dissolve small quantities of Si at high temperatures (0.02 wt % at 1250 °C) but it seems that up to 900 °C, it is not much soluble. The specimens exposed to Pb-17Li were examined by SEM without removing the lead alloy adhering to the surface: a direct view of a specimen is shown in Fig. 2 where particles of solidified lead alloy can be seen. Cross sections were also observed by SEM after polishing of the surface.

Figure 2 : SEM direct view of the 25x3 mm2 surface Concerning the as received surfaces, the cross section does not show significant open porosities from the two sides of the specimen. The X-ray maps do not exhibit liquid penetration into the composite: the lead alloy is only adherent to the surface. Therefore, no corrosion damage is observed for the SiCf/SiC material exposed to Pb-17Li. Concerning the surface which was cut after manufacturing (cut surface), the cross section presented in Figure 3 clearly shows penetration of Pb-17Li into the open porosities. The presence of the liquid alloy in the free space is confirmed X-ray analyses

Figure 3 : Cross section of the SiCf/SiC specimen after exposure to Pb-17Li liquid alloy,

the vertical side corresponds to the cut surface, the horizontal side to the as received surface

CONCLUSION The SiCf/SiC composite Cerasep® N3-1 was exposed to liquid Pb-17Li at 800 °C up to 3000 h. The experiment indicates that the material has not reacted with the liquid alloy. Therefore, it should be stable and compatible with this environment. Penetration of Pb-17Li was only observed into the free space (porosity) as the composite is not highly densified. The open porosity was mainly due to the cutting of the specimen after manufacturing. A more complete study on the effect of the machining on the liquid infiltration into the composite but also on the effect of the temperature is necessary. Moreover, the determination of Si solubility in Pb-17Li as a function of the temperature will be very fruitful: it would allow to determine the temperature range in which a dissolution process could have a significant contribution to the corrosion process. Therefore, in that temperature range, the influence of the liquid flow rate will have to be evaluated.

REFERENCES [1] L. Giancarli, J.P. Bonal, A. Caso, G. LE Marois, N.B.

Morley and J.F. Salavy Design requirements for SiC/SiC composite structural

material in fusion power reactor blankets Fus. Eng. & Des. 41 (1998) [2] W. R. Drawl and K.E. Spear The compatibility of non-oxide materials with molten

lithium - Advances in ceramics 25 (1989) 195-199

Pb-17Li

500 µm


[3] M. Hansen Constitution of binary alloys, McGraw-Hill, 1958, p.

1106 [4] T. Yoneoka, S. Tanaka and T. Terai Compatibilty of SiC-SiC composite materials with

molten lithium metal and Li16-Pb84 eutectic alloy Presented to ISFNT5, 5th International Symposium On

Fusion Nuclear Technology, Rome, September 19-24, Rome (1999)

[5] G. Aiello, L. Giancarli, L. Gobert, E. philippe, Y.

Poitevin, J.F. Salavy and J. Szczepanski A parametric analysis of the TAURO SiC/SiC breeder

blanket CEA report, SERMA/LCA/RT/99-2582/A (1999)

REPORT F. BARBIER, F. HERBERT, Compatibilityof SiC/SiC composite with Pb-17Li static liquid alloy, RT-SCECF 544 (November 2000)

TASK LEADER Françoise BARBIER Person to contact: Anne TERLAIN DEN/DPC/SCCME/LECNA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 16 18 Fax : 33 1 69 08 15 86 e-mail: [email protected]


TTMA-1.6 Task Title : SiC/SiC CERAMIC COMPOSITES Brazing of SiC/SiC composites with thermo-mechanical testing INTRODUCTION Low activation SiC/SiC composites have been proposed as structural materials for Fusion Power Reactor (FPR) breeding blankets. In this prospect, DMT/SERMA studies the possibility to design a breeding blanket, so called TAURO Concept. Nowadays, two different concepts are under consideration, Pb-17Li and He cooled blankets. It was already mentioned that brazing can be used to join different complex parts of the blanket [1]. Some of the junctions can not be obtained with current joining techniques such as welding, diffusion bonding and/or textile densification stage during the manufacture of the composite component. Brazing of SiC/SiC composite parts is thus a key point for the manufacturing of large complex parts as those designed in the TAURO concept. In this context, CEA/CEREM is involved in the development of joining technologies for SiC/SiC composites thanks to the BraSiC process and of mechanical testing of the SiC/SiC junctions. Activities of the year 2000 have been devoted to brazing optimisation, in respect to the 1999 achievements, and to mechanical testing of the junctions. 2000 ACTIVITIES IMPLEMENTATION OF THE BraSiC PROCESS Requirements and recall of the previous work In the previous work, the BraSiC alloys have been selected regarding the requirements for both Pb-17Li and He cooled blankets. Then, the main problem linked to the brazing of SiC/SiC porous composites is the control of the infiltration of the brazing alloys within the composite to assess the filling of the joining area by the filler metal [1-2]. A special care has to be taken to the thermo-mechanical behaviour of the braze against the SiC/SiC composite. In 1999, we have focusing on the brazing of SiC/SiC composites with the BraSiC H and the BraSiC RE grades. These ones gave promising results and successful brazed joints have been obtained. In the case of the BraSiC® RE grade, some improvement is necessary to adapt the thermal expansion mismatch between SiC and the braze. In the 2000 activities, we have worked on the optimisation of the brazing process to get joints that are perfectly adapted to the SiC/SiC based-composite. In addition, previous results have shown a large discrepancy in the mechanical results due to a likely inadequacy of the mechanical test used for the junctions obtained with the composites.

So, it is proposed to use a new mechanical test that is suitable for the testing of brazed SiC/SiC junctions from room temperature to at least 800°C. The test and specimen dimensions have been determined taking into account the features of the composite and the test capabilities. Joining experiments The brazing trials have been performed on both CERASEP N3-1 and N4-1 composites covered by a thin SiC seal coat to limit the open porosity. Samples have been delivered lately in 2000, so to start with the experiments, bulk dense SiC supplied from Boostec company (France) has been used. The samples to be brazed are cleaned in alcohol using a Ultra Sonic device and then covered with the braze powder mixed with an organic binder. The parts are then heat treated in a metallic furnace. After cooling down, samples are cut through the joint in cross section, polished and examined using an optical or Scanning Electron Microscope. The melting point of the BraSiC® H2 grade is roughly 1300°C which is perfectly suitable regarding the maximum working and processing temperatures (1100 and 1400°C respectively). Brazing trials have been performed either under argon or high vacuum. Good results have been obtained using both atmospheres and sound joints are obtained without pores nor cracks. The elastic properties of the braze are well adapted to that of SiC and even in the fillet area where there is a large amount of additional braze, no cracks are observed (see Figure 1). The results obtained with brazing the CERASEP® N3-1 SiC/SiC composite is very satisfying as well (see Figure 2). There is a complete filling of the joint, no reactivity of the braze with the composite and no infiltration despite the porosity of the composite.

Figure 1 : Bulk dense SiC joint with a BraSiC® H alloy

joint fillet

25 µm


Figure 2 : SiC/SiC composite joint with a BraSiC® H alloy

The BraSiC® RE2 grade is interesting due to a lower melting point (roughly 1200°C) allowing to limit the damage of the fibbers when subjected to high temperatures. It has been observed perfect joints with a good micro-structural homogeneity of the joint without neither pores nor large cracks. On the edge of the joint some small cracks are observed in the braze as well as in the SiC. These cracks are due to a thermal expansion mismatch between SiC and the braze that induces thermal stresses up to failure. These cracks probably initiate in the braze and propagate through the SiC/braze interface. One possibility to limit the thermal stresses and thus possible cracks is to change the composition. Such a change is possible but with disadvantages regarding the control of the homogeneity of the microstructure. Another interesting approach is to implement “composite joints” that are made of the non-reactive BraSiC® braze and a reinforcement made of a low CTE material that can be preferably SiC due to its chemical compatibility with BraSiC® grades. Numerous experiments have been carried out using such an approach. In addition to the adaptation of the CTE, such reinforcements increase the toughness of the joints. A micrograph of such a junction is displayed in figure 4. (see figure 3 without reinforcement). It can be seen on the micrograph that such a technique allow the implementation of sound joints with the BraSiC® RE2 grade. Such joints are exempt of cracks and pull-outs featuring no or very limited residual stresses after brazing. Moreover, they allow the brazing of large gaps making easier the brazing of large components with bad tolerances. Indeed, joints with a thickness as high as 700 µm have been produced. Finally, as the reinforcement content can be changed, the properties of the joints are adapted to match one particular feature. During the year 2000, only the BraSiC® H2 grade has been tested with either the CERASEP® N3-1 and CERASEP® N4-1 respectively. The other grade will be tested in 2001.

Figure 3 : SiC/SiC composite joint with a BraSiC RE2 alloy numerous cracks and pullouts

Figure 4 : SiC/SiC composite joint with a reinforced BraSiC RE2 alloy no cracks nor pullouts

MECHANICAL TESTING To be more representative and close to the final probable joints, the joint specimens have been designed as shown on the figure 5.a). Moreover, this simple geometry can be processed without machining of the composite bars (only cutting was performed) and, when tested in compression, the fracture load is related to the shear strength of the brazing alloy (shear strength = load / brazed surface) when failure occurs in the joint.

Figure 5.a) Specimen geometry

Figure 5.b) Brazed Specimen with asymmetrical

fillets

300 µm 330 µm

330 µm

braze

4 mm

5 m

m


Ten brazed specimens have been produced with the BraSiC® H2 grade and each of the CERASEP® N3-1 and N4-1 composite type. Tests have been done in air at 20°C (8 samples) and 800°C (2 samples) for both composite grades. Specimens are tested in compression with a crosshead speed of 0.1 mm/min. Each sample geometry has been observed and measured before testing. At room temperature, the average shear strength is around 100 MPa and depend on the variations of the brazing thickness (a large thickness is thought to be detrimental). The size of fillets is also a relevant parameter. Indeed, some samples with low strength have asymmetrical fillets that seem detrimental to the structure strength (see figure 5.b). In the case of CERASEP® N4-1, experimental shear strengths (40±16 MPa) are lower than with CERASEP® N3-1 composites. The specimen failure occurred mainly by two mechanisms: - failure of the braze as with all CERASEP® N3-1

specimens, - failure of the composite in the inner bar, starting from

the bottom. In this case, the cutting of the composite bar has probably introduced a “weak point” in the specimen structure and the test result is not related to the true shear strength of the joint.

When looking at the shear strength of the samples with regards to the geometrical variations observed, no obvious correlation appears and therefore it seems that the test we developed is fairly well suited to compare different joints or brazed composites. In our case, it appears that CERASEP® N3-1 composite lead to better shear strength values that the CERASEP® N4-1. This feature is evident when looking at the figures obtained but at that time there is no clear explanations for that phenomenon. CONCLUSION BraSiC® H an RE2 alloys have been selected regarding the requirements of the new TAURO concepts, with an estimated maximum working temperature of about 1100°C. Both BraSiC® grades give promising results and successful brazed joints have been obtained using the CERASEP® N3-1 composite. In the case of the BraSiC® RE2 grade, a adaptation of the thermal expansion mismatch between SiC and the braze has been proposed thanks to composite joints including a SiC powder reinforcement. Thus, it allows a limitation and even suppression of thermal stresses that may occur during brazing. We have produced specific shear test specimens with CERASEP® N3-1 and N4-1 composite grades. They have been tested in compression at 20°C and 800°C, with interesting results for the CERASEP® N3-1 composite grade with an average shear strength of 100 MPa at room temperature

REFERENCES [1] G. Chaumat, A. Gasse, G. Coing-Boyat, G. Bourgeois,

"Long term programme - Stage I. Development of joining techniques for the low activation materials (LAMs) SiC-SiC composite", Note technique D.E.M./S.G.M. n°4/96, 9 janvier 1996.

[2] A. Gasse, F. Saint-Antonin, G. Coing-Boyat, Specific

non-reactive BraSiC® alloys for SiC/SiC joining, Note Technique DEM N° DR 25/97, Décembre 1997.

TASK LEADER Adrien GASSE DRT / DTEN / SMP CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 57 44 Fax : 33 4 76 88 51 17

- 177 - EFDA technology / Tritium Breeding and Materials / Structural materials development / RAFM steels

SM-2-3-1 Task Title : MECHANICAL PROPERTIES OF F82H WELDMENTS INTRODUCTION The development of fast induced radioactivity decay alloys has been incorporated into the goals of the IEA collaborative program in order to reduce the quantity of highly radioactive materials from system components. Concerning the innermost parts of a Fusion reactor, the low activation martensitic (LAM) steels and reduced-activation ferritic/martensitic (RAFM) steels are considered as primary candidate alloys as first wall and breeder blanket structural materials because of high irradiation resistance, especially the resistance to the phase instability and the resistance to the shift of the ductile-brittle transition temperature (DBTT) during or after irradiation. Several RAFM steels have been developed specially for this application in the frame of the Long Term Fusion international collaborative test program. Among these materials, the F82H steel developed by JAERI in the mid nineties and the more recent Eurofer 97 proposed by FZK and Böhler are considered to be specially viable candidates. Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding and laser welding processes are envisaged for the fabrication of structural parts of relevant breeding blanket modules. The present task aims to evaluate the tensile and impact test properties of F82H TIG welded joints processed with F82H filler material. During the year 2000, F82H welded joints of 24 mm of thickness were produced. The specimens were sampled in these weldments for tensile and impact test characterisation.

The sampling design has been settled in keeping with the tests previously carried out on weldments processed at JAERI [1]. The work will continue in the frame of the task TTMS-2.3.2. 2000 ACTIVITIES PRODUCTION OF WELDED JOINTS FOR CHARACTERIZATION F82H TIG welded joints of 24 mm of thickness have been processed in order to sample the tests specimens. Two assemblies of 200 mm of length each have been welded. Each assembly was made of two plates coming from the JAERI heat n° 9753. The operating conditions have been selected according to the preliminary weldability results coming from the task SM-4.3, as summarised in table 1. Both weldments have displayed a suitable external aspect. Further, they have followed radiographic control indicating that they are free of internal flaws. The weldments have then followed a post welding heat treatment of 1 h at 720°C. The treatment has lowed the hardness down to about 220 Hv1 in the seam, that is similar to the hardness of the base material. SAMPLING OF TESTS SPECIMENS The test matrix has planned 7 tensile specimens to be sampled either in the longitudinal direction or in the cross direction and two series of 12 impact tests specimens to be sampled in the melted zone or in the heat affected zone (Table 2).

Table 1 : Operating conditions used to process 24 mm TIG welded joints for mechanical characterization

mat

eria

l

nb o

f pa

sses

thic

k (m

m)

clam

ping

U (

V)

I (A

)

wel

ding

spe

ed (

cm/m

in)

wir

e sp

eed

(m/m

in)

wir

e di

amet

er (

mm

)

scan

ning

wid

th (

mm

)

uppe

r sh

ield

ing

(l/m

in)

back

shi

eldi

ng (

l/min

)

back

side

shi

eldi

ng (

l/min

)

F82H 9 24 welded 10 150-230 4 0,5-1,3 1,2 4-9,5 12 20 7,5


Table 2 : Tensile and impact tests matrix

Centred MZ Centred HAZCentred MZ

crossCentred MZ

longi.

A 3 T 08 A 3 T 20 A 1 T 01 A 1 L 03

A 3 T 09 A 3 T 21 A 1 T 02 A 1 L 04

A 3 T 10 A 3 T 22 A 1 L 05

A 3 T 11 A 3 T 23 A 1 L 06

A 3 T 12 A 3 T 24 A 1 L 07

A 3 T 13 A 3 T 25

A 3 T 14 A 3 T 26

A 3 T 15 A 3 T 27

A 3 T 16 A 3 T 28

A 3 T 17 A 3 T 29

A 3 T 18 A 3 T 30

A 3 T 19 A 3 T 31

Material Process

Tensile (20 ; 400 ; 500 ; 600°C)

Cylindrical M7

F82H TIG filler 24

Impact tests

StandardThick. (mm)

A small cylindrical specimen design has been selected to carry out common tensile tests at various temperatures (Figure 1). The standard V-notch design has been used for the impact specimens.

The rough shapes of the specimens have been sampled according to the scheme of Figure 2. They have been etched in order to reveal the seam zones. Then, this allowed to centre precisely the specimens either on the melted zone or on the heat affected zone during the final machining. Each specimen has been engraved at both ends. CONCLUSIONS F82H TIG welded joints of 24 mm of thickness have been processed using F82H filler metal. Tensile and impact tests specimens have been sampled within heat affected zones and melted zones of the joints. The test specimens have been etched before final grinding in order to locate precisely the notch respectively at the centre of the HAZ or MZ. They are now available for mechanical testing in the frame of TTMS-2.3.2.

F0A0302

* 2 * 2 * 2

* 2 * 2 * 2

11,5

A

A

Eprouvettes de résilience F82 H

60

A-A

11,5

* 2 * 2 * 2

* 2 * 2 * 2

A3T08

A3T09

A3T10

A3T11

A3T12

A3T13

A3T14

A3T15

A3T16

A3T17

A3T18

A3T19

A3T20

A3T21

A3T22

A3T23

A3T24

A3T25

A3T26

A3T27

A3T28

A3T29

A3T30

A3T31

Centré en ZF

Centré en ZAT

F0A0303

40

* 2* 2

28 28

* 2

8

Eprouvettes de traction F82 H

Fin de cordon

A-A

8

A

A

* 1 28

A1L03

A1L04

A1L05

A1L06

A1L07 A1T01

A1T02

Figure 2 : Design of the specimen sampling

Figure 1 : Tensile specimen design


REFERENCES [1] A. FONTES, F. CASTILAN, "Coupons d'acier de type

Z 9 CWV 0802 (F82H) soudés par faisceau d'électrons ou par procédé TIG (provenance JAERI) : caractérisation mécanique 2ème partie", rapport CEA STA/LMS/99-RT3794, février-99

REPORTS AND PUBLICATIONS A. FONTES, “F82H TIG weldments and sampling for mechanical testing”, CEA report DECS/STA/LMS/01-RT4004

TASK LEADER André FONTES DRT/DECS/STA/LMS CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 37 21 Fax : 33 1 69 08 75 97 E-mail : [email protected]


SM-3-3 Task Title : CORROSION IN WATER CONDITIONS Corrosion water with additives INTRODUCTION In the framework of the task SM3.3 concerning the Structural Materials, a study is carried out to assess: - the general corrosion behaviour, - the pitting susceptibility and, - the Stress Corrosion Cracking (SCC) susceptibility of

Low Activation (LA) steels. Theses materials are EUROFER97 and F82H.

The first material is a new material whose corrosion behavior is unknown. The second one has already been studied but its comportment in FUSION representative environment remains poorly known. 2000 ACTIVITIES MATERIALS The two LA materials studied are: - F82H (heat NKK RB 801 6.4) and, - EUROFER97 (heat E83698). The first material was supplied by JAERI-NKK as plates of 8 mm thick and the second one by BÖHLER as plates of 1.5 mm thick. The chromium content of the EUROFER97 is higher than the one of the F82H: respectively 7.66 wt% and 8.96 wt%. This concentration range is optimal against neutron irradiation embrittlement. The tungsten concentrations of the F82H and the EUROFER97 are respectively 2.00 wt% and 1.08 wt%. Both materials have some vanadium: respectively 0.16 wt% and 0.20 wt%. The EUROFER97 has some tantalum (0.14 wt %) and the F82H has only 0.02 wt % of this component. The EUROFER97 is tested as received: Normalized 980°C/0.5 hour then Tempered 760°C/1.5 hour (N&T). According to the partners, the F82H is tested for two metallurgical conditions: - as quenched: Normalized 1040°C/0.5hour (N) and, - as received: Normalized 1040°C/0.5hour then Tempered

625°C/1 hour (N&T). During 2000, the microstructural properties and the mechanical properties of these materials were characterized.

The main effects of the heat treatment at 625°C consist in modifying the microstructural properties and the mechanical properties of the material. The microhardness and the 0.2% yield stress are lower than the ones of material only normalized: Rp0.2 = 1250 MPa for F82H N, Rp0.2 = 460 MPa for F82H N&T, HV(4.905N) = 420 for F82H N and HV(4.905N) = 260 for F82H N&T. The microstructure of the F82H N&T consists of laths of tempered martensite within prior austenite grains (α+γ). Partially recristallised zones are observed. The precipitation consists of fine M23C6 precipitates located along grain and lath/subgrain boundaries. The typical composition of these precipitates is (at%) 60Cr-30Fe-5V-5W [1,2]. The average grain size is 100 µm. For the F82H N, the microstructure is a quenched martensite without precipitation. The average grain size is about 100 µm, as for the N&T F82H. It can be noticed that the chromium content of the matrix is about 1% higher because of the lack of precipitation. For the EUROFER97, the Vickers microhardness is about 230 HV(4.905N). The grain size of the high temperature austenitic phase is between 10 and 15 µm. The microstructure is martensitic without delta ferrite. EXPERIMENTAL The tests are performed in a corrosion loop [3,4]. This loop consists of an autoclave linked to a water circuit allowing to insure a continuous water flow and to control the medium chemistry (Figure 1).

Figure 1 : View of the corrosion and permeation loop


The test medium was prepared by adding to deionized water (resistivity 18 MΩ.cm) lithium as lithia (LiOH) in order to have a pH of 7 at the test temperature: 320°C. 2.40 10-5 mol of Li was added to 1 kg of deionized water ([Li] = 166 ppb). The flow rate of the medium was 10 L.h-1 and the working pressure was 15 MPa. A hydrogen overpressure, controlled with a Pd/Ag probe, was applied to maintain a hydrogen concentration in the liquid medium of about 35 mL(SPT).kg-1 at the test temperature (320°C). A total of twenty-one specimens are tested in the corrosion loop, 7 specimens of each material. Four specimens of each material are used for general and pitting corrosion and three specimens of each material for stress corrosion cracking. The same specimens are used for general corrosion and pitting corrosion. SPECIMENS During 2000, the specimens were prepared and characterized. The specimens for general and pitting corrosion are coupons of length 50 mm, width 10 mm and thickness 1 mm. The surfaces of the specimens were machined with a grinding machine. So, they have a cold worked layer of small depth. This depth is about 10-20 µm for all the grades tested in this study. The Vickers microhardness at the surface of the specimen is higher than the one of the bulk metal: - For the F82H N, HV(0.49) = 445 at the surface and

HV(1.981N) = 420 in the core of the coupons. - For the F82H N&T, HV(0.49) = 325 at the surface and

HV(1.981N) = 260 in the core of the coupons. - For the EUROFER97, HV(0.49) = 255 at the surface

and HV(1.981N) = 230 in the core of the coupons. Usually, the cold work of the surface increases general corrosion: many studies have shown that mass variations are more important on specimens with a cold worked layer than on specimens without a cold worked layer. The roughness of the surfaces of the specimens has also been characterized. For all the specimens, the roughnesses are roughly identical: Ra = 0.5 µm. The specimens for the SCC tests are U-Bend specimens (Figure 2). These specimens are obtained by bending the coupons previously described. The bending of the specimens occurs in two steps: - The coupon is plastically strained in order to have an

angle of about 45° between the two legs of the specimen.

- After this first stage, the two legs of the specimens are

made parallel with a nut and a screw.

The total strain on the outside of the bend is roughly equal to 12.5 %. The stress induced depends on the mechanical properties of the tested materials. By taking into account the curves σ versus ε of the materials, the stress level at the apex of the U-bends at 320°C are approximately: - 1200 MPa for the specimens in F82H N, - 400 MPa for the specimens in F82H N&T and, - 600 MPa for the specimens in EUROFER97. It must be noticed that this estimation does not take into account the cold worked layer at the surface of the coupons. As the surface of the material was previously work hardened by the grinding machine, the stress level at the outside of the specimen should be more important than the values given above. Some X-ray diffraction measurements will be performed in 2001 to assess the real value of the main stress at the surface of one specimen of each grade. Before testing these specimens, the outside of the bend was examined with a scanning electron microscope. Some surface defects were observed for all the materials. The density of defects is more important for the F82H N. These defects, mainly oriented in the direction of the applied stress, can reach a length of 100 µm, but have a very small depth (< 20 µm). The depths of the surface defects are in accordance with the depth of the cold worked layer induced by the grinding machining. We have observed in the framework of an other study carried out with Alloy 690 that by removing the cold worked layer by the mean of an electrolytic polishing, such a surface defects do not appear at the surface specimen during the bending. The reasons why these defects appear in the cold worked layer are both the "low" ductility of the material and mainly the high level of deformation at the surface of the cold worked layer. In this study, the cold worked layer has not been removed in order to have a surface representative of the surface of the components of the future FUSION reactor.

Figure2 : Schematic representation of a U-Bend specimen for SCC test

RESULTS According to the test matrix and the time schedule given in the CEA Report 540 [5], the first test of duration 1000 hours was carried out in 2000.


Some mass variation measurements were performed on 4 specimens of each material (Table 1). For all the materials, the mass of the specimens decreases during the corrosion test. This result is induced by the surface film and/or matrix dissolution, or by a bad adhesion of the surface film. For a stainless steel, a low increase of the mass is usually observed. These experimental results will be complemented with mass variation measurements obtained after 2000, 3000, 4000 and 5000 hours to evaluate the kinetic of general corrosion.

Table 1 : Mass variation measurements after the first period of duration 1000 hours

Materials F82H N F82H N&T EUROFER97

Average mass variation after

1000 h. (mg/dm2) -33 -27 -38

Lower value/ higher value

-28 / -38 -23 / -35 -36 / -41

Standard deviation 4 5 3

One coupon of each material was weighted after removal of the adherent corrosion products. The decreasing of the mass were 77 mg/dm2 for EUROFER97, 75 mg/dm2 forF82H N and 68 mg/dm2 for F82H N&T. The mass variations after removal of the corrosion products will be also obtained after 2000, 3000, 4000 and 5000 hours for two or three coupons of each materials. These experimental data will allow to assess the evolution versus time of the amount of metal oxidized. The coupons were also observed with a microscope (4 specimens of each material) in order to evaluate the pitting susceptibility of the three materials. No pits were observed after 1000 hours of test duration. Three U-bends of each material were observed with a microscope. No cracks were observed after 1000 hours. After this examination, the U-bends were replaced in the autoclave for a test duration of 2000 hours. CONCLUSION The tests performed in the framework of this study will allow a better knowledge of the corrosion resistance of F82H N, F82H N&T and EUROFER97 in a hydrogenated aqueous medium. EUROFER97 is a new material that has not yet been characterized. For the F82H, its corrosion comportment is very poorly known as for many low activation ferritic-martensitic materials. The results obtained in 2000 are not enough numerous to give conclusions about the corrosion behavior of the materials but it can be noticed, from the results obtained after the first period of duration 1000 hours, that the corrosion behavior of the three materials are roughly identical.

REFERENCES [1] "Metallurgical and mechanical characterization of

reduced activation ferritic-martensitic steels", WP SM 2-1 and WP SM 2-2, A. Alamo, Fusion technology, Annual report of the Association EUROMAT/CEA, 1999.

[2] Physical metallurgy and mechanical behavior of

FeCRWTaV low activation martensitic steels: effects of chemical composition", A. Alamo, J.C. Brachet, A. Castaing, C. Lepoittevin, F. Barcelo, Journal of nuclear materials, 258-263, 1998, pp 1228-1235.

[3] "Tritium permeation from Pb-17Li towards water and

corrosion loop", T. DUFRENOY, A. TERLAIN, RT-SCECF 513 (December 1999).

[4] "Material compatibility in fusion environment – WP

SM3-5.1: corrosion in water with additives – final report", O. RAQUET, CEA report N°535 (July 2000).

[5] "Material compatibility in fusion environment – WP

SM3.3: aqueous corrosion of low activation steels – progress report", O. RAQUET, CEA report N°540 (october 2000).

TASK LEADER O. RAQUET DEN/DPC/SCCME/LECA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 15 92 Fax : 33 1 69 08 15 86 E-mail : [email protected]


SM-3-5.1 Task Title : COMPATIBILITY WITH HYDROGEN AND LIQUID SCC behaviour of EUROFER in aqueous environments INTRODUCTION The main objective of this study is to perform Constant Extension Rate Tests (CERTs) to characterize the Stress Corrosion Cracking (SCC) susceptibility of EUROFER in aqueous environments. This study complements the results obtained in the framework of the SM3-3. 2000 ACTIVITIES

MATERIAL The material studied in the framework of this study is EUROFER97 (heat E83698). This material was supplied by BÖHLER as plates of 1.5 mm thick. The chemical composition of this material is given in Table 1.

Table 1 : Chemical composition of EUROFER97 (in Wt %)

Material Eurofer97

C 0.11

Cr 8.96

Mn 0.49

Si 0.04

W 1.08

V 0.20

Ni 0.02

Ta 0.14

N 0.021

P <0.005

S 0.004

Fe Bal.

The material is tested as received: Normalized 980°C/0.5 hour then Tempered 760°C/1.5 hour (N&T). During 2000, the microstructural properties and the mechanical properties of the material were characterized.

The microstructure is martensitic without delta ferrite. The grain size of the high temperature austenitic phase is between 10 and 15 µm. The Vickers microhardness of the bulk metal is about 230 HV(4.905N).

Figure 1 : Microstructure of the EUROFER97 N&T (Magnification 200)

Chemical etching HNO3 + HCl + glycerol The tensile properties of this material at 325°C are: - 0.2% yield stress: 550 MPa, - ultimate tensile stress: 675 MPa, - total elongation: 25%. EXPERIMENTAL The CERTs will be performed with the facility device given in Figure 2. The autoclave is made of AISI 316L stainless steel and has a 1.23-litre inner volume. On the bottom of the autoclave is fixed a Pd/Ag membrane for measuring the hydrogen pressure in equilibrium with the dissolved hydrogen in the water. The autoclave heating is insured by 3 heating resistances rolled round the external surface of the autoclave. The design temperature and pressure of the autoclave have been respectively 400°C and 20 MPa. All the tests will be performed at the same elongation rate: 5x10-8 s-1.


Figure 2 : View of the facility device for CERTs The corrosion tests will be performed in water with additives at 320°C. The exact water chemistry will be defined after discussions with all the partners. SPECIMENS During 2000, the specimens were machined. The specimens are tensile specimens of length 50 mm, width 4 mm and thickness 1 mm. The surfaces of the specimens were machined with a grinding machine. During 2001, the cold worked layer and the surface roughness will be characterized. PROGRAM In a first time, one CERT will be performed with the same test medium as the one used in the framework of the SM3.3: hydrogenated water of resistivity 18 MΩ.cm with 166 ppb of Li (as LiOH) in order to have a pH of 7 at 320°C. In a second time, if cracking occurs, some CERTs interrupted after various durations will be performed in order to characterize crack initiation and crack propagation. If cracking do not occur with the first test medium, some CERTs will be performed with others test mediums containing some additives.

CONCLUSION The tests performed in the framework of this study will allow a better knowledge of the SCC resistance of EUROFER97 in a hydrogenated aqueous medium with or without additives. This program complements the program performed in the framework of the task SM3-3. TASK LEADER O. RAQUET DEN/DPC/SCCME/LECA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 15 92 Fax : 33 1 69 08 15 86 E-mail : [email protected]


SM-4-3 Task Title : F82H WELDABILITY WITH FILLER METAL INTRODUCTION The development of fast induced radioactivity decay alloys has been incorporated into the goals of the IEA collaborative program in order to reduce the quantity of highly radioactive materials from system components. Concerning the innermost parts of a Fusion reactor, the low activation martensitic (LAM) steels and reduced-activation ferritic/martensitic (RAFM) steels are considered as primary candidate alloys as first wall and breeder blanket structural materials because of high irradiation resistance, especially the resistance to the phase instability and the resistance to the shift of the ductile-brittle transition temperature (DBTT) during or after irradiation. Several RAFM steels have been developed specially for this application in the frame of the Long Term Fusion international collaborative test program. Among these materials, the F82H steel developed by JAERI in the mid nineties and the more recent Eurofer 97 proposed by FZK and Böhler are considered to be specially viable candidates. Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding and laser welding processes are envisaged for the fabrication of structural parts of relevant breeding blanket modules. The present study aims to evaluate the TIG weldability of the F82H steel by using F82H filler material in order to provide EU references to be compared to JAERI previous weldability results [1]. The 1999 work has led to sound F82H welded joints of 14 mm of thickness. 2000 results concern the improvement of the welding conditions of plates of 14 mm of thickness and the study of the weldability for higher thickness (24 mm). This work is continued in the frame of the task TTMS-4.3 on Eurofer 97. 2000 ACTIVITIES OPERATING CONDITIONS The 14-mm thick specimens come from both F82H heats (9741 and 9753) that have been elaborated by NKK Corporation according to requirements from JAERI. The 24-mm thick specimens come exclusively from the 9753 heat. The 24-mm weldability samples have been designed as assemblies of plates with chamfer. A specific U-shape chamfer has been designed as in the case of the 14-mm thick plates (Figure 1). This design is supposed to prevent high welding strains.

The plates to be welded are 200 mm of length and 80 mm of width. Five assemblies (10 plates) have been used specifically for the weldability study. Two more assemblies have been produced in the frame of the task SM-2.3.1 for mechanical characterisation. F82H wire of 1.2 mm of diameter delivered by JAERI has been used as filler.

8 mm

R : 5 mm

2 mm

R : 6 mm

2 mm

7 mm

14 mm

24 mm

Figure 1 : U-shape chamfers of the weldability specimens

Before joining, the plates are directly welded around upon a rigid support which is a steel plate of 40 mm of thickness (Figure 2). This strong clamping of the plates is supposed to be somehow representative of a mechanical tightening of a large structural part during welding. It aims to estimate the risk of eventual cracking due to welding (hot tearing, cold cracking).

Figure 2 : Welded clamping of the plates to be joined

The operating conditions used for the 24-mm plates have come from the 1999 preliminary results on 14-mm plates. They are summarised in Table 1. The deep chamfer shape implies high length of uncovered electrode for the first passes (Figure 3). In order to avoid porosities due to oxidation, a backside shielding gas is user for these first passes.


Table 1 : General welding conditions of F82H plates of 14 and 24 mm of thicknesses

mat

eria

l

nb o

f pa

sses

thic

k (m

m)

clam

ping

U (

V)

I (A

)

wel

ding

spe

ed

(cm

/min

)

wir

e sp

eed

(m/m

in)

wir

e di

amet

er (

mm

)

scan

ning

wid

th

(mm

)

uppe

r sh

ield

ing

(l/m

in)

back

shi

eldi

ng

(l/m

in)

back

side

shi

eldi

ng

(l/m

in)

F82H 9 24 welded 10 150-230 4 0,5-1,3 1,2 4-9,5 12 20 7,5

Figure 3 : High length of uncovered electrode in the bottom of deep chamfer

As previously for 14-mm plates, one 24-mm assembly has been instrumented with thermocouples in order to measure the cooling time between 800 and 500°C (∆t800-500) and also in order to follow the interpass temperature. However, in contrast with the 1999 work, no preheating was used so that the operating conditions were supposed to be more severe in terms of risk for cold cracking. Each weldments has been X-Rayed. Microhardness measurements (Hv1) have been carried out in the seams at the as-welded state.

RESULTS If respecting suitable shielding conditions in order to avoid oxidation cavities, the TIG weldability of the F82H was shown to be easy. Only very few internal flaws occured in terms of small cavities which diameter were less than about 1 mm. Even when applying severe clamping without any preheating, no cracking was observed within the 24-mm weldments as for 14-mm weldments. The thermal cycles kinetics in the heat affected zone of 24-mm weldments during welding is close to that of 14-mm weldments preheated at 250°C and slightly faster than that of unpreheated 14-mm weldments. A representative thermal evolution in the heat affected zone of 24-mm weldments during the whole welding process (9 passes) is given in figure 4 as it has been measured by thermocouples. The thermal results are summarised in table 2 for both thicknesses.

Evolution de la température - F0A0201

0

100

200

300

400

500

600

700

800

900

0 50 100 150 200 250 300 350 400

Temps en minutes

Tem

pér

atu

re e

n °

C

T1 T2 T3 T4

T

Tr 800-500 = 19 s

T T T

hA

B

Figure 4 : Thermal evolution in the heat affected zone during the welding of 24-mm plates


Table 2 : Comparative results of thermal instrumentation

Material Test Thickness (mm) Preheating Pass h (mm) T0 (°C) Energy (J/cm) Dt (s)

F9A0101 4e 40 33000 42,5

F9A0101 5e 57 33000 26,5

F9A0102 4e 3,75 72 33000 41,0

F9A0102 5e 0,95 24 33000 22,5

F0A0101 3e 6,85 209 30000 22,0

F0A0101 4e 4,6 208 33000 19,0

F0A0201 24 no 7e 4,2 25 34500 19,0

?

F82 H14

no

250° C

800500

In each case, the cooling rates led to similar hard martensite of hardness of about 400 Hv1. This seam microstructure needs to follow a post welding heat treatment which conditions have been defined previously. CONCLUSIONS The TIG weldability of thick F82H plates (24 mm of thickness) has been studied by using matching filler material previously delivered by JAERI. Severe operating conditions have been imposed to evaluate the risk of eventual cracking : the plates have been strongly clamped (by welding) to a rigid support and without any preheating. Even in these conditions, the weldments were sound or displayed only acceptable (in RCCM terms) small cavities. In particular, no cracking has occurred. In conclusion, the F82H was shown to be easily weldable at the welding speed that has been used (4 cm/min) which is close to an average manual welding speed. This result is in keeping with the weldability result from JAERI. REFERENCES [1] A. Fontes, "Welded mod. F82H steel from JAERI

intended for the EEC participants", CEA report, STA/LMS/96-NT798, November-96.

REPORTS AND PUBLICATIONS A. Fontes, O. Dubet, "TIG weldability of F82H with filler matching filler metal", CEA report DECS/STA/LMS/01-RT4005 TASK LEADER André FONTES DRT/DECS/STA/LMS CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 37 21 Fax : 33 1 69 08 75 97 E-mail : [email protected]


SM-5-7 Task Title : SMALL SIZE SPECIMEN TECHNOLOGY INTRODUCTION Using testing volumes in neutron irradiation devices such as: ion beams, spallation-, 14 MeV neutron sources and material test reactors, requires large fractions of R&D budgets. There is a continuous push for irradiation of smaller test specimens to reduce exposure expenditure. With decreasing volume the representativity of the specimens for bulk properties might decrease. This is certainly the case for shape sensitive samples such as those with artificial defects (toughness and crack propagation testing), and discontinuities; cylindrical fatigue specimens with weldments also cause problems. The decrease in size results in non-standard specimens, usually with a near standard shape but a considerably smaller scale. The validity of non-standard specimens must be proven.

The development of smaller specimen is now aiming for sizes beneficial for neutron sources. The results will also be valuable for other irradiation programmes. 2000 ACTIVITIES The year 1999 was devoted to the feasibility of small tensile specimen testing. Two sizes of specimen were considered. The largest one has a section of 2 x 1 mm and the smallest one has a section of 1 x 0.4 mm. For the largest specimen, an in-situ observation technique has been developed in 2000. This technique was used to characterise the strain localisation and the plastic behaviour beyond necking during a tensile test at room temperature (Figure 1).

image 1 image 2

image 3 image 4

0

300

600

900

1200

1500

0 0,5 1 1,5 2Elongation in mm

Lo

ad in

N

images 1 to 4

4

32

1

Figure 1 : In-situ observations and corresponding tensile curve on tempered EM10 steel at room temperature

- 192 - EFDA technology / Tritium Breeding and Materials / Structural materials development / RAFM steels This technique will provide interesting information on the ductility of irradiated steels showing low uniform strains properties. For the smallest specimen, the qualification of the tests rigs has been performed in 2000 (Figure 2).

0

100

200

300

400

500

600

700

800

0 0,02 0,04 0,06 0,08 0,1 0,12 0,14

Strain

Str

ess

in M

Pa

small specimens tested at 20°C

full scale specimen tested at 30°C

full scale specimen tested at 0°C

Figure 2 : Tensile test at room temperature on low alloy steel using small specimens

Comparison with full scale specimens During 2000, the feasibility of toughness tests on small specimens has been studied. Fatigue pre-cracking and J-da tests have been performed on round compact specimens of 13mm diameter and 5mm thick (CTR5). First results have been obtained at room temperature on 316L steel strained at 45% (Figure 3). The complete qualification of the test will be achieved in 2001.

0

10

20

30

40

50

0 0,2 0,4 0,6 0,8 1 1,2 1,4

∆∆a in mm

J in

kJ/

m²

J1c = 17,2 kJ/m²J0,2 = 19,6 kJ/m²dJ/da = 11,8 MPa

Figure 3 : Toughness results on CTR5 specimens tested at room temperature. 316L steel strained at 45%.

REFERENCES "SM5.7 : SMALL SIZE SPECIMEN TECHNOLOGY. PROGRESS REPORT" P. Lamagnère, P. Wident, B.Marini NT SRMA, to be issued. TASK LEADER B. MARINI DEN/Saclay/DMN/SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 85 99 Fax : 33 1 69 08 85 99 E-mail : [email protected]


SM-6-1 Task Title : DEVELOPMENT OF HIGH DOSE HIGH TEMPERATURE

Low temperature, high dose irradiation test with ISTC INTRODUCTION In fusion reactors, blanket damage levels over 70 dpa are anticipated. Present day beam devices and Material Test Reactors cannot produce such displacement levels in a reasonable time. Liquid Metal cooled Fast Breeder Reactors are, in contrast, capable of providing high dose levels in acceptable times, but at temperatures 45O°C and above. While such temperatures are pertinent to most fusion components, there are those that will be fully or partially exposed to lower temperatures. As a result, the effects of high dose neutron irradiation have to be investigated at both low and high temperatures. The overall objective of SM 6 is to investigate the effects of high dose neutron irradiation on Reduced Activation Ferritic/Martensitic (RAFM) Steels, i.e. the primary structural material retained for ITER Test blanket Modules and DEMO. The sub-task SM 6.1 allocated to CEA has three objectives: - Contribute to a European irradiation experiment in

Russian reactors. - Investigate the possibility of placing a low temperature

rig in Phénix. - Perform a high dose irradiation experiment in Phénix. This report concludes the activities engaged at CEA under this task. 2000 ACTIVITIES CONTRIBUTION TO THE EUROPEAN IRRADIA-TION EXPERIMENT IN RF REACTORS This action has been performed in two stages. Visit of Russian center and reactor The visit and meeting took place at Dimitrograd. It was concluded that performing an irradiation experiment in BOR 60 reactor is feasible. The experiment would offer the following possibilities: - Material capsule with a low longitudinal temperature

gradient and loading/unloading capability at intermediate exposure. Example of specimen loadings are: 12 pressurized tubes, 66 tensile specimens of different sizes, 80 sub-size Charpy specimens.

- Irradiation temperature in the range of 320-330°C (can be adjusted to other values) and dose rates of about 20 dpa/year.

- All PIE can be carried out at Dimitrovgrad. Main negotiations These negotiations were led by FZK. The irradiation parameters retained are 30 dpa at 320 to 340°C. The experiment, to be loaded in BOR 60 reactor, would contain more than 300 specimens. They are of types: - Tension. - Charpy-V impact. - Fatigue. The materials to be tested are: - Eurofer 97 steel with different heat treatments. - F82H steel (IEA heat). - Optifer IV steel. - Small Eurofer heat with different amounts of boron. HIGH DOSE, LOW TEMPERATURE IRRADIATION IN PHÉNIX A project code named CALIPH was defined with the help of specialists from Phénix. A loop, cooled to about 300°C, would be placed inside the reactor to allow low temperature high dose irradiation of specimens. The project underwent extensive scrutiny, mainly to reduce cost and satisfy the reactor safety requirements. Despite all these efforts, the task had to be reoriented towards higher temperature irradiations. HIGH DOSE, HIGH TEMPERATURE IRRADIA-TION IN PHÉNIX This experiment is code named ELIXIR. The irradiation rig is a DIMEP “Diluent” (DDC0003), comprising of a central capsule and 6 steel pins around. The rig is placed in the position 15-25 of reactor. The temperature inside the rig is controlled by argon circulating inside the rig. Specimens are placed inside baskets, they are filled according to 7 plans and positioned at 15 levels inside the rig (level 1 is situated at the bottom of the rig and level 15 at the top of the rig), see Table 1.


Table 1 : Specimens placed in the ELIXIR rig

Type of specimen Basket No. In rig from

top to Bottom

Φt/Φt max

Target Dose dpa

T calcul. Specimen

°C Fatigue (GI 119)

Tension (GI 116)

Impact (GI 90-1)

Impact Sz (GI 91)

Fracture (GI 94-10)

Creep (GI 87c)

15 0.29 9.3 433 F82H (3)

OPT Ia (8) OPT II (8)

14 0.45 14.4 431 OPT IV (8)

F82H Mod (8)

OPT Ia (8) OPT II (8)

13 0.6 19.2 429 OPT IV (8)

F82H Mod (8)

12 0.73 23.4 419

F82H (2) JLF1 (2)

(75 & 150MPa) plus lames

F82H (8) JLF1 (8)

11 0.83 26.6 416 LA12 LC (8)

EM10 (8)

10 0.9 28.8 411 FMS PNC (7)

FMS PNC Soud (3)

9 0.95 30.4 409 F82H (6) F82H (8)

8 0.99 31.7 403

F82H (2) JLF1 (2)


7 1 32 399 (Fl) 407 (CT) 304-1H (2)

316-1E (2) (180 & 250MPa)

plus lames

304-1H (2) 316-1E (2) 316-1H (2) 304-2H (2)

6 0.99 31.7 394 316TiE 74 (2) 316TiE 73 (2) 316TiH 73 (2)

316TiE+R 73 (2)

316-2E (2) Nitro 50 (2) 5 0.94 30.1 390

304-2H (4)

304-1H (2)


0.87 27.8 390 (CT) 386 (Tr)

304-1H (2) 316-1E (2)

304-1H (2)

316-1H (2) 316TiH 73 (2)

3 0.71 22.7 384 Uranus S1N (2)

Nitro 50 (2)

316-1E (2)


304-1H (2) 316-1E (2) 316-1H (2) 316-2E (2)

2 0.57 18.2 381 316TiE 74 (2) 316TiE 73 (2) 316TiH 73 (2)

316TiE+R 73 (2)

316TiE 73 (2) Nitro 50 (2)

1 0.39 12.5 380 304-1H (2) 316-1E (2)

304-1H (2)


E= Cold Worked, H= Quenched, R= Tempered, Tension specimens have d= 2, gauge length= 12, total length = 33, Impact specimens are quasi standard Charpy-V size 2.5x10x55 and sub-size 3X4X27, Fracture specimens are CTJ10 type but with thickness reduced to 5, fatigue specimens are cylindrical with d=4, parallel gauge= 10, length= 66, Irradiation creep specimens are pressurized tubes external diameter = 6.55 and total length = 55 mm


The ELIXIR irradiation started with the Phénix’s 50th operation cycle on may 24, 1998. Following anomalies observed in the secondary circuit of reactor, this cycle was stopped at the end of September 1998 after 77.34 running days (EFPD), well before the planned 120 EFPD. Since then the rig has recived small doses. The rig is still in reactor but the uncertainty over the start date of 51st cycle remains. Up to now the damage in specimens loaded in the ELIXIR experiment has reached a max dose of about 7 dpa. Depending on the future of Phénix, the irradiation will either be continued to higher dose levels or the rig will be removed at the current dose level. CONCLUSIONS - Conditions of an irradiation experiment in BOR 60

reactor have been finalised. - High dose irradiation experiment started in Phénix has

reached 7 dpa, but due to uncertainties over the future operation schedule of the reactor it has been put on hold.

REFERENCES [1] Compte Rendu, Centrale Phénix, 28 June 1995,

Project CALIPH. [2] J-L. Bertin, C. Milet, “Expérience ELIXIR, dossier de

définition technique, éléments de surêté”, NT. SRMA 95-2148.

[3] F. Tavassoli, J-L. Séran and J-L. Bertin, SM 6.1 final

report, NT DECM/Dir 2000-83., December 2000.

TASK LEADER TAVASSOLI Farhad DEN/DMN CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 60 21 Fax : 33 1 69 08 80 70 E-mail : [email protected]


TTMS-1.1 Task Title : RAFM STEELS - IRRADIATION PERFORMANCE Neutron irradiation of RAFM steels to 30-35 dpa at 325°C INTRODUCTION The main objective of this proposal is to study the sensitivity to hardening/embrittlement of RAFM steels in relation to compositional effects at irradiation temperatures relevant for EBP structural materials. In particular, to assess the behaviour of RAFM steels data for high doses are needed at irradiation temperatures lower than 400°C, where materials could reach a high level of hardening and embrittlement,. In the period 95-99, different neutron irradiations of RAFM steels have been performed, that is, in HFR reactor at temperatures ranging from 250 to 450°C for a dose of 2.4 dpa and in Osiris reactor at 325°C for 5 dose levels from 1 to 9 dpa. To complete the available data, it is proposed for the period 2001-2002 to irradiate the precursor EU heat EUROFER97 as well as other RAFM steels and ODS alloys in the BOR60 reactor at 325°C up to a dose of about 30-35 dpa. These results will be compared to mechanical testing data of several RAFM steels previously obtained and will contribute to the final specification for the industrial heat of EU RAFM steel. 2000 ACTIVITIES Activities during 2000 were mainly devoted to the selection of materials, the machining of specimens and the preparation of the irradiation. This experiment was initially planned to be performed in the Osiris reactor, but due to some technical problems in the irradiation loop, it was decided to perform this experiment in the BOR60 reactor. The irradiation rig is sheared with other CEA materials project. MATERIALS Materials selected for this experiment are : a) Eurofer 97 steel (9Cr1W VTa) in the as-received condition, that is normalised and tempered (N&T) condition ; b) two experimental RAFM steels 9Cr2W VTa and 9Cr2W VTaB in the N&T condition; c) advanced ODS alloys, MA957, that is 13Cr-1Ti-0.3Mo-0.25Y2O3 obtained with two different fabrication routes to produce different microstructures, that is, fine grains and recrystallised microstructures.

IRRADIATION EXPERIMENT This one will be performed in BOR60 reactor of the Research Institute of Atomic Reactors (RIAR, Dimitrovgrad, Russia). The irradiation is planned to start on May 2001 and the completion is foreseen during 2002. The main characteristics of the experiment are : Irradiation temperature : 325°C. Dose : 30-35 dpa. Environment : Liquid Na. Specimens : Tensile and Charpy-V (subsize)

specimens and pressurised tubes.

Materials : EUROFER97, 9CrWTaV

RAFM steels and ODS alloys. After dismantling the capsule, post-irradiation examinations and testing will include tensile tests mainly performed at 20°C and 325°C. Charpy-V tests will be performed to obtain the transition impact energy curves (typically from -200°C to +200°C). Irradiation-creep data will be obtained by profilometry measurements on pressurised tubes of EUROFER97 and 9CrWTaV RAFM steels having 2 different stress levels. PIE of irradiated specimens are planned to be performed in RIAR’s hot cells. REPORT Progress Report TTMS 1.1, December 2000, A. Alamo, CEA/DMN/SRMA/01.07. TASK LEADER A. ALAMO DEN / DMN / SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 67 26 Fax : 33 1 69 08 61 75 Email : [email protected]


TTMS-2.1 SM-2-1

Task Title : RAFM STEELS - METALLURGICAL AND MECHANICAL

CHARACTERIZATION Metallurgical characterization of EUROFER 97 steel INTRODUCTION The main objective of this task is to characterise the metallurgical behaviour of the new EU heat EUROFER 97 in comparison with other RAFM steels, like F82H and experimental CrWTaV alloys, previously studied in the period 95-99. The main items developed in the frame of this task concerning the EUROFER 97 steel are: a) The metallurgical characterisation of products in the as-

received condition. b) The physical metallurgy (phase transformations,

precipitation reactions) during different thermal cycles. c) The effects on microstructure of thermal ageing

simulating in-service conditions (range 250-550°C, time up to 10000h).

2000 ACTIVITIES The task named SM2.1 corresponds to the programme 1999 and TTMS2.1 to the same programme developed on 2000. Due to the late delivery of the Eurofer 97 products (December 20, 1999) and the further time needed to plan, cut and machining the specimens (3 months), activities on the 1st trimester 2000 were oriented to finish some experiments started on 1999 concerning the physical metallurgy of experimental CrWTaV steels having a chemical composition close to Eurofer 97. In the second time, the metallurgical characterisation of different products of Eurofer 97 (item a) was carried out and completed on December 2000. On the other hand, the phase transformation behaviour is actually investigated (item b) to set the phase transformation diagrams (CCT, TTT diagrams) determined under different heating and cooling conditions. The thermal ageing treatments were also started on mid-2000 for temperatures in the range 250-550°C for 10000h and they will finish on November 2001.

CrWTaV STEELS: EFFECTS OF THE INITIAL METALLURGICAL CONDITION ON THE PHASE TRANSFORMATION TEMPERATURES The temperatures for starting and ending the transformations of ferrite/martensite → austenite on heating and austenite → martensite on cooling have been determined for different initial metallurgical states of the alloys LA12Ta and LA12TaLN [1]. The objective of this investigation was to determine the effects of the initial interstitial content in solid solution on the matrix on the phase transformation behaviour. For this purpose, six different initial conditions were considered, that is, a fully ferritic matrix, the as-quenched martensite and martensite obtained after tempering at 550, 650, 750 and 800°C. The different temperatures of tempering induce a variable fraction of precipitated carbides and so, different interstitial contents in solid solution in the matrix. The chemical composition of both considered alloys is quite similar except for the N content as shown in table 1.

Table1 : Chemical composition of the studied steels (weight %)

C Si Mn Cr V W Ta N

LA12TaLN 0.17 0.02 0.74 9.1 0.25 0.77 0.1 0.004

LA12TA 0.16 0.03 0.80 9.8 0.27 0.85 0.1 0.042

Figure 1 shows the temperatures to reach 50% of austenite transformed during heating in the same conditions for the RAFM steels, LA12TaLN and LA12Ta. The lower temperature is obtained for the as-quenched martensite and the higher ones for the fully ferritic matrix. 9Cr-1Mo (EM10) conventional steel displays the same behaviour. The same trend is obtained for the temperatures corresponding to start (As) and to finish (Af) the austenite transformation. The results obtained for the transformation temperatures As and Af could be related, in first approximation, to the variation of the interstitial content (mainly C) in solid solution, according to the initial metallurgical state. Taking into account the "extreme" conditions of our samples, this content is nearly 1600 - 1700 ppm for the as-quenched martensitic state and nearly 100 ppm for the fully ferritic state.


The martensite start temperatures Ms on fast cooling have been measured as well, for the same initial states and after a short incursion in the austenite phase field. It was also found an increase of these values from the as-quenched martensite initial state (320°C) to the fully ferritic matrix initial state (380°C), with intermediate values for the temperatures corresponding to the tempered martensite initial states. Thus, the evolution of the phase transformation temperatures on-heating and on-cooling (As, Af and Ms) is correlated to the fraction of carbon in solid solution in the initial state, the temperatures increases with the decreasing carbon content. No clear nitrogen influence was found on transformation temperatures, for the studied composition range.

880

900

920

940

960

980

FerriteAs-quenched

martensite

800750700650600550

LA12TaLN

LA12Ta

9Cr1Mo (EM10)

A5

0%T

em

pe

ratu

re (

°C)

Tempering Temperature (°C)

Figure 1 : Temperatures to reach 50% of austenite (A50%)

during heating for different steels METALLURGICAL CHARACTERIZATION OF EUROFER 97 PRODUCTS Eurofer 97 is the first European RAFM steel produced at industrial scale. Several heats have been manufactured by Böhler steel maker (Austria). The metallurgical characterization of Eurofer97 products in the as-received conditions consisted on the determination of chemical composition, grain size distribution, inclusion content, optical microscopy examinations and hardness measurements. This is the first step in the metallurgical characterization of these products in view to the qualification of their mechanical properties, thermal ageing and irradiation behaviour. Plates of four different thicknesses and a bar of Eurofer97 have been assigned to CEA. They were delivered in the normalized and tempered condition (N&T). According to the supplier, heat treatments of plates consisted in normalisation for 27-31 min. at 980 °C followed by air cooling and tempering for 1h 30 min. at 760 °C. In the case of the bar, the austenitisation treatment was performed for 1h 51m. at 980 °C, followed by air cooling and tempering for 3h. 42 m. at 740 °C.

The metallurgical characterization was carried out on one plate of each thickness. For the Ø100mm bar, examinations were performed at different points from the centre up to the external surface. The identification and dimensions of the characterized products are summarised on Table 2.

Table 2 : References and dimensions of Eurofer97 products

Product Heat Reference CEA

Reference Thickness

(mm)

Plates E83698 2/3 930 1.5

Plates E83698 66/5 931 8

Plates E83698 4/9 932 14

Plates E83694 1/6 933 25

Bar E83699 M7CEA-S 928 100

(diameter)

Chemical composition Chemical compositions are summarized in Table 3. All of the products exhibited a similar chemical composition; nevertheless, some differences were found with respect to the specified contents. Values out of specification are given with a grey background. The main differences detected were: 1) P, Ta, Ni, Nb and Co were found to be systematically

out of specification and the same was true for Cu in five of the six products.

2) Mo, Si and W displayed deviations in some cases. 3) Sb exhibited an elevated concentration for one of the

products. Inclusion content Inclusion content was determined according to the standard ASTM E45 by comparison with standard charts (Jernkontoret charts). Determinations were performed on polished specimens (without chemical etching) in longitudinal cross sections parallel to the rolling direction. For each sample, 20 fields were observed corresponding to a surface of 0.5 mm2 for each field and the total examined area is 10 mm2. For the 100 mm diameter bar, five samples were taken from the external surface up to the centre. Detected particles are classified by similarities in the morphology (size, shape and distribution) in the following types: A (sulfides), B (alumina), C (silicates) and D (oxides). The inclusions detected are mainly of type D (globular oxide type, thin series) and, in minor amounts, A, B, and C. The specified average rating number must be <1 in all cases. The plates of 8, 14 and 25 mm thickness were found to be slightly out of specification (values ranging from 1.0 to 1.4) for the type D, whereas the bar displayed values within the specified range.


Table 3 : Chemical composition of the different Eurofer97 products

Alloy Heat

Reference

Specification EUROFER 97

EUROFER E83698

2/3

EUROFER E83698

66/5

EUROFER E83698

4/9

EUROFER E83694

1/6

EUROFER E83699

M/7CEA S(1)

EUROFER E83699

M/7CEA S(2)

Cr 8,5 - 9,5 [ 9 ] 9,04 8,96 9,01 8,85 8,71 8,68

C 0,09 - 0,12 [ 0,11 ] 0,109 0,117 0,109 0,092 0,115 0,116

Mn 0,20 - 0,60 [ 0,40 ] 0,49 0,48 0,5 0,44 0,41 0,41

P < 0,005 0,011 0,011 0,012 0,013 0,013 0,013

S < 0,005 0,003 0,003 0,003 <0,003 0,003 0,003

V 0, 15 - 0,25 0,19 0,18 0,18 0,18 0,17 0,18

B < 0,001 <0,001 <0,001 <0,001 <0,001 <0,001 <0,001

N2 0,015 - 0,045

[ 0,030 ] 0,0226 0,0219 0,020 0,0180 0,019 0,019

A)

O2 < 0,01 0,0011 0,0011 0,0005 0,0008 0,0011 0,0009

W 1,0 - 1,2 [ 1,1 ] 1,03 1,04 1,04 1,08 1,26 1,26

Ta 0,06 - 0,09 0,14 0,15 0,10 0,15 0,12 0,12 B)

Ti < 0,01 (100 ppm) 0,002 0,002 0,003 0,001 0,005 0,005

Nb [< 0,001 (10 ppm)] <0,002 <0,002 <0,002 <0,002 <0,002 <0,002

Mo [< 0,005 (50 ppm)] <0,01 <0,01 <0,002 <0,01 <0,002 <0,002

Ni [< 0,005 (50 ppm)] 0,06 0,06 0,07 0,07 0,04 0,04

Cu [< 0,005 (50 ppm)] 0,04 ** 0,04 ** 0,003 0,04 ** 0,02 ** 0,02 **

Al [< 0,01 (100 ppm)] 0,004 0,003 0,003 0,002 0,004 0,004

Si <0,05 (500 ppm) 0,03 0,03 0,04 0,04 0,06 0,06

Co [< 0,005 (50 ppm)] 0,02 0,02 0,010 0,02 0,008 0,008

Sn *) <0,002 <0,002 0,004 <0,002 0,002 0,002

As *) <0,002 <0,002 <0,002 <0,002 <0,002 <0,002

Sb *) 0,010 0,006 <0,002 0,003 <0,002 <0,002

C)

Zr *) <0,0005 <0,0005 <0,0005 <0,0005 <0,0005 <0,0005

Legend: A) basic composition, B) varied substitution elements C) radiological undesired trap elements, [ ] target values, *) As + Sn + Sb + Zr < 100 ppm; (1) center of the bar, (2) external surface of the bar, ** : analysis to be confirmed.

Hardness measurements Vickers hardness measurements were performed on the complete set of as-received products. The applied load was 10 Kg. For the plate products, measurements were done on the longitudinal and transverse cross-sections. Each value represents the average of five indentations. The average values of hardness for these products are presented in Table 4. For the 100 mm-diameter bar, five cross sections were obtained from the external surface down to the centre. Similar values (HV10 : 218-220) are obtained for all the axial and radial cross-sections showing a good homogeneity of the structure.

Table 4 : Average hardness values for plate products

Thickness 25 14 8 1.4

Longitudinal Average hardness 218 215 219 216

Transverse Average hardness 222 216 215 212

Optical microscopy observations Optical microscopy observations were performed on longitudinal cross sections, parallel to the rolling direction. A homogeneous martensitic morphology was observed for the whole of the samples. Plates of 1.4 mm thickness exhibited a more refined structure of the martensitic matrix. No traces of ä-ferrite were detected. Prior austenite grain size A solution of saturated picric acid was used to detect prior austenitic grain boundaries [2]. Approximately 600 grains were sampled for each specimen. Binary grain boundary images were obtained from numerical ones. Two methods were used to determine the prior austenite grain size: a) Determination of the ASTM G standard index for grain

size (ASTM standard E112) by counting the average number of grains by unit of surface area. The specified grain size index must be G ≥ 9.


b) A quantitative statistical determination of the size distribution and the average size value using image analysis techniques.

For the first method, the ASTM G values are deduced from the “m” values, where m corresponds to the number of grains by mm2 in the observed surface. In the case of equiaxed grains, the standard ASTM E112 gives the G index value as a function of m (Table III of the ASTM standard). In the case of second method, the mean diameter prior austenite grains are determined from the size distribution curves obtained from measuring about 600 grains for each sample. A good agreement is found between both methods as shown in table 5. In general, a fine structure characterised by a prior austenite grain size in the range 13-15 microns was found for all products en good agreement with the specification requirements. In the case of plates 1.4mm thick a more fine grain is observed with a mean value of about 10-11 microns.

Table 5 : Comparison of mean grain diameters obtained by both methods

Mean diameter (µm)

Reference Thickness or

diameter (in mm)

ASTM standards

Image analysis

G ASTM index

2/3 1.4 10.18 11.20 10

66/5 8 13.32 14.56 9

4/9 14 13.77 14.84 9

1/6 25 14.88 15.27 9

M7CEA-S (c)* ∅100 14.51 15.27 9

M7CEA-S (s)* ∅100 14.18 14.77 9

∗ M7CEA – S (c) and (s) correspond to the core and the surface of the bar respectively CONCLUSIONS - Investigations performed on CrWTaV experimental

alloys have shown that phase transformation temperatures depend on the initial metallurgical conditions and in particular on the carbon content in solid solution.

- The different control tests performed on Eurofer97

products have shown homogeneous properties for all of them.

• All products exhibited nearly the same chemical

composition. Nevertheless, some differences were detected with respect to the specified range. In particular, some elements (P, Ta, Ni, Nb, Co) were systematically found to be at higher contents.

• Optical microscopy observations showed a fully martensitic structure in all cases; no traces of delta ferrite were detected.

• Prior austenite grain size was determined using the

ASTM standard procedure and quantitative image analysis. The mean grain diameter ranged from 13 to 15 microns and the obtained G value was 9 for the different products except for the plates of 1.4 mm, which exhibited values of 10-11 microns and G = 10 respectively.

• Inclusion content was determined according to the

ASTM standard by comparison with standard charts. Inclusions detected are mainly of type D and, in minor amounts, A, B, and C. The averaged inclusion rating numbers ranged from 1.1 to 1.4 in the case of 14 and 25 mm-thick plates.

• Hardness measurements displayed similar values (about

220 HV10) for different products and differently oriented cross sections.

REPORTS [1] A. Danon, A. Alamo, «Effects of the initial

metallurgical condition on the phase transformation temperatures of FeCrWTaV steels », Progress Report SM2.1, NT SRMA 00-2399, December 2000.

[2] A. Danon, S. Urvoy, A. Alamo, « Metallurgical

characterisation of as-received EUROFER 97 products», Progress Report TTMS2.1, NT SRMA 01-2418.

TASK LEADER A. ALAMO DMN / SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 67 26 Fax : 33 1 69 08 61 75 E-mail : [email protected]


TTMS-2.2 Task Title : RAFM STEELS - METALLURGICAL AND MECHANICAL

CHARACTERIZATION Mechanical characterization of EUROFER 97 :

Tensile and impact properties INTRODUCTION EUROFER 97 is the first European RAFM steel produced at industrial scale. The main objective of this task is provide data concerning the tensile and impact properties of EUROFER 97 steel in different metallurgical conditions (as-received and thermal aged states). The main items that will be developed in the frame of this task are: a) The characterization of tensile and impact properties of

Eurofer 97 in the as-received condition. b) The determination of tensile and impact properties of

Eurofer 97 after ageing performed in the range 250-550°C up to 10000 hours.

c) Comparison of the mechanical behaviour evolution of

Eurofer 97 and other RAFM steels in particular after thermal ageing.

2000 ACTIVITIES During 2000 activities were focused on the determination of mechanical properties of Eurofer 97 in the as-received condition (item a). For this purpose, tensile and Charpy specimens were machined on plates of 14mm thick along different directions. All mechanical tests were completed on December 2000. The thermal ageing treatments of specimens for mechanical tests were also started on mid-2000 at temperatures in the range 250-550°C for times up to 10000h. They will finish on November 2001. On the other hand, following to the requirements of the European Task Leader a complete compilation of CEA data concerning the mechanical properties of the F82H steels was prepared and presented in the “Workshop on Data Base Evaluation of RAFM steels” held at Brasimone, Italy, on November 27-28, 2000. This compilation, object of the report referenced later [1], deals with the tensile and impact properties of F82H in the as-received and thermal aged conditions as well as the mechanical behaviour after neutron irradiations performed in HFR and Osiris reactors in the range 250-450°C with doses up to 3.4 dpa.

MECHANICAL PROPERTIES OF EUROFER 97 IN THE AS RECEIVED CONDITION The mechanical characterization was mainly conducted on plates of 14 mm thick reference 4/9 issued from the heat n° E83698 manufactured by Böhler (Austria) [2]. Plates were distributed by FZK. According to the supplier, heat treatments of plates consisted in normalisation for 27-31 min. at 980 °C followed by air cooling and tempering for 1h 30 min. at 760 °C. Details of the metallurgical characterization (grain size, inclusions, hardness) are given in the report TTMS2.1. Tensile properties Tensile properties were performed on cylindrical specimens of 2 mm of diameter and 12 mm of gauge length. The strain rate was 7. 10-4 s-1 and the test temperature ranged from 20 to 700°C. Compared to the rolling direction, tensile specimens were machined along the parallel (longitudinal) and the perpendicular (transverse) directions of plates. No differences were observed of tensile properties for longitudinal and transverse specimens, which mean that Eurofer 97 plates present quite isotropic properties. Figure 1 shows the evolution of tensile strength, elongations and reduction in area values as a function of the test temperature for specimens obtained along the longitudinal direction. Tensile properties of Eurofer 97 are very close to that of F82H plates of the same thickness (15mm). Tensile strength of Eurofer is slightly higher than F82H in the range 20-450°C, but the inverse situation is observed for higher test temperatures. Concerning the ductility, both materials display nearly the same levels of elongation and reduction in area values. Impact properties Impact tests were performed on specimens having different dimensions, that is: - Charpy V subsize (KLST) specimens of 3mm thick,

4mm wide and 27mm long with a cross-section are of 0.09 cm2.

- Charpy V standard specimens of 10mm thick, 10mm

wide and 55mm long with a cross-section are of 0.8 cm2.

- Charpy V quasi-standard specimens of 3.5mm thick,

10mm wide and 55mm long with a cross-section are of 0.28 cm2.


In the case of Charpy V subsize, specimens were machined with LT and TL orientations. For the other types (standard and quasi-standard), samples were obtained only with the TL orientation.

In general, tests were conducted over the temperature range –200 to +100°C to produce the full energy transition curves. The ductile-brittle transition temperature (DBTT) defined as the 50%ductile-50%cleavage fracture mode corresponds to the midway between the upper lower shelves of the transition curves. The accuracy of DBTT values is about ± 10°C.

100

200

300

400

500

600

700

0 100 200 300 400 500 600 700 800

Eurofer 97

F82H

0.2

% P

roof

Str

ess

(MP

a)

Temperature (°C)

As-received conditionLongitudinal Direction

100

200

300

400

500

600

700

0 100 200 300 400 500 600 700 800

Eurofer 97

F82H

U.T

.S. (M

Pa)

Temperature (°C)


0

5

10

15

20

25

30

0 100 200 300 400 500 600 700 800

Eurofer 97

Eurofer 97

F82H

F82H

Elo

ngat

ion

(%)

Temperature (°C)


Total Elongation

Uniform elongation

75

80

85

90

95

100

0 100 200 300 400 500 600 700 800

Eurofer 97

F82H

Red

uct

ion in A

rea

(%)

Temperature (°C)


Figure 1 : Comparison of tensile properties of the EUROFER 97 and F82H steels in the as-received condition, determined in the longitudinal direction


Figure 2 presents the transition curves obtained with Charpy V specimens machined with the LT and TL orientation. As shown, impact properties are quite similar in both directions. The Upper Shelf Energy (USE) level is slightly lower for TL specimens. DBTT reach very low values for both directions.

0

2

4

6

8

10

12

-200 -150 -100 -50 0 50 100 150 200

LT orientation

TL orientation

Ene

rgy

( J)

Temperature (°C)

EUROFER steelSpecimens : 3 x 4 x 27 mm

3

USE (J)DBTT (°C)Spec. Orientation

9.0- 99LT

8.4- 97TL

Figure 2 : Impact properties of EUROFER 97 in the as-received condition determined for Charpy V

subsize specimens with LT and TL orientations. Figure 3 and 4 compare the energy transition curves of Eurofer 97 and F82H steels determined respectively with subsize (KLST) and standard Charpy V specimens with TL orientation. In both cases, DBTT values of Eurofer 97 are lower than those obtained for F82H. But, F82H exhibits higher USE levels. The same situation is also observed with quasi-standard specimens.

0

2

4

6

8

10

12

-200 -150 -100 -50 0 50 100 150 200

Eurofer 97

F82H

Ene

rgy

( J)

Temperature (°C)

EUROFER

Specimens : 3 x 4 x 27 mm3

TL-orientation

F82H

USE (J)DBTT (°C)Material

8.4- 97Eurofer 97

9.5- 59F82H

Figure 3 : Comparison of impact properties of EUROFER 97 and F82H steels

(Charpy V subsize specimens with TL orientation)

0

50

100

150

200

250

300

350

-200 -100 0 100 200 300 400

Ene

rgy

( J)

Temperature (°C)

EUROFER

Specimens : 10 x 10 x 55 mm3

TL orientation

F82H

USE (J)DBTT (°C)Material

220- 62Eurofer 97

290- 18F82H

Figure 4 : Comparison of impact properties

of EUROFER 97 and F82H steels (Charpy V standard specimens with TL orientation)

CONCLUSIONS The tensile and impact properties were determined for plates of Eurofer 97 steel in the as-received condition. This material presents quite isotropic mechanical properties, similar values are obtained along the longitudinal and transverse directions of plates. Compared to F82H steel, Eurofer 97 exhibit similar tensile strength and ductility values, lower DBTT, but also lower USE level. During 2001, the evolution of mechanical properties of aged specimens of Eurofer 97 will be studied for ageing treatments performed for short times (2000h) at different ageing temperatures. REPORTS [1] A. Alamo, “Mechanical properties of F82H steel in

the as-received, aged and irradiated conditions. Compilation of CEA-data”, CEA report NT SRMA 00-2395, December 2000.

[2] A. Alamo, P. Wident, Y. Tournier, S. Urvoy,

“Mechanical properties of EUROFER 97 steel in the as-received condition”, CEA report NT SRMA 01-2419.


TASK LEADER A. ALAMO DMN/SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 67 26 Fax : 33 1 69 08 61 75 E-mail : [email protected]


TTMS-2.3.1 Task Title : RAFM STEELS - METALLURGICAL AND MECHANICAL

CHARACTERIZATION Mechanical properties of diffusion bonded welds

(RAFM/RAFM HIP joint) INTRODUCTION The objective of this task is to supply a data base on mechanical properties of the Eurofer/Eurofer junction at different temperatures. In 1999, tensile properties and impact properties have been determined. This year, low cycle fatigue properties at room temperature and creep properties at 550 °C have been measured. 2000 ACTIVITIES MATERIAL Böhler has elaborated the reduced activation ferritic martensitic steel or Eurofer 97. The chemical composition (Heat N° E83699) is given in table 1.

Table 1 : Eurofer 97 composition

wt % C Si Mn P S Cr

Value 0.12 0.06 0.42 0.004 0.003 8.87

wt % Mo Ni V W Cu Co

Value <0.001 0.0075 0.19 1.10 0.021 0.005

wt % Ti Al Nb B N Ta

Value 0.008 0.008 <0.001 <0.0005 0.018 0.14

wt % O As Sn Sb Zr

Value 0.0013 <0.005 <0.005 <0.005 <0.005

MANUFACTURING OF JUNCTION The conditions of manufacturing have been presented in [1] [2]. The different thermals treatments are just recalled here. The HIP treatment is the following : - 2hrs/2hrs/2hrs -1100°C - 1000 bars A Post Bond HIP Treatment (PBHT) was applied to restore the microstructural and mechanical properties of Eurofer : - An austenitisation treatment : 1hour at 950°C followed

by an air cooling. - A tempering treatment : 1 hour at 750°C followed by an

air cooling.

PROPERTIES OF JUNCTION Metallurgical examinations Very often, junctions elaborated by diffusion bonding present oxides precipitates along the bond. A SEM observation has been realized on the Eurofer/Eurofer junction elaborated at 1100°C after having applied a smooth etching (Villella) to the specimen. The process route (surface preparation, degassing conditions) and the HIP conditions have allowed to obtain a excellent quality of the junction since no oxides precipitates were detected along the joint (Figure 1).

Figure 1 : SEM observation of the Eurofer/Eurofer junction

Mechanical properties Summary of results obtained in 1999 - Tensile properties:

Figure 2 : Tensile properties of the Eurofer/Eurofer junction

Junction

50 µm


- Impact properties:

Figure 3 : Impact properties (KCU) of the Eurofer/Eurofer junction and bulk material

The mechanical properties of the joints were excellent, the failure never occurred at the interface but was always located in the bulk material. The impact properties of the bulk material and of the joint were similar. Low Cycle Fatigue properties The low cycle fatigue tests (LCF) have been carried out at room temperature. The joint was in the centre of the gauge length and oriented perpendicular to the loading axis. The specimens were tested under imposed strain. A triangular signal was used with a strain rate of about 2. 10-3 s-1. The test was stopped for a load drop of about 25%. Results are displayed in table 2 and on figures 4 and 5.

Table 2 : Results of LCF tests

Strain ∆σ∆σ/2 (MPa)

σσo (MPa) (1/4cycle)

Nf 25

Junction ± 0.2 % 411 440 No

failure 86000

Junction ± 0.5 % 512 647 4000

Junction ± 1 % 590 694 563

Bulk material

± 0.5 % 515 634 2818

Figure 4 : Curves of minimum and maximum stresses bonded specimens

Figure 5 : Curves of minimum and maximum stresses at ± 0.5 % on bulk material specimen

and on bonded specimen For the imposed strain of about ± 0.5 %, a specimen taken in the junction and one specimen taken in the bulk material block have been tested under the same loading conditions to check that they have the same behaviour. The number of cycles at failure were different but the evolution of the tension and compression stresses was similar. This difference in Nf 25 is probably due to a scatter in the material properties. For the test at ± 0.2 %, no failure has been detected at 8600 cycles. For the 3 tests on junction no damage has occurred in the joint. A cross section has been realized on the bonded specimen (figure 6) tested at ± 0.5 %. A large failure appears, but not along the joint.

Figure 6 : Cross section of the specimen tested at ± 0.5% Creep properties at 550 °C The creep tests of the Eurofer/Eurofer junction have been carried out at 550°C under air. The loads have been chosen in order to obtain both small time and large time tests. The creep tests have been conducted under constant load. The results are displayed on figure 7.

0

50

100

150

200

250

300

-200 -100 0 100 200 300 400

Temperature

Impa

ct t

ough

ness

(J/

cm2)

Bulk Material

Junction


Figure 7 : Results of creep tests at 550°C As for the other tests, fracture has never occurred in the joint. We may thus consider the behaviour of these specimens as relevant of the behaviour of the bulk material. CONCLUSION The mechanical properties of the Eurofer/Eurofer junction are excellent. As for tensile tests, impact tests carried out in 1999, low cycle fatigue and creep tests realized in 2000, the fracture never occurred in the joint, but always in the bulk material. In 2001, low cycle fatigue tests at 300 and 550°C and creep tests at 450 °C will be finished. Fracture toughness tests at both room temperature and at temperature will be performed to complete the mechanical properties on the Eurofer/Eurofer Junction. REFERENCES [1] Mechanical properties of RAFM/RAFM HIP Joint..

Bedel L., Rigal E., Lemoine P., SM 2-3 Note technique DEM N° 99/87.

[2] Development of martensitic steel / martensitic steel

HIP joints with improved mechanical properties. SM 4-5. E. Rigal, L. Bedel. Note technique DEM N° 91/99.

REPORT AND PUBLICATION [R1] Mechanical properties of A RAFM/RAFM HIP Joint

(2000). L. Bedel, P. Lemoine TTMS 2-3. Note Technique DEM 2000-105,

Décembre 2000 TASK LEADER Laurent BEDEL DRT/DTEN/SMP/LS2M CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 57 20 Fax : 33 4 76 88 58 91 E-mail : [email protected]

250 MPa

220 MPa

235 MPa


TTMS-2.3.2 Task Title : RAFM STEELS - METALLURGICAL AND MECHANICAL

CHARACTERIZATION Mechanical properties of weldments INTRODUCTION The development of fast induced radioactivity decay alloys has been incorporated into the goals of the IEA collaborative program in order to reduce the quantity of highly radioactive materials from system components. Concerning the innermost parts of a Fusion reactor, the low activation martensitic (LAM) steels and reduced-activation ferritic/martensitic (RAFM) steels are considered as primary candidate alloys as first wall and breeder blanket structural materials because of high irradiation resistance, especially the resistance to the phase instability and the resistance to the shift of the ductile-brittle transition temperature (DBTT) during or after irradiation. Several RAFM steels have been developed specially for this application in the frame of the Long Term Fusion international collaborative test program. Among these materials, the F82H steel developed by JAERI in the mid nineties and the more recent Eurofer 97 proposed by FZK and Böhler are considered to be specially viable candidates. Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding and laser welding processes are envisaged for the fabrication of structural parts of relevant breeding blanket modules. Mechanical tests are carried out on F82H welded joints in the frame of the task SM-2.3.1. The present study is the continuation of the task SM-2.3.1 by carrying out mechanical tests on Eurofer 97 weldments. During the year 2000, welded joints of 24 mm of thickness were produced. The specimens were sampled in these weldments for tensile and impact test characterization.

2000 ACTIVITIES PRODUCTION OF WELDED JOINTS FOR CHARACTERIZATION Two assemblies of two Eurofer 97 plates of 24 mm of thickness and of 200 mm of length each have been welded using the TIG process with matching filler metal. Each assembly was made of two plates (Figure 1). The operating conditions have been selected according to the results of the task TTMS-4.3 (Table 1). They were very close to those applied on the F82H in the frame of the task SM-2.3.1.

Figure 1 : TIG welding of Eurofer 97 plates of 24 mm of thickness

Table 1 : Operating conditions used to process 24 mm TIG welded joints

of Eurofer 97 for mechanical characterization

mat

eria

l

nb p

asse

s

thic

k (m

m)

clam

ping

U (

V)

I (A

)

wel

ding

spe

ed (

cm/m

in)

wir

e sp

eed

(m/m

in)

wir

e di

amet

er (

mm

)

scan

ning

wid

th (

mm

)

uppe

r sh

ield

ing

(l/m

in)

back

shi

eldi

ng (

l/min

)

back

side

shi

eldi

ng (

l/min

)

Eurofer 8 24 welded 10 150-230 4 0,7-2 1 3,5-9,5 15 20 7,5


The weldments have followed radiographic control indicating that they are free of internal flaws. The post welding heat treatment conditions were 1 h at 710°C. The treatment has lowed the hardness down to about 320 Hv1 in the seam. This value is higher than the base metal hardness (about 220 Hv1) indicating that the duration of the treatment should be longer. SAMPLING OF TESTS SPECIMENS The test matrix is given in Table 2. As in the task SM-2.3.1, 24 impact tests have been planned (12 specimens with the notch centred on the melted zone and 12 with the notch centred on the heat affected zone). Two types of tensile specimens have been planned : a cylindrical design for the conventional tensile tests at various temperatures and a plate design dedicated to specific tensile tests which exploitations give a precise strains distribution over the specimen.The plate specimen design is displayed in Figure 2. The standard V-notch design has been used for the impact specimens. The specimens sampling is displayed in figure 3. The rough shapes have been etched in order to centre precisely the specimens either on the melted zone or on the heat affected zone during the final machining. Each specimen has been engraved at both ends.

Figure 2 : Plate tensile specimen design

Table 2 : Tensile and impact tests matrix

Material Process Thick. Plane 1.5 mm

(mm) Centred MZ Centred HAZCentred MZ

crossCentred MZ

longi.BM

Centred MZ cross

B 3 T 32 B 3 T 44 B 1 T 57 B 1 L 59 B 1 L 64 B 2 T 55B 3 T 33 B 3 T 45 B 1 T 58 B 1 L 60 B1 L 65 B 2 T 56B 3 T 34 B 3 T 46 B 1 L 61 B 1 L 66B 3 T 35 B 3 T 47 B 1 L 62B 3 T 36 B 3 T 48 B 1 L 63B 3 T 37 B 3 T 49B 3 T 38 B 3 T 50B 3 T 39 B 3 T 51B 3 T 40 B 3 T 52B 3 T 41 B 3 T 53B 3 T 42 B 3 T 54B 3 T 43 B 3 T 55

Eurofer 97 TIG filler 24

Impact tests

Standard Cylindrical M7

Tensile (20 ; 400 ; 500 ; 600°C)

165 : Zone exploitable

F0B0306

A-A

A

A

Eprouvettes de résilience EUROFER 97

11,511,5

60

Centré en ZF

Centré en ZAT

Cordon

B3T32

B3T33

B3T34

B3T35

B3T36

B3T37

B3T38

B3T39

B3T40

B3T41

B3T42

B3T43

B3T44

B3T45

B3T46

B3T47

B3T48

B3T49

B3T50

B3T51

B3T52

B3T53

B3T54

B3T55

* 2

28

* 1

8

B1L64 B1L65 B1L66

F0B0304

170 : Zone exploitable

* 2* 2* 2

28 28

2

8

A

A

A-A

8

Eprouvettes de traction EUROFER 97

Cordon* 1

B1L59

B1L60

B1L61

B1L62

B1L63B1T57

B1T58

22

61

28

B

B

2

B-BB2T55

B2T56

Figure 3 : Design of the specimens sampling


CONCLUSIONS Sound TIG welded joints of 24-mm plates of Eurofer 97 have been processed with matching filler metal. After a post welding heat treatment, the weldments have been sampled. The rough shapes of the specimens have been etched in order to reveal the seam and to provide a precise sampling. The exploitation of the tests is in progress. The test specimens are now available for the next step. REPORTS AND PUBLICATIONS A. FONTES “Eurofer TIG weldments and sampling for mechanical characterization” CEA report DECS/STA/LMS/01-RT4006.

TASK LEADER André FONTES DRT/DECS/STA/LMS CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 37 21 Fax : 33 1 69 08 75 97 E-mail : [email protected]


TTMS-2.4 Task Title : RAFM STEELS - METALLURGICAL AND MECHANICAL

CHARACTERIZATION Mechanical properties of powder HIP Eurofer INTRODUCTION The fabrication of some DEMO blanket piecesare foreseen by Powder metallurgy which offers several advantages compared to forged material. In parallel with task TTM.S 4.1 which optimise the Eurofer powder and the HIP process the main objective of this task is to determine the mechanical properties of the powder HIPped RAFM material in relevant Test Blanket Module operating conditions. Mechanical properties such as tensile , impact tests, creep at 550°C and fatigue properties at same temperature will be determined. The influence of specific powder metallurgy microstructure and defects (powder boundary and PPB) on tensile ; impact results and creep behaviour will be analysed in relationship with the optimisation of elaboration parameters in task TTMS 4.1. Optimised Eurofer HIPed material specified in task TTMS 4 will be used for that task. 2000 ACTIVITIES The following consolidation conditions were selected to produce Eurofer powder HIPed material used in the mechanical test program temperature [1040°C; time 4 hours, pressure 140MPa]. The microstructure of tempered material (2h at 750°C) after HIP at 1040°C and post austenization at 980°C show martensite laths with carbide precipitation on figure 1.

Figure 1 : Microstructure of Eurofer Powder HIPed

at 1040°C/ tempered for 2h at 750°C (magnification x500)

The tempering effect on hardness is illustrated on figure 2 and it appears that the stabilization of the microstructure is only obtained between one and two hours.

Hardness Results

0

50

100

150

200

250

300

350

400

450

0 0,5 1 1,5 2 2,5 3

tempering time (h) at 750°C

hard

ness

( H

V30

)

P2000-1040°C

PS*

Figure 2 : Hardness value versus tempering time treatment ((PS* : Post Treatment 1h at 950°C /WQ )+ 2h at 750°C)

The tensile Charpy V and fatigue specimens were machined parallel to the canister axis but HIPed material does not present anisotropic properties. Tensile and Impact tests were performed at 20, 300, 550°C under air whereas fatigue tests were performed at room temperature. The tensile properties show monotonic decrease of strength (yield and Ultimate) and uniform elongation with temperature and total elongation increase with temperature.

Table 1 : Tensile results at different temperature on Eurofer HIPed material

Test tempe-rature

Yield stress 0.2% MPa

U.T.S MPa

Uniform Elongation

%

Total Elongation

%

Reduction of area %

20* 561* 668* 5,7* 14,1* 68*

20 489 617 6.4 15.5 76

300 445 516 3.7 15.5 75

550 353 369 1.5 21.8 76

*As HIPed +2h at 750°C (whitout PNT)


Due to the limited volume of powder material only a few impact tests were performed at both 20 and 300°C. The room temperature results reveal high values : 387J/cm2 and the impact toughness decreases sligtly with temperature (308 J/cm2 at 300°C). But other tests are in progress. The fatigue tests were performed at 0.2 and 0.4 % deformation amplitude with R=-1 and the table 2 summarizes the results. The life time at the lowest strain amplitude is between 6610 and 8800cycles and is over 30000 cycle (specimen no broken) at the upper strain amplitude. These results seems too be acceptable one for powder material behaviour.

Table 2 : Summary of fatigue results at room temperature

Total strain amplitude

(%)

True Stress (MPa)

Life time Nr( cycle)

Elastic Modulus E

(Gpa) Remark

0.4 426 6610 193

0.4 421 8860 187

0.2 385 >150000 197 No Broken

0.2 378 >30000 193 No Broken

The tensile, impact and fatigue specimen were observed by SEM. CONCLUSION The first metallurgical and mechanical tests reveals acceptable tensile and impact properties of Eurofer powder HIPed material . Further tests are necessary to confirm the results.

REPORTS AND PUBLICATIONS Note technique DEM N° 2001/28 Mechanical properties of Eurofer HIP powder steel Eurofer J.M. GENTZBITTEL, I. CHU TASK LEADER Jean Marie GENTZBITTEL DTEN/SMP/LS2M CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 95 01 Fax : 33 4 76 88 58 91 E-mail : [email protected]


TTMS-4.1 Task Title : RAFM STEELS - QUALIFICATION FABRICATION PROCESSES Eurofer powder HIP development and characterization INTRODUCTION Eurofer is the european Reduced Activation Ferritic/Martensitic steel candidate material for different nuclear applications and especially for DEMO blanket materials. The production of solid alloy by compacting Eurofer powder metal offer the following advantages: - production near net shape parts without melting, this

avoids the severe segregation associated with melting and permits complex design,

- reduction of the cost by a strong reduction of the

machining and the number of welds. The evaluation and the development of RAFM powder HIPed steel is the first step for estimate the potentiality of this technique followed by the mock up fabrication. The final objective of the task TTMS 4.1 is to develop the Eurofer powder HIP process in order to obtained a Reduced Activation Ferritic/Martensitic solid HIPed steel with similar mechanical properties to that of forged material. The technical steps are: - evaluate and optimise the powder characteristics

(chemical composition, granulometry, homogeneity, cleanless….),

- optimise the powder compaction process (powder and

canister preparation, temperature, time, pression of HIP cycle) in relation with the microstructure and mechanical properties of the Eurofer HIPed material,

- provide powder HIPed material for mechanical

properties measurements (task 2.4 and irradiation testing (ECN Petten)).

2000 ACTIVITIES POWDER FABRICATION A new batch of Eurofer powder has been atomized under argon from a forged Eurofer bar (heat E83699).and the powder has been characterized in terms of granulometry and chemical composition. The powders have been screened and sorted into 3 size classes : large particles (over 250 µm), medium particles (45µm < d < 250 µm) and small particles (d < 45µm).

The chemical composition of medium size and small size powde is comparable to the of the cast material, except for the oxygen and nitrogen content. There is no loss of major elements but one recommendation would be to have an initial cast alloy with major element in the upper range of composition. The higher oxygen content of small particles can be explained by their higher specific surface. We keep for this work only medium particles in order to limit the O content and to control the powder filling. The particles present usually a spherical shape except some larges particles which exhibit some satellites due to collision during atomisation (figure 1). The metallographic observations of particles show a homogeneous structure but no local chemical analysis was performed on the powder. A previous atomisation (1999) has been done on the first Eurofer heat received in Grenoble (heat D83334) and the resulting powder was called Eurofer powder 1999, whereas the heat used mainly in this study is the heat E83699 and the atomized powder was called Eurofer 2000.

Figure 1 : Large particles with satellites Powder Eurofer 2000. (Magnification x100)

Both powders were used and tested in order to control the reproducibility of atomisation process. The powder 99 and Eurofer powder 2000 present same shape and size aspects and chemical composition. Compared with F82 H powder previously prepared [2] , the shape of Eurofer powder looks less spherical and large particles present more inclusions or carbides. POWDER CONSOLIDATION Different compaction High Isostaticl Pressure cycle parameters in the temperature range [950-1100°C] have been tested on the powder99 or powder2000 and are detailed on the table 1.


All the HIP tests were performed with a gas pressure of 140MPa. The cooling rate was limited by the canister size and HIP facility to 15°C/minute.

Temperature

Time 950 °C 1000°C 1040°C 1100°C

2h P99 P99/P2000

4h P2000

5h P2000

10h P99

Table 1 : HIP powder compaction temperature and time parameters investigated on Powder99 or Powder2000

The microstructure has been observed in the as HIPed + tempered 2h at 750° condition by optical microscopy and SEM. Further investigations and especially TEM observations are in progress. MICROSTRUCTURE The microstructure of tempered material (2 h at 750°C) after HIP at 1040°C show martensite laths with carbide precipitation (figure 2). The prior austenite grain size seems to be between 10 and 20µm but would be precisely determined next year and the prior particle boundaries are sometimes. The microstructure The material HIPed at 1100°C revealed the same microstructure.

Figure 2 : Microstructure of Eurofer Powder 2000 HIPed at 1040°C /tempered for 2h at 750°C

The mechanical properties at room temperature of the solid HIPed materials were tested in the following metallurgical conditions : as HIPed with a cooling rate of 15°C/mn; after HIPing and tempering 2h at 750°C ; some tests were performed after 1hour in order to follow the hardness and impact properties evolution; as HIPed + normalisation 1h at 950°C/water quenched (called Post Normalisation Treatment : PNT) and tempered at 750°C -2h. The hardness results indicate that even if the cooling rate is low, martensitic structure is obtained in all cases. Nevertheless , the higher the HIP temperature is , the lower the martensitic hardness is.

MECHANICAL PROPERTIES Results of impact tests, performed at room temperature most on charpy V notch specimens , are presented on table 1. The Eurofer powder consolidated at 1040 or 1100°C present acceptable impact results. The Impact properties for 1100°C HIP consolidation temperature followed by a Post Normalisation Treatment reached high value 320J/ cm2.

Table 1 : Impact properties at room temperature for Eurofer Powder P99 and P2000

Powder HIP

conditions

Energy J/cm2

P2000

Tempering time (h) at 750°C

Energy J/cm2 P99

HIP conditions

206 1

155 2 163-U 1000°C- 5 h

248 *PNT +

2h 156-U

950°C- 10 h

228 1

237 2 226 1040°C- 4 h

276 *PNT +

2h 288

1040°C- 2 h

218 1

267 2 192

Powder

2000

1100°C- 2 h

321 *PNT +

2h 324

1100°C- 2 h

Powder

99

* PNT : Post Normalisation Treatment : + 950°C-1h + water quenched / U =charpy U

The post heat treatment improve the impact properties. The impact properties of the powder 99 and Eurofer powder2000, consolidated under same conditions present the same impact properties after Post Normalisation Treatment : 324 / 320J/cm2. That means that the process : atomisation, filling, HIPing is relatively reproducible. TENSILE RESULTS The HIP temperature effect on tensile properties is lower than for impact properties. The typical tensile properties are given on table 2 and the best HIP conditions seems to be 1040°C -4h or 1100°C - 2h.

Table 2 : Typical tensile results at room temperature on HIPed Eurofer P2000

HIP Post Heat

treatment

Yield Stress 0.2% MPa

U.T.S MPa

Uniform Elongation

%

Total Elongation

%

Reduction of area %

1040°C- 4 h

*PNT + 2 h at 750°C

514 632 7,0 16,1 71

1100°C-2 h

*PNT + 2 h at 750°C

506 623 6.5 15.9 72


CONCLUSION Two Eurofer cast alloys were atomised and different HIP consolidation conditions were tested. The chemical analysis does not reveal a large deviation in chemical composition between cast alloys and atomised powder. The consolidation process parameters examined were : the powder size, the filling process and the different HIP cycles temperature varying in the range [950°C-1100°C]. The microstructure and mechanical results reveal that the best compromise is obtained with the 1040°C HIP consolidation temperature and the 1100°C cycle reveal the highest impact properties. Further microstructure observations must be realised to confirm these first results. REPORTS AND PUBLICATIONS Note technique DEM N° 120/00 Eurofer Powder HIP development and Characterization J.M. GENTZBITTEL, I. CHU

TASK LEADER Jean Marie GENTZBITTEL DTEN/SMP/LS2M CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 95 01 Fax : 33 4 76 88 58 91 E-mail : [email protected]


TTMS-4.2 Task Title : RAFM STEELS - QUALIFICATION FABRICATION PROCESSES Mechanical alloying including ODS INTRODUCTION Raising the operating temperature of fusion reactors is one of the keys, which enable to improve their efficiency. Present use of Reduced Activation Ferritic Martensitic (RAFM) steels allows reactors’ operating temperatures around 550°C. The use of Oxide Dispersion Strengthened (ODS) RAFM steels offers the potential of temperatures up to 650°C. Until now, no ODS steels with reduced activation have been commercially available. Our main objective in this task is to demonstrate the benefit of ODS RAFM steels for blanket design solutions. As powder Hot Isostatic Pressing (HIP) is the principal and a very interesting process to manufacture large net-shape or near net-shape parts for structural applications, our work is focussed on this consolidation technique. The study mainly consists in the optimisation of the complete process of elaboration of ODS RAFM steels. The work has started last year with the production of different batches of ODS EUROFER powder and their consolidation [1]. The work programme of this year was to characterize the microstructure and the mechanical properties of these samples. In addition, consolidation by extrusion has been investigated and new batches of ODS EUROFER powder have been produced with different reinforcement powders. 2000 ACTIVITIES INITIAL POWDERS Two batches of EUROFER powders have been used as initial powders. Both have been atomised under argon by Osprey (one in 1999, one in 2000) from the same forged bar of one heat of “EUROFER 97” supplied by Böhler (E83699). In order to get better homogeneity after mechanical alloying, only the powder grains having a size less than 45 µm have been used. The chemical compositions of the heat from Böhler and of the two batches of atomised powder are reported in Table 1 with the specifications for the EUROFER 1997 label [2]. Two kinds of fine oxide powders have been used as reinforcement: Y2O3 and MgAl2O4. Both oxides have a good stability under irradiation. Most of the samples have been reinforced with yttrium oxide, which is the most commonly used for ODS alloys.

Table 1 : Chemical composition in weight % of EUROFER alloys before and after atomisation

Elements Specifications

EUROFER 1997

Böhler Forged

bar

Osprey Powder

1999

Osprey Powder

2000

Fe Bal. Bal. Bal. Bal.

Cr 8.50 – 9.50 8.87 8.6 8.97

C 0.090 – 0.120 0.12 0.091 0.095

W 1.00 – 1.20 1.10 1.05 1.07

V 0.150 – 0.250 0.19 0.19 0.18

Ta 0.050 – 0.090 0.14 0.086 0.1 ± 0.02

Mo < 0.0050 < 0.0010 0.0027 < 0.0010

Nb < 0.0010 < 0.0010 0.0020

Mn 0.20 – 0.60 0.42 0.38 0.40

Si < 0.050 0.06 0.091 0.050

O < 0.010 0.0013 0.055 0.043

N 0.015 – 0.045 0.018 0.024 0.013

Ni < 0.0050 0.0075 0.14 0.078

Co < 0.0050 0.0095

Cu < 0.0050 0.021 0.030

P < 0.0050 0.004 0.0056 0.0026

S < 0.0050 0.003 0.005 0.0040

Ti < 0.0100 0.008 0.0060

MECHANICAL ALLOYING ODS EUROFER powders have been produced by the mechanical alloying technique, which consists in milling together the EUROFER powder with the reinforcement powder. Two kinds of mills have been used: ball mills with a capacity of 5 kg per batch, and attritor mills with a capacity of 400 g. Both these kinds of mills are commonly used in industry. Reinforcement with Y2O3 Batches of ODS EUROFER powders have been produced in ball mill: - from powder 1999 with 0,5 and 1 wt % of Y2O3; - from powder 2000 with 1 wt % of Y2O3. A complete control of the milled powders has been done for each batch [3]: - Morphological characterization by SEM (figure 1). - Optical microstructural characterization (figure 2). - Particle size distribution analysis. - Chemical analysis of the potential contaminations

(Oxygen and Carbon).


For powders reinforced with 1% of Y2O3, the mean particle size is 50 µm and 80% of the particles are smaller than 100 µm. The particle size of the powders reinforced with 0,5% of Y2O3 is 35% bigger than the particle size of the 1% reinforced powders. The milling induces no carbon contamination. The average oxygen content after milling of 1% reinforced powders is 3200 weight ppm. The level of oxygen content due to global contamination (atomisation and milling) is relatively low (0.1%) and can be considered as acceptable.

Figure 1 : Morphology of a 1% Y2O3 reinforced powder

Figure 2: Microstructure of a 1% Y2O3 reinforced powder

Reinforcement with MgAl2O4 Batches of ODS EUROFER powders have been elaborated in attritor mill from powder 2000 with 1 weight % of MgAl2O4 spinel. A sum of 2,4 kg of powder have been produced. A complete control of the milled powders has been done [3]. The powder particles reinforced with MgAl2O4 are more flat than the powder particles reinforced with Y2O3. The microstructure is also thinner. The mean particle size is 75 µm and 75% of the particles are smaller than 100 µm. Indeed, all these differences are due to the higher milling efficiency of the attritor mill compared to the ball mill, rather than to the nature of the reinforcement oxide. The milling induces no carbon contamination. The oxygen content after milling is 4500 ± 200 ppm, which corresponds to the theoretical oxygen content brought by 1 weight % of MgAl2O4. CONSOLIDATION Hot Isostatic Pressing Each kind of available powder has been consolidated by HIP with the following cycle : 1020°C/2h/1000bar under argon. In addition to the samples consolidated from powder produced this year, samples of the four ODS EUROFER grades consolidated last year (from powders milled in attritor or in ball mill, reinforced with 0,5 and 1% of Y2O3) have been characterized by hardness measurements and metallographical observations. The material obtained from Y2O3 reinforced powder is fully dense and has the same microstructure as the milled powder (figure 3). The hardness of 0,5% reinforced alloys is 475 ± 10 HV1, and 505 ± 10 HV1 for 1% reinforced alloys, whatever the powders were milled in attritor or ball mill. Alloys produced from powder 2000 (including the spinel-reinforced alloy) have not been observed yet.

Figure 3 : Microstructure of a 0,5% Y2O3 reinforced sample consolidated by HIP

40 µm

100 µm 20 µm


Hot extrusion Powders produced in ball mill with 0,5 and 1% of Y2O3 have been extruded into bars of 16 mm diameter, at 1100°C with an area reduction ratio of 14:1. The material seems fully dense, but the microstructure is much thinner than for HIP samples, so that grain size and possible porosity cannot be determined except by TEM observations. Contrary to the HIP samples, the microstructure is elongated parallel to the extrusion axis. Hardness measurements (both in the longitudinal and transversal sections) give the same values than for HIPped samples. HEAT TREATMENTS One HIPped sample of EUROFER reinforced with 1% of Y2O3 has been heated at 950°C for 1 hour and water quenched. No major differences have been noted compared to the as-HIPped state regarding the microstructure and the hardness (502 HV1). Dilatometric analyses have been performed on a as-HIPped sample reinforced with 1% of Y2O3, in order to identify the occurrence or not of martensitic phase transformation since the difference between ferritic and martensitic microstructures is not resoluble by metallographical observations, due to the particularly fine microstructure of ODS steels. Different cycles have been defined from the conditions of HIPping (1020°C, cooling rate –15°c/min) and quenching (950°C, dilatometer maximum cooling rate: -100°C/min). The results show that both as-HIPped and quenched states are fully ferritic (Table 2), while a fully martensitic state is obtained after furnace cooling on non-ODS EUROFER steels produced by powder metallurgy [4].

Table 2 : Results of dilatometric analyses Temperature (°C) 1020 1020 950

Cooling rate (°C/min) -15 -100 -100

Ferrite start Fs (°C) 730 705 705

Ferrite finish Ff (°C) 665 645 640

Martensite start Ms (°C)

Martensite finish Mf (°C)

Final microstructure 100% ferritic

100% ferritic

100% ferritic

The martensitic transformation occurs when the transformation of austenite into ferrite and cementite, cannot take place due to too low kinetics of diffusion. As the ferritic transformation has occurred with ODS EUROFER alloys, even at high cooling rate, it seems that the diffusion has been strongly enhanced. This may result from the high density of grainboundaries since the grain size is very fine (< 5 µm), as revealed by TEM observations in an other study [5]. Indeed, the high density of grainboundaries could also explain why the hardness of ODS EUROFER alloys, which microstructure is ferritic, is so high when the hardness of tempered ferritic steels is generally around 220-240 HV1.

If we want martensitic transformation to occur, one solution can lie in the grain coarsening by an appropriate heat treatment. However, E. Brion [6] has shown that a fully martensitic structure can be obtained on a 9Cr-1Mo-0,1C ferrito-martensitic steel reinforced with 0,25% of Y2O3 after a heat treatment of 30 min at 1000°C and a minimum cooling rate of –15°C/min. For the following of the study, lower reinforcement contents should be produced and characterized in term of heat treatment conditions to obtain a fully martensitic structure. MECHANICAL TESTING As-HIPped and as-extruded EUROFER alloys reinforced with 0,5 and 1% of Y2O3 have been tensile tested at room temperature at a strain rate of 8.10-4s-1. All grades show very high yield stress and ultimate tensile stress (Table 3), far beyond the specifications required for non-ODS EUROFER steels, but HIPped samples have insufficient elongation. As expected, extruded samples proved to be more resistant and more ductile than HIPped samples. With appropriate forming processes and heat treatments one can expect the elongation of HIPped ODS EUROFER steels to be improved beyond 15% while keeping high level of tensile resistance.

Table 3 : Room temperature tensile properties

0.2% yield

stress (MPa)

Ultimate tensile stress (MPa)

Elongation to rupture

(%)

Specifications for EUROFER steels [2]

500 600 15

EUROFER + 0,5% Y2O3 HIPped

870 1175 1.8

EUROFER + 1% Y2O3 HIPped 910 1205 5.0

EUROFER + 0,5% Y2O3 Extruded 1100 1565 12.5

EUROFER + 1% Y2O3 Extruded 1240 1665 8.7

CONCLUSION The four ODS EUROFER grades consolidated last year by HIP (from powders milled in attritor or in ball mill, reinforced with 0,5 and 1% of Y2O3) have been characterized by hardness measurements and metallographical observations. New batches of ODS powders have been produced from a new batch of atomised EUROFER powder with either 1% of Y2O3 or MgAl2O4 powder as reinforcement. Powders mechanically alloyed in ball mill with 0,5 and 1% of Y2O3 have been also consolidated by extrusion. An attempt of quenching has been performed on a 1% Y2O3 reinforced sample consolidated by HIP. Dilatometric analyses have revealed that the microstructure was fully ferritic after both HIP- and quenching-like cycles.


It was suggested that the ferritic transformation was promoted by enhanced diffusion kinetics due to the high density of grainboundaries. Room temperature tensile tests have been performed on as-HIPped and as-extruded materials reinforced with 0,5 and 1% of Y2O3. The next steps will be to produce new grades of ODS EUROFER with lower reinforcement contents and to investigate, for each composition, the possible heat treatment conditions to obtain a martensitic transformation. High temperature tensile tests and room temperature impact tests will be performed. Finally, microstructural and mechanical characterization of the material reinforced with MgAl2O4 will be realized. REPORTS AND PUBLICATIONS [1] S. Revol, C. Labonne, “Task SM4.2 : Mechanical

alloying including ODS”, CEA report NT DEM N°71/99, nov. 1999.

[2] W. Dietz, A. Alamo, R. Lindau, “Technical

specification : supply of reduced activation ferritic martensitic steel type 9CrWTaV”, NET 97/917, European Blanket project – Structural Material Task 5.3, 23/07/97.

[3] S. Launois, S. Revol, C. Labonne, “Task TTMS-4.2 :

Mechanical alloying including ODS”, CEA report NT DEM N°106/2000, dec. 2000.

[4] L. Federzoni, “Powder HIP development and

characterization – Task SM4.1”, CEA report NT DEM N°80/99, dec. 1999.

[5] S. Launois, S. Revol, C. Labonne, “Task UT-SM&C-

ODS : Development fo forming and joining technologies for ODS steels”, CEA report NT DEM N°117/2000, dec. 2000.

[6] E. Brion, “Etudes métallurgiques des alliages base fer

renforcés par dispersion d’oxydes”, CEA report, NT SRMA 00-2362, april 2000.

TASK LEADER Stéphane REVOL DRT/DTEN/SMP CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 98 42 Fax : 33 4 76 88 54 79 E-mail : [email protected]


TTMS-4.3 Task Title : RAFM STEELS - QUALIFICATION FABRICATION PROCESSES RAFM weldability INTRODUCTION The development of fast induced radioactivity decay alloys has been incorporated into the goals of the IEA collaborative program in order to reduce the quantity of highly radioactive materials from system components. Concerning the innermost parts of a Fusion reactor, the low activation martensitic (LAM) steels and reduced-activation ferritic/martensitic (RAFM) steels are considered as primary candidate alloys as first wall and breeder blanket structural materials because of high irradiation resistance, especially the resistance to the phase instability and the resistance to the shift of the ductile-brittle transition temperature (DBTT) during or after irradiation. Several RAFM steels have been developed specially for this application in the frame of the Long Term Fusion international collaborative test program. Among these materials, the F82H steel developed by JAERI in the mid nineties and the more recent Eurofer 97 proposed by FZK and Böhler are considered to be specially viable candidates. Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding and laser welding processes are envisaged for the fabrication of structural parts of relevant breeding blanket modules. A preliminary study of the weldability of F82H has been carried out in the frame of the task SM-4.3. The present study is the following of SM-4.3. It aims to evaluate both TIG (with and without filler material) and EB weldabilities of Eurofer 97 and to compare it to the F82H weldability. During the year 2000, the TIG weldability of Eurofer 97 plates of 14 and 24 mm of thicknesses has been investigated by using Eurofer 97 filler metal.

2000 ACTIVITIES OPERATING CONDITIONS The Eurofer 97 cuts and filler wire have been delivered in January 2000. Plates of 200 * 80 mm2 have been sampled in the cuts of 14 and 24 mm of thicknesses in order to weld assemblies of 160 mm of width for both thicknesses : 1 14-mm assembly and 7 24-mm assemblies (Table 1). The general operating conditions were taken from those applied on F82H in the frame of the task SM-4.3. In particular, the same deep chamfer design, the same clamping conditions (by welding the plates to a rigid support) and the same welding speed have been (4 cm/min) used for comparison. The general welding parameters are summarised in Table 2. As previously for the F82H, the filling of the chamfer at each pass has been measured in order to settle a predictive law of filling, knowing that at the end, the total thickness of the seam may not be more than 1.1 times the plate thickness according to the RCCM rules. Each weldments has been X-rayed. Further, between the first passes, the welding has been interrupted in order to carry out US control of the bottom weldments. Each weldment has followed metallographical examinations in cross section. A post welding heat treatment of 1 hour at 710°C has been applied on a selected weldments. Microhardness measurements have been carried out in cross section before and after the post welding heat treatment.

Table 1 : Series of weldability tests

material tests series thickness (mm) clamping pre-heating thermal instr. strains evaluation

F0B0101 14 weld No No No

F0B0201

F0B0301

F0B0302

F0B0303

F0B0304

F0B0305

F0B0306

Eurofer 9724 NoNo Noweld


Table 2 : General welding conditions of Eurofer 97 plates of 14 and 24 mm of thicknesses

mat

eria

l

nb p

asse

s

thic

k (m

m)

clam

ping

U (

V)

I (A

)

wel

ding

spe

ed (

cm/m

in)

wir

e sp

eed

(m/m

in)

wir

e di

amet

er (

mm

)

scan

ning

wid

th (

mm

)

uppe

r sh

ield

ing

(l/m

in)

back

shi

eldi

ng (

l/min

)

back

side

shi

eldi

ng (

l/min

)

5 14 welded 10 120-240 4 0,7-1,6 3,5-8,5 14 0 7,5

8 24 welded 10 150-230 4 0,7-2 1 3,5-9,5 15 20 7,5Eurofer

RESULTS The operating conditions settled for the TIG welding of the F82H were shown to be rather suitable when applied to the Eurofer 97 in both cases (14-mm and 24-mm thick plates). Sound weldments have been obtained in such conditions. However, it has been observed during welding, that the molten pool tend to stick to colder zones, leading to sticking flaws. This particular behaviour imposed to carry out fusion lines on such passes in order to re-melt the flaw. This method led to sound weldments of both thicknesses (figure 1). Using the same cumulative filler deposition, the filling rate is lower in the case of the Eurofer 97 (figure 2).

FOB0201

Eurofer97 14 mm Eurofer97 24 mm

Figure 1 : Macroscopic aspect of Eurofer 97 weldments of 14 and 24 mm of width in cross section

Figure 2 : Chamfer filling laws

2 mm

y = 5E-05x2 - 0,0891x + 20,461

y = -2E-06x2 - 0,0877x + 10,735

y = 4E-05x2 - 0.0708x + 20.471

-5

0

5

10

15

20

25

0 50 100 150 200 250 300 350 400 450

cumulative filler deposition in section (mm2)

dep

th le

ft (

mm

)

F82H 24 mm

Eurofer 14 mm

Eurofer 24 mm

2 mm

2 mm


In the as-welded state, the melted zone displayed a high microhardness (420 Hv1). The post welding heat treatment that has been applied did not low down the hardness sufficiently. The treated melted zone displayed 320 Hv1, that is about 100 Hv1 more than the base material. This result indicated that the duration or the nominal temperature of the treatment should be increased (figure 3).

FOB0201

Eurofer97 24 mm as-welded

FOB0201 TT (710°C : 1 h)

Eurofer97 24 mm PWHT

Figure 1 : Macroscopic aspect of Eurofer 97 weldments

of 24 mm of width in cross section, before and after post welding heat treatment of 1 h at 710°C

CONCLUSIONS Matching filler material has been used to study the TIG weldability of thick plates of Eurofer 97 (thicknesses of 14 and 24 mm). This study has been completed by the second half of 2000. It has shown that accurate operating conditions lead to sound welds despite the need of intermediate fusion lines between the deposit passes. The conditions (time, temperature) of post welding heat treatment (PWHT) that have been selected have led to insufficient softening. REPORTS AND PUBLICATIONS A Fontes, "TIG weldability of Eurofer 97 with matching filler material", CEA report DECS/STA/LMS/01-RT4007 TASK LEADER André FONTES DRT/DECS/STA/LMS CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 37 21 Fax : 33 1 69 08 75 97 E-mail : [email protected]

2 mm

2 mm


TTMS-4.4 Task Title : RAFM STEELS - QUALIFICATION FABRICATION PROCESSES Dissimilar welding RAFM/316LN INTRODUCTION The development of fast induced radioactivity decay alloys has been incorporated into the goals of the IEA collaborative program in order to reduce the quantity of highly radioactive materials from system components. Concerning the innermost parts of a Fusion reactor, the low activation martensitic (LAM) steels and reduced-activation ferritic/martensitic (RAFM) steels are considered as primary candidate alloys as first wall and breeder blanket structural materials because of high irradiation resistance, especially the resistance to the phase instability and the resistance to the shift of the ductile-brittle transition temperature (DBTT) during or after irradiation. Several RAFM steels have been developed specially for this application in the frame of the Long Term Fusion international collaborative test program. Among these materials, the F82H steel developed by JAERI in the mid nineties and the more recent Eurofer 97 proposed by FZK and Böhler are considered to be specially viable candidates. Tungsten Inert Gas (TIG) welding, Electron Beam (EB) welding and laser welding processes are envisaged for the fabrication of structural parts of relevant breeding blanket modules. External parts of RAFM cooling tubes will be joined with 316LN tubes. The present study aims to evaluate the weldability of transition joint between RAFM steel and stainless steel 316LN. The 2000 work was dedicated to a literature review on the TIG feasibility of transition weldments. Further, some quantities of suitable industrial filler material has been supplied. 2000 ACTIVITIES LITERATURE RESULTS Dissimilar joining between austenitic and ferritic-martensitic steels are commonly employed in high-temperature applications in energy conversion systems. Differences in coefficients of thermal expansion between the two types of steels induce thermal stresses at the welds and local metallurgical changes near the low alloy steel/weld metal interface due to prolonged service at an elevated temperature [1],[2].

Concerning specifically the welding conditions, it is well known that both the ferritic and the martensitic steels can easily be welded by using austenitic stainless steel or nickel-based solid consumables that can avoid the occurrence of cold cracking in the weldments.

However, such fillers may display two major disadvantages as regards the application : they low down the mechanical characteristics of the welded material and they bring to the weldment undesirable elements of high activation, such as nickel [3], [4], [5]. The first criteria to be used to select a suitable filler material would be to avoid as possible both risks of hot cracking and cold cracking [2]. SUPPLYING OF FILLER MATERIAL Several filler products have been investigated with a view of supplying some quantities for a preliminary weldability study. The 1675 product from ESAB has been selected because of its high chromium content. If used at 20% of dilution in the weld, it should provide austenitic-ferritic-martensitic microstructure (Figure 1) avoiding a cracking risk.

0,00

5,00

10,00

15,00

20,00

0,00 5,00 10,00 15,00 20,00 25,00 30,00

Eurofer

316LN

F82H

ESAB1675 20%

ESAB1675 10%

ESAB1675

Figure 1 : Schaeffler diagram : prediction

of the microstructure of the melted zone according to various dilution conditions of the base materials

and of the filler product CONCLUSIONS The composition range for filler material dedicated to the TIG welding of RAFM steel with SS 316LN has been defined according to the literature review carried out in the frame of the 94-98 task SM4-5 and according to the Schaeffer diagram. This study allowed to select an commercial filler product. This material has been delivered by December 2000. A technical report is now in progress.


REFERENCES [1] Roberts, DI, Performance of dissimilar welds in

service, J. of Pressure Vessel Techno., Vol. 107, n°3, pp. 247-254, Aug. 1985.

[2] Feldstein, JG, Dissimilar welds with stainless steels,

ASM International, ASM Handbook, Vol. 6, Welding, Brazing and Soldering, 1993, pp. 500-504.

[3] Tamura M., Hayakawa H., Tanimura M., Hishinuma

A., Kondo T., Development of potential Low Activation Ferritic and Austenitic Steels, J. of Nucl. Mater., 141-143 (1986) 1067-1073.

[4] Davis J.R., Selection of wrought martensitic stainless

steels, ASM Handbook, Vol. 6, Welding, Brazing and Soldering, Ed. J.R. Davis, 1993, pp. 432-442.

[5] Cousement C., Backer T. de, Dewitte M., Verelst L.,

Internal pressure creep testing on welded components of modified 9%Cr and 12%Cr steels, Proc. of Eurojoin 2, 2nd European Conference on Joining Technology, Florence, Italy, 16-18 May 1994, Publ. 15-16141 Genova, Italy, 1994, pp. 461-479.

REPORTS AND PUBLICATIONS A Fontes, "Filler material for the TIG welding of RAFM/316LN transition joints", CEA report DECS/STA/LMS/01-RT4008. TASK LEADER André FONTES DRT/DECS/STA/LMS CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 37 21 Fax : 33 1 69 08 75 97 E-mail : [email protected]


TTMS-4.5 Task Title : RAFM STEELS - QUALIFICATION FABRICATION PROCESSES Solid HIP process and qualification INTRODUCTION Diffusion welding is considered as a manufacturing technique for the fabrication of DEMO European blankets. The structural material is a 9%Cr - 1%W martensitic steel. Previous work on HIP diffusion welding of Eurofer has shown that good results could be obtained in the range 1050-1100°C. The full mechanical characterization of joints HIPed at 1100°C is currently made in the frame of task TTMS 2. In this task, the work aimed at the assessment of the possibility to decrease the HIP temperature to 1050°C in view of the WCLL blanket design that includes copper compliant layers in the first wall. Encouraging results were obtained in 1999 [1]. 2000 ACTIVITIES The joining process involves the following steps : surface preparation, HIP and cooling, normalisation at 950°C and quenching, tempering at 750°C. The HIP & post - treatment parameters are shown on figure 1.

0

300

600

900

1200

1500

0 2 4 6 8 10 12 14 16 18 20time (h)

T(°

C)

HIP 100MPaNormalisation

950°C, 2htempering750°C, 2h

quench

Figure 1 : HIP parameters and post treatment QUENCHING RATE Experiments were made to assess the influence of the quenching rate on the HIP joints mechanical properties. Samples were as shown on figure 2. The surface preparation was the same than in [1], i.e. polishing the surfaces to be joined with a SiC paper grade 1200, degreasing the samples and the canister with a mixture of alcohol, acetone and ether, drying with a neutral gas and outgassing the parts in a furnace under vacuum at 950°C for 2h.

The parameters and the results are given in table 1. The results are rather dispersed. All tensile specimens broke away from the joints except one specimen from sample 8. The impact toughness is very low, except for sample 4. All KCU specimens broke at the joint, however the rupture was partly away from it for specimens from sample 4. Pores were found at the joints but only few inclusions were observed. Sample 4 was not HIPed in the same batch than the following ones, and some temperature difference might be suspected, as can be seen from the difference of base metal properties. A slight strength decrease can be noticed for the sample quenched at the lowest rate.

Figure 2 : Scheme of the diffusion welding experiments EFFECT OF HIP TEMPERATURE The objectives were to assess the sensitivity of the joining process to the HIP temperature. Experiments were made at 1050°C, 1065°C, 1080°C and 1100°C for 4h, followed by post heat treatments defined in figure 1 (water quench was used). Again, all tensile specimens broke away from the joints. The impact toughness increases with temperature. The microstructure of the joints is shown on figure 3. No porosity was noticed contrary to sample 6. The joints are marked by alignments of particles, most probably carbides. When the HIP temperature increases, the joints become less marked, which indicates better carbide dissolution.

Vacuum pump

canister

Eurofer

samples

KCU for joint testing

KCU for base material testing


Table 1 : Room temperature mechanical properties of joints HIPed at 1050°C and quenched at various rates

Tensile (2 specimens) 4 spec. 3 spec.

Sample Quench (°C/h)

Yield stress 0.2% (MPa)

Max Strength (MPa)

Elongation (%)

Reduction of area

(%)

KCU joint (J/cm²)

KCU base metal

(J/cm²) HV5

4 Water 550 667 17.5 68 143 187 224

5 Air 532 648 20 80 43 218 216

6 300 524 645 19.5 78 15 231 216

7 200 521 645 20.5 77 52 224 220

8 100 511 638 20 78/53 15 226 211

Table 2 : Room temperature mechanical properties of joints HIPed at various temperatures

Tensile (2 specimens) 4 spec. 3 spec.

Sample HIP (°C)

Yield stress 0.2% (MPa)

Max Strength (MPa)

Elongation (%)

Reduction of area

(%)

KCU joint (J/cm²)

KCU base metal

(J/cm²) HV5

9 1050 481 629 21.5 n/a 13 n/a 210

10 1065 518 614 24 n/a 6 n/a 206

11 1080 549 629 27 n/a 60 207 191

12 1100 534 623 25 n/a 147 220 204

Sample 9, HIP 1050°C/4h

1cm = 8µm


1cm = 8µm


1cm = 8µm


1cm = 8µm

Figure 3 : Microstructure of the joints HIPed at various temperatures


CONCLUSIONS The good results obtained during 1999 with a HIP temperature of 1050°C were not confirmed this year. This is probably due to the fact that during 1999 the samples were HIPed twice, which seems beneficial. If a single HIP cycle is used, residual porosity and particle alignments (most probably carbides) are visible at the interface. Tensile properties of the joints are excellent whatever the HIP temperature in the range studied (1050-1100°C). Acceptable joint impact toughness cannot be obtained in a reliable manner with a HIP temperature less than 1100°C. This is coherent with the fact that dissolution of refractory carbides (that tend to prevent grain boundary mobility) begins at a temperature higher than 1050°C. Basic surface preparation consisting in polishing and solvent degreasing is efficient. A post-HIP heat treatment consisting in a low temperature austenitisation at (950°C, 2h), quench and tempering at (750°C, 2h) allows to restore a fine grain size and to achieve acceptable hardness level in the range 200-220HV. The quenching rate must not be less than 200°C/h. REFERENCES [1] E. Rigal, Task SM45, NT DEM 91/99 REPORTS AND PUBLICATION [1] E. Rigal, task TTMS4.5 final report, NT DEM

n°111/2000

TASK LEADER Emmanuel RIGAL DRT/DTEN/SMP/LS2M CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 97 22 Fax : 33 4 76 88 54 79 E-mail : [email protected]


TTMS-5.1 Task Title : RAFM STEELS - RULES FOR DESIGN AND INSPECTION

Design code assessment and development INTRODUCTION The main objective of this sub-task is to write the Appendix A, " Materials Data and Analyses" for interim design rules of ITER Test Blanket Modules and DEMO components (DISDC). The work to be done is similar to the one carried out for ITER. In both cases designers use these rules as guidelines for design of reactor components and systems. Currently, ITER documents do not include rules and data for RAFM (Reduced Activation Ferritic/Martensitic) steels, i.e. the structural materials selected for the ITER Test Modules (ITM) and DEMO. The behaviour of these steels is, in some aspects, very different from those of the austenitic steels used in ITER and hence ITER rules cannot be reliably extended to them. In addition, the operating conditions of ITM and DEMO components differ from those of ITER (higher service temperatures and higher radiation doses) provoking damage mechanisms not considered in IISDC. This sub-task is closely related to sub-task TTMS 5.4, where, materials properties data are collected. 2000 ACTIVITIES PART 1: DISDC/A. GEN: GENERAL Appendix A begins with the section A.GEN, which contains the definitions of the physical and mechanical properties used in the DISDC together with the formulae used for calculating various design limits used in materials properties sections of Appendix A. The first interim version of this report was issued in 1999. It was updated in 2000. PART 2: APPENDIX A: A3S18: MODIFIED 9Cr-1Mo STEEL This report covers rules for the conventional 9Cr-1Mo steels. The behaviour of this material is considered to be relatively close to that of the RAFM steels. As a results, while awaiting results of tests on RAFM steels, Appendix A3S18 recommendations can be used as preliminary guidelines. The first interim version of this report was issued in 1999. It was updated in 2000.

PART 3: APPENDIX A: A3S18F: F82H STEEL The main part of the 2000 activities was dedicated to writing an interim Appendix A for F82H steel. Table A3.S18F.1.2.1 shows the IEA chemical compositions specified for this steel as well as the compositions obtained with the two large IEA heats. Compositions of TIG wire and welds produced in Japan are also shown.

Table A3.S18F.1.2.1. Chemical compositions of F82H steel and its TIG welds. Numbers are in wt % unless

otherwise indicated

IEA 1993

Specification

F82H Ladle Analysis

Heats

TIG Welds (Typical)

Element min max 9741 9753 Wire Weld

C 0.10 0.12 0.09 0.09 0.07 0.07

Si 0.05 0.20 0.11 0.08 0.09 0.10

Mn 0.05 0.20 0.16 0.1 0.46 0.46

P 0.01 0.002 0.003 0.003 0.004

S 0.01 0.002 0.001 0.003 0.003

Ni 0.1 0.02 0.02 0.005 0.008

Cr 7.5 8.5 7.66 7.89 7.41 7.57

Mo 0.05 <0.01 0.003 <0.005 <0.005

Cu 0.05 <0.01 0.01 0.005 0.005

Al 0.1 0.001 0.001 <0.002 0.002

V 0.15 0.25 0.16 0.19 0.18 0.18

W 1.8 2.2 2.00 1.99 2.02 2.02

Ti 0.004 0.012 0.01 0.004

Nb 0.0002 <0.01 0.0002 <0.005 <0.005

B 0.001 0.0002 0.0002 <0.0002 0.0004

N2 0.02 0.005 0.006 0.0051 0.0045

Ta 0.01 0.06 0.02 0.02 0.023 0.026

Sn 0.004

As 0.005

Sb 0.004

Co 0.01 <0.01 0.003

Ag 0.002

O 0.005

Tb

Ho

Zr

Bi

Eu

Physical properties sections of the Appendix A were mostly derived from the JAERI work.

- 236 - EFDA technology / Tritium Breeding and Materials / Structural materials development / RAFM steels Mechanical properties data were extracted from the databases established in TTMS 5.4, with data mainly supplied by EU, Japan and US. However, the results of tests performed on F82H steel are still coming in and a major revision of this Appendix will be issued in 2001. When writing the current version of Appendix A, the available data had to be filtered to retain only those satisfying the code requirements (e.g. RCC-MR). In this process, results from many small heats tested during alloy development stage had to be discarded and only those from the large commercial heats (5000 kg) to be retained. Even then, test results deviating from the reference heat treatments or testing conditions had to be omitted. Despite the above restrictions, in most cases the databases were sufficient in content to allow preliminary analyses. An example of such analysis with its associated recommendations is calculation of Sm values from tension test results. In figure A3.S18F 5.1.1, it is shown that Sm values of these alloys are governed by 1/3 of Su (ultimate tensile strength). This is due to the fact that the difference between the yield and tensile strengths of these alloys is small and hence 1/3Su is significantly lower than 2/3Sy. The unirradiated Sm values also apply to irradiated state, since Sy and Su are increased after irradiation and, in general, the difference between Sy and Su is further reduced.

100

150

200

250

300

350

0 100 200 300 400 500 600 700

F82H (IEA Heat)

Base Metal

Sm

, M

Pa

Temperature, C

2/3 Sy(min, T)

1/3 Su(min, T)

2/3 Sy(min, 20C)

Sm

Figure A3.S18F.5.1.1. : Variation of Sm versus temperature for F82H steel

Another example of analysis made is the creep rupture master curve calculation. Figure A3.S18F 3.5.3 shows that a good correlation is obtained between experimental results and trends derived from the master creep rupture equation.

Another example, but very preliminary, is establishment of a fatigue design curve. Here, not only the data are few, but they come from three different sources and show significant scatter.

10

100

1000

100

101

102

103

104

105

F82H (IEA Heat)

Comparison of Exp. & Fitted Values

Str

ess,

MP

aTime to Rupture, h

450 C

500 C

550 C

600 C

650 C

700 C

Figure A3.S18F.3.5.3. : Comparison of experimental creep

rupture results with trends derived from master curve

0.1

1

10

102

103

104

105

106

F82H (IEA Heat)

20 - 550 °C

Tentative Fatigue Design

Tot

al S

trai

n R

ange

, %

Cycles to failure, Nr

Average (20-550°C)

Nr

/20

∆εt

/2

Design

Fig. A3.S18F.6.4.2. : Tentative fatigue design curve to be updated in future editions

CONCLUSIONS The two interim reports issued in 1999 are updated in 2000: - Appendix A. General Rules - Appendix A. Rules for a conventional 9Cr steel An interim Appendix A report for FH2H steel is issued.


REPORTS AND PUBLICATIONS [1] Appendix A. Material Design Limit Data, AGEN:

General, Interim SM 5.1 & 5.4 Reports, NT DECM/Dir 2000-11, Jan. 2000

[2] Appendix A. Material Design Limit Data, A3.S18,

Modified 9Cr-1Mo Steel, Interim SM 5.1 & 5.4 Report, NT DECM/Dir 2000-10, Nov. 2000

[3] Appendix A. Material Design Limit Data, A3.S18F,

F82H Steel, Draft, Interim SM 5.1 & 5.4 Report, NT DECM/Dir 2000-42, Dec. 2000

[4] “EU Contribution to RAFM steel database”, IEA

Workshop on Reduced Activation Ferritic/Martensitic Steels, 2-3 Nov. 2000, Tokyo.

TASK LEADER Farhad TAVASSOLI DEN/DMN CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 60 21 Fax : 33 1 69 08 8070 E-mail : [email protected]


TTMS-5.4 Task Title : RAFM STEELS - RULES FOR DESIGN AND INSPECTION

Data collection and data base maintenance INTRODUCTION Reduced activation ferritic/martensitic steels (RAFM steels) are reference structural materials for ITER Test Blanket Modules and DEMO reactor. After initial screening tests, using mainly small heats, composition of these steels were narrowed down to 7-10% chromium steels. JAERI in Japan ordered a 5000 kg heat to NKK and designated it as F82H steel. Small amounts of this heat was sent to Europe and US. Under the auspices of the Annex II IEA Fusion Implementing agreement, the composition of F82H was later re-specified (1993), particularly to reduce activation levels after irradiation. Two new 5000 kg heats (9741, 9753) were subsequently produced by NKK according to these specifications. These new heats are known as IEA heats and the one before as pre-IEA heat. The first IEA heat was rolled into plates of 7.5 and 15 mm and the second into plates of 15 and 25 mm thickness (products). These plates were then cut into smaller coupons (sub-products) and dispatched to IEA partners. Several TIG and EB welded joints were also produced from these plates in Japan and shared with the partners. Europe, in particular CEA and NRG, also fabricated F82H welded joints, to which one should add powder HIP and diffusion bonded joints. Characterization of the sub-products has been underway since and results obtained by various laboratories have been regularly discussed at the IEA sponsored workshops. The European Union for her part has defined a new RAFM composition (Eurofer) and has procured from Boehler several 2000 kg heats of this alloy (Eurofer 97). These heats have also been shared with the IEA partners. The main difference between Eurofer and F82H steel resides in lower tungsten (1% instead of 2%) and higher chromium contents (9% instead of 8%) of the former as compared with the latter. EU is planning to fine-tune the composition of Eurofer and order new heats in 2001. The primary objective of this sub-task is to collect, validate, harmonize the results generated under the above collaborative programs. Initially the work will focus on F82H steel and later it will extend to Eurofer.

2000 ACTIVITIES Work continued in 2000 with the establishment of a product database (605 records). Information on RAFM steels, starting from melting procedures, thermo-mechanical treatments, fabrication of large plates (products) down to cutting of small plates (sub-products) or fabrication of various joints are collected, structured and entered in the database. As a result, each sub-product received in a laboratory can be fully traced back to its origin and manufacturing procedures. Its position in the parent plate vis-à -vis other sub-products can also be situated. The second database expanded in 2000 is the chemical compositions database (476 records). Here, the available chemical compositions for heats, products, sub-products or even given zones of a sub-product are collected. Not only the conventional ladle or heat analyses are entered, but also, when available, those performed at the laboratories on sub-products, including e.g. on the weld deposits of TIG joints. The mechanical properties databases are also expanded and now include tensile (784 records), impact (1112 records), impact plots (120 records), fracture (26 records), creep (127 records) and fatigue (26 records) databases. Each mechanical properties database is structured at the specimen level. Information about specimens is documented with respect to their geometry, position and orientation in the sub-product, as well as their thermo-mechanical, aging and irradiation history. Testing parameters, standards and analysis procedures are also recorded. The mechanical properties databases are organised in a way that allow sorting and finding of conventional results as well as those needed for advanced analysis. For instance, in the case of tension tests, in addition to yield stress and ultimate tensile strength, provisions are made for values of stresses at strains 0 to 10%, thus allowing (when these values are reported) full description of tension curves. Likewise, in addition to uniform elongation, total elongation and reduction of area at rupture, values of strain to necking for irradiated materials are recorded. This information is particularly useful for materials that show a pseudo peak in tension curves just after yielding. The impact properties database has provisions for conventional and instrumented impact testing results. Impact tests performed on a given serie of specimens (with the same thermo-mechanical and irradiation history) are grouped together and linked to an Impact plot database.


There, the results are plotted and fitted with a tangent hyperbolic sign curves. From these curves values of upper shelf energy, DBTT, etc. are calculated. The fracture database is detailed for full description of results from such tests. However, the number of tests reported so far is small. This database will be expanded in 2001. The creep database is relatively large and covers both creep rupture and creep deformation. Good correlations are obtained between predicted and experimental results. The fatigue database includes information on continuous fatigue cycling, fatigue with hold times in tension and/or in compression, fatigue relaxation and fatigue hardening. However, the number of tests reported so far is small or are for different types of testing, e.g. thermal fatigue. This database will be expanded in 2001. A link is established between all mechanical properties databases and the product and the chemical compositions databases, allowing full traceablity. The package as it stands, allows searching and sorting of results according to different criteria. This is very useful for sorting and finding code qualified data. For instance, for Appendix A, all results were initially sorted out for commercial heats of F82H.

Then the results were further filtered to find those with reference thermo-mechanical treatments. Finally, each mechanical property database was filtered according to the required testing procedures. The outcome was then used in TTMS 5.1. Figure 1 shows an example of search results found for creep tests on F82H at 550C. Another possibility offered by the package is full correlation between various databases. A summary database is established, where for a given heat, a given product, or a given sub-product, all relevant information from the databases are extracted and summarized on single page. Figure 2 presents an example of a summary data sheet produced for heat 9741. CONCLUSIONS Significant progress has been made in establishing a comprehensive materials properties database for RAFM steels. Work will continue in 2001 with databases of F82H steel and Eurofer.

Figure 1 : An example of search performed in the creep database to find out tests performed on F82H steel at 550 C


Figure 2 : Summary sheet showing available results in the databases on heat 9741 of F82H steel REPORTS AND PUBLICATIONS [1] Appendix A. Material Design Limit Data, AGEN:

General, Interim SM 5.1 & 5.4 Reports, NT DECM/Dir 2000-11, Jan. 2000.

[2] Appendix A. Material Design Limit Data, A3.S18,

Modified 9Cr-1Mo Steel, Interim SM 5.1 & 5.4 Report, NT DECM/Dir 2000-10, Nov. 2000.

[3] Appendix A. Material Design Limit Data, A3.S18F,

F82H Steel, Draft, Interim SM 5.1 & 5.4 Report, NT DECM/Dir 2000-42, Dec. 2000.

[4] “EU Contribution to RAFM steel database”, IEA

Workshop on Reduced Activation Ferritic/Martensitic Steels, 2-3 Nov. 2000, Tokyo.


TASK LEADER Farhad TAVASSOLI DEN/DMN CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 60 21 Fax : 33 1 69 08 8070 E-mail : [email protected]


TTMS-6.3.1 Task Title : RAFM STEELS - QUALIFICATION OF HIGH PERFORMANCE

STEELS Development of RAFM-ODS steels INTRODUCTION High efficiency of power reactors could be reached with higher operating temperatures. Conventional RAFM steels allow operation somewhat below a temperature of 550°C. Oxide Dispersion Strengthened (ODS) materials offer the potential of an operating temperature of about 600-800°C. The main objective of this task is to assess ODS-RAFM alloys, based in a Fe-Cr ferritic or martensitic matrix strengthened by oxide dispersion particles, for applications as structural materials in the high temperature range (T > 550°C). ODS are obtained by particular manufacturing methods, i.e., mechanical alloying of metallic and oxide powders, which are subsequently consolidated by hot-extrusion or HIP processes. The metallurgical and mechanical behaviour of ODS is strongly dependent on the fabrication processes used. 2000 ACTIVITIES Activities performed during 2000 were essentially oriented to define the most interesting chemical compositions of ODS and to prospect the feasibility of manufacturing such as type of materials. It is worthwhile to point out that only few possibilities exist in Europe to perform the mechanical alloying at industrial as well as laboratory scale. As discussed in different technical meetings, FZK ordered the manufacturing of two ODS-materials at Plansee (Austria). These materials have been delivered to CEA-Saclay on December 2000. The programme of the metallurgical and mechanical characterization has been defined and is actually in progress. ODS IRON-BASED ALLOYS AND CHARACTERI-ZATION PROGRAM The typical microstructure of ODS Fe-base materials consists on a metallic matrix reinforced by a distribution of nanometric particles of oxides, which determines the strength level. Thus, the high temperature strength of ODS is strongly dependent on the microstructure, that is, the size distribution of oxide particles, the homogeneity of oxide particle dispersion, the morphology and grain size of the matrix, the texture.

Depending on the in-service temperature range envisaged, two types of Fe-Cr ODS alloys need to be considered for the assessment and qualification of ODS materials: a) Steels with 7-11%Cr that could present a martensitic

transformation. In this case, a martensitic matrix strengthened by oxide dispersion and having isotropic properties could be obtained for applications as structural materials in the range 600-700°C. In this case, in-service temperatures are limited by the occurrence of the martensite to austenite phase transformation at about 800°C.

b) Steels with 12-14%Cr with a ferritic matrix structure,

which does not present any phase transformation. They could be envisaged for higher in-service temperatures (700-800°C). Nevertheless, this type of ODS alloys need an optimisation of their microstructure by controlling the fabrication route to minimise the eventual anisotropy of metallurgical and mechanical properties induced by the particular methods of manufacturing used.

In the first time, the objective of this task is to characterize the mechanical (tensile) properties in relation to their microstructure of two ODS-martensitic alloys based on the EUROFER97 (Fe-9Cr-1W-0.1C) steel (type a). These steels have been manufactured from atomised powders of EUROFER97, mechanically alloyed with two different fractions of Yttrium oxide and consolidated by HIP. The nominal compositions (in wt. %) are: 1) Fe-9Cr-1W-0.1C + 0.3 Y2O3 (ref. HXN 954/4) 2) Fe-9Cr-1W-0.1C + 0.5 Y2O3 (ref. HXN 958/2) Materials have been supplied as consolidated round bars of 60mm in diameter. On the other hand, mechanically alloyed powders with both Yttrium oxide concentrations have been also delivered to be consolidated in Saclay by hot-extrusion. The programme of characterization will include the following steps: - Control of the as received bars (as-consolidated

condition) by optical and electron scanning microscopy, hardness measurements to determine grain size, inclusion content, eventual fabrication defects (porosity).

- Control of the chemical composition homogeneity of

the as received bars, in particular for the Yttrium oxide distribution. For this purpose, composition profiles will be determined with the electron microprobe to detect the eventual local variation of different elements.


- Effects of austenitisation and tempering treatments on the microstructure and hardness values.

- Definition of the most suitable heat-treatment conditions

to be applied. - Determination of the tensile properties. This program recently started on the as-received bars. Similar tests will be also performed on materials that will be compacted by hot-extrusion at Saclay. REPORT - Progress Report TTMS 6.3.1, December 2000, A.

Alamo, CEA/DMN/SRMA/01.07.

TASK LEADER A. ALAMO DMN / SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 67 26 Fax : 33 1 69 08 61 75 E-mail : [email protected]

- 245 - EFDA technology / Tritium Breeding and Materials / Structural materials development / IFMIF facilities

TTMI-001 Task Title : IFMIF - ACCELERATOR FACILITY INTRODUCTION The IFMIF CDE has been carried out at the CEA in the framework of a considerably larger activity presently undergoing in the field of high-intensity linear accelerators. The activity specific to IFMIF has been concentrated on: - Task 1: ECR source long run test, beam transport and

analyses. (RAMI, beam quality, perfor-mances).

- Task 4: 4-vane RFQ design, optimization of RFQ

electric field on cold model, RFQ-DTL transition.

- Task 5: Collection of RF tube progress – Monitoring

of long test tube. - Task 6: Detailed DTL design, beam dynamics, drift

tubes, multiparticle simulations, error studies, halo studies.

- Task 9: ECR source to DTL high power diagnostics. The accelerator must demonstrate reliability commensurate the maintenance schedule of the facility. Low to very low spark rates have to be shown, as well as reliability in term of linac performances. Explicit contact with every task specific to IFMIF and the results that have been obtained are summarized in the next sections. 2000 ACTIVITIES TASK 1 : ECR SOURCE AND LEBT SILHI is in operation at CEA-Saclay. It delivered its first proton beam in July 1996. This ECR (Electron Cyclotron Resonance) source is designed to produce continuous high-intensity proton and deuteron beams. Calculations for deuteron operation have been specifically done to reach the IFMIF requirements and a new extraction systems has been calculated and installed. SILHI is operated with proton beams because it is presently installed in an unshielded area that precludes deuteron acceleration. Nevertheless, all the results obtained using protons can be directly applied to the deuteron beam required by the IFMIF accelerator, for example: - control of high currents, - stability of the operation in continuous mode,

- microwave-plasma coupling, - design/technologies used for the 100 kV extraction

system, - materials used for the plasma chamber. The LEBT is also constructed and a full set of diagnostics is now in operation to allow a complete analysis of the beam characteristics. Again, all these diagnostics have been designed to reach the IFMIF goals. The measurements done using high-intensity continuous proton beams show that they are perfectly suitable for IFMIF. This equipment designed and tested at CEA-Saclay on the SILHI source for IFMIF includes several systems: - Focusing system and mechanical set up. - Pumping system. - CCD camera and associated software for beam profile

measurements. - Emittance measurement units, classical and using the

CCD camera technique. - Beam intensity monitors. After several studies of the microwave-plasma coupling and several tests of material (plasma chamber and extraction system), the following results have been obtained: - Intensity = 157 mA in the continuous mode. - Plasma density = 247 mA/cm2 (φ = 9 mm). - Proton fraction = 80%. - Intensity stability = ± 1%. - Normalized rms emittance = 0.11-0.33 π.mm.mrad. Since the new extraction system installation, leading to lower beam losses and higher LEBT transparency, SILHI currently produces 95 keV - 120 mA total beam with a proton fraction higher than 80 %. Up to 157 mA (247 mA/cm2) total CW beam has been extracted with a single ECR zone inside the plasma chamber. The new remote control based on EPICS system using EMC devices and automatic procedures has been installed to perform new long runs (several weeks of continuous operation). To analyze these upgrades a 4-week long run test is presently under progress.


The authorization to extract a pulsed deuterium beam has been obtained. In May, in the framework of the IFMIF project CEA participation, a 3 GHz klystron will allow deuteron pulsed beam production and analysis. TASK 4 : FOUR VANES RFQ The IFMIF accelerators begin with a deuteron source and a low-energy beam transport to a radiofrequency quadrupole (RFQ) up to ~5 MeV, then a DTL. Very low beam losses along the accelerator are required in order to allow maintenance without requiring remote manipulators. Sparks lead to RFQ damage and beam shut down. The task concentrates in improvement of the ratio RFQ length/maximum field in the RFQ, while the transmission is kept as high as possible. According to other high power RFQs, the big effect observed at the RFQ front end (not fully understood) has to be mastered. The new RFQ tool named TOUTATIS was developed in CEA-Saclay to reach higher accuracy in the beam dynamics simulations. The RFQ beam dynamics optimization is almost finished. The transmission through the RFQ was raised up to 96.3% and the transmission of accelerated particles up to 95.9%. More important the output beam emittance in both the longitudinal and transverse plane are optimized for injection into the DTL. RFQs are powerful accelerators to accelerate and bunch high intensity beams. Nevertheless, the IFMIF beam power and IFMIF RFQ design complicates the RF tuning of this part of the accelerator. The task also consists in using the CEA-Saclay IPHI RFQ cold model to develop the tuning method and associated software and hardware for the IFMIF RFQ RF tuning. Great advance was made on the tuning procedure of the 4-vanes RFQ. A 6-meters long cold model has been installed and the tuning process formalism is investigated. First results show the ability to detect a 30 µm vane misplacement before a brazing step. The ability to tune the field within 0.1% of the theoretical value in only three steps on the 32 slugs 1 m long element was demonstrated. TASK 5 : RF SYSTEM It could be recalled that the RF system is one of the biggest challenges for the IFMIF accelerator. The IFMIF RF system frequency is 175 MHz, tetrode or diacrode tubes must be used instead of the classical klystrons used for the other projects requiring higher frequencies. The CEA-Saclay team activity concerning the 1.0 MW - 175 MHz RF systems is then only devoted to the IFMIF project. CEA-Saclay activity has been mainly to interact with Thomson Tubes Electronics (TTE) on the IFMIF aims and to collect, evaluate, and transmit all the available information to the IFMIF team. TTE has produced a very substantial effort to demonstrate the feasibility of such RF systems.

Following the IFMIF accelerator team request, a 1000 h long test run is under discussion with TTE. Their test stand needs to be upgraded to proceed to the test at 1 MW cw. These preliminary results obtained by TTE are very promising, and suggest convincingly that a RF system, up to full IFMIF performance, could be built and operated satisfactorily. TASK 6 : DTL STUDIES IFMIF accelerators must deliver 250 mA continuous current of deuterons at a nominal energy of 40 MeV, with provision for operation at 30 MeV and 35 MeV. The basic approach is to provide two linacs modules, each delivering 125 mA to a common target. This approach has been taken to reach the requested availability and to give operational flexibility advantages. Beam loss limitation is a critical issue for every high-power accelerator (structure activation), especially for IFMIF due to the acceleration of deuterons. An important point concerning the beam loss problem is the ability designers have to predict them up to extremely low levels. For IFMIF, relative beam loss must be below 10-10 of the total beam. As development of the RFQ and RFQ-DTL arises, the DTL beam dynamics needs to be adjusted keeping the losses under the limits. Using the new tools developed at CEA-Saclay, a DTL beam dynamic design has been proposed in August 2000. Improvements of this design are in progress to take into account the post couplers limitations and the new RFQ design (beam matching to the DTL). The drift tube engineering is the delicate parts of DTL accelerators. The drift tube and quadrupoles optimization involves many calculations of the cavity with different shapes to minimize the RF power lost in copper, the IFMIF design will be consolidated by experiments on the IPHI DTL hot model (measured power deposition, cooling, manufacturing tolerances…). The 2 IPHI drift tubes hot models have been successfully electron beam welded and show good thermal performances. The DTL tank is under construction but shows vacuum leaks. Troubles for the repairs and delays are unfortunately expected. The hot tests will take place in 2001 in CERN. Some delays occurs also last years with the postpone LEP shutdown to November 2000. TASK 9 : HIGH POWER DIAGNOSTICS IFMIF beam power is about 15 kW between the source and the RFQ, 625 kW before the DTL and 5 MW in the HEBT. The relatively low energy (40 MeV) means that this power is deposited in any interceptive diagnostics. New kind of diagnostics has to be developed including intensified CCD analyses, spectroscopy analyses, fast wire scanners, laser analyses… A vigorous effort is undertaken for the diagnostics of CW high power beams. The emittance measurement unit has been upgraded (smaller size makes easier displacements and faster analyses).


A large number of experiences are made in fluorescence and absorption techniques to analyze the beam position, size and profile. The absorption techniques lead to some difficulties up to now but the other techniques may still be used. CCD camera images and profiles are crosschecked with classical electric measurements in order to control their accuracy. CONCLUSIONS Very significant progresses have been obtained recently, especially on the source and RFQ. The soon expected long runs on both ECR SILHI source and TTE Diacrode are the next important milestones for the accelerator part of the IFMIF studies. PUBLICATIONS - TOUTATIS, the CEA-Saclay RFQ Code, R. Duperrier,

ICAP 2000 (Darmstadt). - Status Report on the Saclay High-Intensity Proton

Injector Project (IPHI), P-Y. Beauvais et al., EPAC 2000 (Vienne).

- HIGH-POWER PROTON LINAC FOR A MULTI-

USER FACILITY, J-M Lagniel, EPAC 2000 (Vienne). - Linac architecture for high power proton sources, J-M

Lagniel, LINAC 2000, Monterey. - Status report on the 5 MeV IPHI RFQ, R. Ferdinand et

al, LINAC 2000, Monterey.

TASK LEADER Jean-Michel LAGNIEL DSM/DAPNIA/SEA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 53 65 Fax : 33 1 69 08 85 62 E-mail : [email protected]

- 357 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials

UT-SM&C-LAM/Mic Task Title : INFLUENCE OF THE MARTENSITE MORPHOLOGY ON THE

PLASTICITY BEHAVIOUR OF THE EUROFER STEEL INTRODUCTION During processes used during manufacturing (heat treatments, HIP, joining) and under service conditions, the microstructure is modified and induces changes in the mechanical properties of materials. A better understanding of the relationships between the microstructural parameters and the plasticity is necessary to improve the mechanical behaviour. So, the main objective of this task is to perform a generic study on the influence of the morphology and the crystallography of the laths and lath packets of martensite on the deformation mechanisms of the alloy. 2000 ACTIVITIES Activities developed on 2000 consisted in: - a bibliographic study on the martensitic transformations

in steels [1], - the determination of heat treatments to control the prior

austenite grain size, - characterization by Transmission Electron Microscopy

of laths and lath packets orientation in the Eurofer steel in the as received (AR) condition [2],

- determination of the dislocation type present in the

Eurofer steel (AR) [2]. MATERIAL The material studied in this task is the reduced activation martensitic steel EUROFER 97. The chemical composition of this steel developed by Europeans involved in fusion programs is the following :

Fe C N Cr W Mn V Ta Si

bal. 0.117 0.0219 8.96 1.04 0.48 0.18 0.15 0.03

The as-received (AR) conditions are : normalisation at 980°C during 30 min and tempering at 760°C during 90 min. CONTROL OF THE SIZE OF THE PRIOR AUSTENITIC GRAIN One point of this study consist in obtaining different microstructures with a Prior Austenite Grain (PAG) size ranging from few microns to more than one hundred microns.

Indeed, this size govern in part the morphology and the crystallography of martensite laths. So, the first part of this study was to determine the PAG size as a function of the normalisation temperature (see figure1). This was done by performing a special heat treatment that allows to induce a precipitation of carbides on PAG boundaries after the austenitisation treatment and to obtain the PAG size.

1000 1050 1100 1150 12000

20

40

60

80

100

120

140

150Alliage EUROFER

Dia

mèt

re M

oyen

Pro

jeté

Temperature (°C)

Figure 1 : Evolution of the size of the prior austenite grain

as a function of the austenitisation temperature (30 minutes at the austenitisation temperature)

MICROSTRUCTURAL OBSERVATIONS Orientation relation ships between laths and lath packets In the Fe-9Cr steels with a low carbon content, as Eurofer steel, the orientation relation ships between the prior austenite and the martensite are the Kurdjumov-Sachs relations :

(011)α // (111)γ et [1 1 1]α // [0 1 1]γ The analysis of this relations indicates that four family of martensite laths have a habit plane corresponding to the four prior (111)γ planes. Each family is constituted by laths with six possible orientations. So, in theory it is possible to obtain 24 orientations for the laths within a PAG. Different observations were conduct to determine experimentally the orientation relation ships between packets of laths in PAG. One of them consists on the determination of the orientation of each lath by indexing, with a software developed by the University of Metz (Laboratoire d’Etudes des Textures et Application aux Matériaux) the Kickuchi pattern obtained on the lath. The orientation cartography (see figure 2) and the data obtained allow to measure the disorientation between lath packets and to obtain a morphologic and cristallographic description of the laths and laths packets.


reference triangle

Figure 2 : Orientation map obtain by TEM on the Eurofer steel

Two laths parallel with the same habit plane can be separated by a dislocation wall (see figure 3).

Figure 3 : Laths with a low angle boundary. The boundary is formed by a dislocation wall.

Laths belong to the same variant.

In that case, laths have a low angle boundaries (< 3-4°) and have the same variant. They can also have an higher angle boundaries (10°) and correspond to two different variants (see figure 4).

Figure 4 : Laths (1 and 5) with an angle boundary close to 10°. This boundary can not be analysed as a

dislocation wall. Laths belong to two different variants.


Dislocations network The density of dislocations in the Eurofer in the AR conditions depends of the area considered. Areas with a high-entangled dislocation density are detected but also areas with a low dislocation density are also observed (figure 5). In that case the identification of the dislocation type was possible. They are screw dislocations with a Bürgers vector ½ a 0<111>.

Figure 5 : Screw dislocations (arrowed in black) in the as-received Eurofer steels, g = 101 .

CONCLUSIONS The determination of heat treatments to control the PAG size was performed. They can now be applied to obtain relevant microstructures for this study. The morphology and the crystallography of the martensite laths in the as-received Eurofer 97 steel was studied and the orientation maps of laths were obtained by TEM. Two parallel laths with the same habit plane could have a low angle boundary. In that case, the boundary is constituted by a dislocation wall and laths belong to the same variant. They can also have a higher angle boundary (>10°) and correspond to different variants. The analysis of dislocations present in the as-received Eurofer 97 shows that they are screw dislocations with a Bürgers vector ½ a 0<111>. Investigations on deformed materials and materials with a larger PAG size are in progress.

REFERENCES [1] J. Laurencin, “ Microstructure et dureté de la

martensite dans les aciers ”, Note Technique D.E.M. N°2000/71, septembre 2000.

[2] Y. de Carlan, L. Guetaz, “Analyse morphologique et

cristallographique de la structure martensitique dans l’acier Eurofer“ Progress Report UT-SM&C-LAM/Mic, CEA report, NT SRMA 00-2400, décembre 2000.

TASK LEADER Y. de CARLAN DEN/DMN/SRMA CEA Saclay 91191 Gif-sur-Yvette Cedex Tél. : 33 1 69 08 61 75 Fax : 33 1 69 08 71 30 E-mail : [email protected]


UT-SM&C-LAM/Weld Task Title : LASER WELDABILITY OF LAM STEELS (Eurofer97) INTRODUCTION The aim of the task consists in evaluating the laser-weldability of low activation martensitic steels and characterizing the mechanical properties of welded joints, with specific attention given, in 2000-2001, to Eurofer97 (9% Cr 1% W). More, results have to be compared with previous work on F82H (8% Cr - 2 % W) steels. Different laser sources have to be used, depending on the thicknesses of the assemblies (6 to 9 mm). They range from the pulsed regime (1.2 kW YAG), to the continuous wave configuration (4-7 kW YAG and 18 kW CO2). 2000 ACTIVITIES To begin with, the last report on the laser weldability of F82H ( “laser welding of F82H with dual cw+pulsed mode) has been delivered in 07/2000. Between 05/2000 and 03/2001, three weeks of laser experiments have been performed : 1 week with 1.2 kW pulsed YAG, and 2 weeks with 4 kW cw YAG, with objective to optimize the laser conditions driving to good quality butt-welded joints. Due to to the small amount of metal available for the experiments, 75 mm x 50 mm x 6 mm specimens have been used. All experiments have been performed in flat position. PULSED 1.1 kW YAG WELDING First welding trials on Eurofer were carried out with nearly the same conditions than those previously used on F82H in flat position : 30 Hz – 4.5 ms – 1 kW and a 200 mm focal lens resulting in a 1 mm spot.

With such conditions, full penetration was achieved on 6 mm tools below V = 1 mm/s. Most of times, it occurred together with matter ejection at the rear and front surfaces of the assembly and collapse of the joint (Fig 2.a). This was attributed to the excessive Pp / V (with Pp= peak power density ≈ 106 W/cm²) input, driving to a drilling-like effect. Therefore, it seems that the molten pool stability during pulsed laser welding is more critical on Eurofer than on F82H, even when defocusing to reduce the peak power densities. This may be attributed to different viscous behaviours. Despite surface aspects, x-ray inspections indicate a very low porosity content in pulsed regime, even for partially penetrated butt-welds. Cw 4 kW YAG WELDING Cw welding experiments have been performed on 6 mm and 9 mm thick tools, with a 150 mm focal lens (0.45 mm spot), a 4 kW mean power, and welding speeds between 3 and 14 mm/sec. The defocus distance was taken at –1 mm (1 mm below the surface) to favour the penetration and increase the welding speed. Full penetration was achieved below 3 mm/sec on 9 mm tools, and below 12 mm/s on 6 mm tools. However, even if surface aspects are much better than after pulsed welding (Fig 1a and 1b), the porosity content is increased (nearly 10 pores of 50-300 µm diameter / cm). These pores are mainly localized in the bottom part of the bead and attributed to instabilities of the bottom part of the key-hole. Furthermore, the porosity ratio tends to increase at low welding speeds (below 9 mm/sec), indicating a detrimental influence of convective effects. Both cw and pulsed regime drive to the same hardenings in the weld and heat affected zone (WZ and HAZ) when compared to the base material (BM) : from 230 HV in the BM to 450 HV in the WZ and HAZ (Fig 3). This factor 2 increase is close from those achieved on F82H.

Figure 1 : Surface aspects after 1.1 kW pulsed (a) and 4 kW cw (b) YAG butt-welding on 6 mm tools


Figure 2 : Cross-sections (x 12) of butt-assemblies after 1.1 kW pulsed welding (0.9 mm/s), and 4 kW cw welding (10 mm/s)

200

250

300

350

400

450

500

-5 0 5

pulsed 1.1 kW - 1 mm/scw 4 kW - 11 mm/s

HV0.5

position (mm)

WZ + HAZ

Figure 3 : Cross sectional Hardness measurements after pulsed and cw YAG laser welding

CONCLUSIONS AND ONCOMING WORK The first YAG laser trials on 6 mm and 9 mm Eurofer97 without filler wire have indicated a good weldability in cw regime, but difficulties to find accurate welding conditions (≈ penetration without collapse) in pulsed mode, maybe due to the insufficient mean power (1.05 kW max versus 1.2 kW for the F82H experiments), and to different viscous behaviors in the molten pool. In cw mode, good quality beads were obtained, despite the occurrence of porosities, mainly at low welding speed. Metallurgical structures were found to be nearly the same for both welding conditions. In the following months, mechanical testings (Charpy and tensile tests) will be carried out on laser-welded joints (pulsed and cw conditions), before and after tempering treatments of the coarse martensite structure. Results on Eurofer will be compared with those from F82H welding. Last, 18 kW CO2 laser and 4+3 kW YAG (twin optical fiber) will be carried out on 9 mm tools, with and without filler wire.

PUBLICATIONS « Pulsed and CW YAG Laser Welding of F82H Desactivated Steel for Nuclear Applications ». P. Peyre, L. Sabatier, P. Aubert and R. Fabbro, in ICALEO ‘2000 (Int.Conf.Appl.Lasers and Electro Optics), Detroit (USA), Oct 2000. TASK LEADER P. AUBERT P. PEYRE (CLFA) LIST/DECS/CLFA CEA Fontenay aux Roses Boîte Postale 6 92265 Fontenay aux Roses Cedex Tél. : 33 1 42 31 87 09 Fax : 33 1 42 53 97 47 E-mail : [email protected]


UT-SM&C-LAM2 Task Title : IRRADIATED BEHAVIOUR OF REDUCED ACTIVATION (RA)

MARTENSITIC STEELS AFTER NEUTRON IRRADIATION AT 325°C

INTRODUCTION The objective of this task is to study the metallurgical and mechanical behaviour of FeCrW Reduced Activation (RA) compared to the conventional FeCrMo martensitic steels during neutron irradiation at 325°C, relevant temperature for fusion reactor applications. Irradiated specimens of each material are available for 5 levels of radiation, damage ranging from 1 to 9 dpa. These doses appear as enough to investigate the first step of irradiation-induced embrittlement and hardening of such materials. Post-irradiation examinations (PIE) involve tensile tests, measurements of reduction in area to rupture, fractographic examinations and microstructural studies by transmission electron microscopy (TEM) of irradiated specimens. PIE also include dose decay rate and spectrometric measurements of activated elements after irradiation. 2000 ACTIVITIES The irradiation experiment performed in Osiris reactor finished on November 1999 and last specimens irradiated at the highest dose (about 9 dpa) were unloaded on November 99. All the samples were transferred to the hot-cells and mechanical tests on specimens irradiated at 9 dpa began at the end of 2000. Due to the planning of works in hot cells, all the tensile tests initially scheduled for 2000 were not carry out but the gamma spectrometry facility was put in order and measurements on selected steels were performed. On the other hand, measurements of dose decay rate as a function of the cooling time have been continued for specimens irradiated with a dose of 0.9 dpa. Thus, data are available now up to 53 months after unloading. All materials with an activity sufficiently low to be analysed by Small Angle Neutron Scattering (SANS) were prepared and transferred to the Laboratoire Léon Brillouin (see UT-SM&C-LAM3). Thin foils on F82H and LA12LC steels irradiated at the lowest dose (0.9 dpa) were prepared during 2000 and first examinations on these steels were performed by Transmission Electron Microscopy (TEM).

IRRADIATION CONDITIONS Irradiation in Osiris reactor started on January 96 and finished on the second semester 1999. Because of partial unloadings, irradiated samples are available with 5 different dose levels, that is, 1, 2 , 3.4, 5 and 9 dpa. Materials were irradiated at 325°C (+5°C, -10°C) in pressurised water at 155 bars as tensile specimens and plate samples intended for different post-irradiation examinations. Irradiation was carried out in a mixed neutron spectrum. The maximal neutron flux was about 2.1014 n/cm2s (E > 1 MeV) and the fast to thermal flux ratio is about 1.1. MATERIALS Two types of relevant martensitic materials for fusion applications were irradiated in the present experiment, that is, 7/11CrW RA-steels and 9/12CrMo conventional martensitic steels for comparison (see UT-SM&C-LAM3 for the chemical composition of the materials) The following nuances of RA martensitic steels were irradiated : - F82H is a 7.5CrWTaV RA-steel developed by JAERI

(Japan). - LA12LC, LA4Ta and LA13Ta are RA experimental

alloys of 9/11CrWTaV type with different contents of Cr, W, Ta.

Conventional martensitic steels are commercial alloys which present different contents of Cr (9 to 12%), Mo and stabilising elements (V, Nb). Dose rate decay measurements and gamma spectro-metry To check the concept of " reduced activation ", the dose rate measurements are performed on several types of steels, including stainless steels, conventional and reduced activation martensitic alloys, all of them irradiated in the same conditions. Measurements have been conducted on specimens of the first unloading (0.9 dpa), which was obtained on March 1996. The dose rate is measured with a Geiger-Müller counter for different distances between the samples and the counter, that is, 50, 100 and 185 mm.

- 364 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials The first measurements have been performed 9 months after unloading. Six months later these measurements were performed again and after that every 12 months. They are now available after 9, 15, 27, 39 and 53 months. The results obtained during 2000 [1] confirm the trends already observed. After 9 months of cooling (see figure 1), the group of stainless steels (type 304 and 316) present the highest residual activity (25-45mGy/h) and both types of martensitic steels are characterised by dose rates in the range 8-18mGy/h. Measurements performed after different cooling-times show a decrease of the residual activity for all materials.

Figure 1 compares the dose rate values measured after 9 months and after 53 months for a distance sample-counter of 100 mm. After this period of time, the activity of stainless steels is always the highest and ranges from 6 to 28 mGy/h, the conventional martensitic steels CrMo from 1 to 6 mGy/h and in the case of the reduced activation CrW mart ensitic steels the dose rate is close to 1.0 mGy/h. Experimental data are in quite good agreement with decreasing activity calculated with the Microshield code. As shown in figure 2, the relative rate of the decreasing activity is different for each class of steels.

0

5

10

15

20

25

30

35

40

45

50

304

H30

4 H

316

E

316

LN

316

Ti E

EM10

St

EM10

St

EM10

E10

%

T91 S

td

MANET 2

Std

M3

StdT5

E

2 1/

4 Cr

LA 1

2 LC

E

LA 1

2 LC

E

LA 1

3 Ta

E

LA 4

Ta

E

LA 4

Ta

E

F 82

H Std

F 82

H Std

Steels

Do

se r

ate

(mG

y/h

).

9 months after unloading

53 months after unloading

Figure 1 : Dose rate measured after 53 months of unloading for different classes of steels irradiated at 0.9 dpa in the Osiris reactor

0%

10%

20%

30%

40%

50%

60%

70%

80%

90%

100%

304

H30

4 H

316

E

316

LN

316

Ti E

EM10

St

EM10

St

EM10

E10

%

T91 S

td

MANET 2

Std

M3

StdT5

E

2 1/

4 Cr

LA 1

2 LC

E

LA 1

2 LC

E

LA 1

3 Ta

E

LA 4

Ta

E

LA 4

Ta

E

F 82

H Std

F 82

H Std

Steels

Var

iati

on

of

the

do

se r

ate

(%).

Variation between 9 and 15 months Variation between 9 and 27 monthsVariation between 9 and 39 months Variation between 9 and 53 months

Figure 2 : Relative decrease of the dose rate measured after 53 months of unloading for different classes of steels irradiated at 0.9 dpa in the Osiris reactor

- 365 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials Compared to the first measurements (9 months), 304 and 316 steels exhibit a reduction of the dose rate of 38-70%, 42-80% in the case of conventional martensitic alloys, and 89-95% for the reduced activation steels during a period of 4.5 years. During 2000 the gamma spectrometry facility was put in order and measurements on the two less active austenitic steels, three conventional martensitic steels and four low activation martensitic steels were performed. In this equipment, the spectrometry is performed on massif samples. It is worthwhile to note that in these steels, the majority of the dose rate comes from the gamma ray and that the technique allows to measure the concentration of the radio element in the steel. In all the analysed steels, only two radio elements were identified by gamma spectrometry, the 54Mn and the 60Co. The figure 3 shows the activity of 60Co (in Bq/g) in the steels, it follows qualitatively the dose rate of the considered material.

1.00E+06

1.00E+07

1.00E+08

1.00E+09

304H 316Ti EM10std T91 T5 LA13Ta LA4Ta F82H LA12LC

Steels

Act

ivit

y o

f th

e 60

Co

th

e 12

/07/

00 (

Bq

/g).

Figure 3 : Activity of 60Co (in Bq/g) in steels irradiated in

Osiris at 0.9 dpa – 53 months after unloading The lowest 54Mn activity are observed for austenitic steels and martensitic steels exhibit virtually the same 54Mn activity. Those results are in good agreement with the estimations performed by RAEPSAET [2], who identified the radio elements that should be present in the steels after irradiation. After a cooling time of 39 and 51 months, for all the considered steels, the majority of the activity comes from 54Fe, 60Co, 63Ni, and 54Mn and among these elements only 54Mn and 60Co induce gamma radiation. The precursor of 54Mn is 54Fe (54Fe(n,p)54Mn) and the one of 60Co is 59Co (59Co(n,γ)60Co). As 54Fe and 60Co are present in the material before irradiation and do not come from a decay of other isotopes, it appears that after irradiation in PWR conditions and after a cooling time close to four years, the dose rate is essentially governed by the content of iron and cobalt in the alloy before irradiation. For longer cooling time (from 10 years to 1000 years for example) the contents of Ni, W, Mo and Nb have also to be considered. These experimental data assess the good performance of " reduced activation steels " from the point of view of radiological effects and environmental considerations after irradiation in PWR conditions. Consequently, reduced activation materials seems very encouraging for applications as structural materials for fusion reactors and other projects where components are submitted to in-service high doses.

CHARACTERIZATION BY TRANSMISSION ELECTRON MICROSCOPY Thin foils of F82H and LA12LC steels irradiated 0.9 dpa at 325°C were prepared in hot cells for examination by TEM. The 80 µm thick TEM disks were prepared electrochemically and TEM observations [3] were conducted on a Philips EM430 operating at 300 kV. F82H steel was irradiated in the as received condition (normalised and tempered) whereas LA12LC was irradiated in the normalised, tempered and 10% cold work conditions but first examinations we performed by TEM did not reveal any significant difference between the two materials. The figure 4 show a general view of the microstructure of the F82H steel after irradiation. As for the unirradiated material, the martensite laths within prior austenite grains are clearly observable and carbides are present along grain/lath boundaries. The irradiation does not seem to modify the precipitation and no recovery of the microstructure is observed. As unirradiated materials, irradiated samples (see figure 4) present dislocation networks in the matrix. In the irradiated specimens we also observed “black dots” and dislocation loops characteristic of the irradiation damage.

Figure 4 : Microstructure of F82H steel irradiated 0.9 dpa at 325°C

The size of the dislocation loops vary from 15 to 30 nm. As shown in figure 5, LA12LC steel exhibit similar microstructural evolutions as F82H steel. In both steels, no He bubbles or cavities were observed.

1 µm

300 nm


Figure 5 : Black dots and dislocation loop in the LA12LC steel irradiated 0.9 dpa at 325°C

CONCLUSIONS The irradiation experiment performed in Osiris reactor at 325°C finished on November 1999. Several Cr-W low activation and Cr-Mo conventional martensitic steels have been irradiated for five dose levels : 1, 2, 3.4, 5 and 9 dpa. Most of the last PIE were performed during 2000. Measurements of the dose rate have been performed after several cooling-times up to 53 months to validate the concept of " reduced activation materials ". For this purpose, specimens of different types of steels have been examined : austenitic steels, conventional martensitic and reduced activation martensitic alloys, after irradiation for a dose of 0.9 dpa. Compared to the conventional martensitic and to the austenitic steels, experimental data show that the reduced activation steels exhibit : − the lowest residual activity. − the highest rate of decreasing activity.(90%). The dose rate of steels is mainly due to the gamma-ray emission. Gamma spectrometry performed on the different types of steels shows that the major contribution to gamma radiation comes from 54Mn and 60Co. This fact confirms that after irradiation in PWR conditions and after a cooling time close to 5 years the iron and cobalt contents in these steels before irradiation govern the dose rate of the material. Preliminary examinations by TEM on F82H and LA12LC steels irradiated 0.9 dpa at 325°C were performed. No recovery of the microstructure is observed. Black dots and dislocation loops were detected but no cavities or He bubbles.

REFERENCES [2] C. Raepsaet, “Alexandre : calculs d'activité et de

débits d'équivalent de dose”, CEA report SERMA/LEPP/RT/00-2767/A

REPORTS [1] X. Averty, P. Coffre, G.M., Decroix, P. Duigou

“Irradiation Alexandre – débit de dose pour le panier 1 à To+53 mo is“ CEA Report, SEMI/LCMI/RT/00/040/A, octobre 2000.

[3] Y. de Carlan "Progress report 2000 - UT-SM&C-

LAM2- Post-irradiation examinations performed on RA and conventional martensitic steels irradiated at 325°C in Osiris reactor", Progress Report UT-SM&C-LAM2, CEA report, CR SRMA 01-1657.


100 nm

Dislocation loops

Black dots

g=110


UT-SM&C-LAM3 Task Title : MICROSTRUCTURAL INVESTIGATION OF REDUCED

ACTIVATION FERRITIC-MARTENSITIC (RAFM) STEELS BY SMALL ANGLE NEUTRON SCATTERING (SANS)

INTRODUCTION In order to assess reduced activation materials and to understand the effect of the substitution of W to Mo and the role of Ta on the microstructural evolutions during long-term thermal ageing or neutron irradiation, several RA martensitic steels of different compositions, with a chromium content between 7 and 12 at.%, were studied by Small Angle Neutron Scattering (SANS). Indeed, SANS allows, in this kind of materials, to characterize the mean size, shape and number density of very small precipitates, even at early ageing time when they are not detected by TEM. 2000 ACTIVITIES The complete characterisation of RAFM steels thermal aged in the range 250°C to 550°C was performed in 1999. During 2000, SANS was used to continue the characterisation of neutron irradiated materials for a dose of 0.8 dpa and the synthesis of obtained results was presented during the 20th International Symposium on the effects of Radiation on Materials [1]. Preliminary SANS experiments (not presented here) were also performed during this year on ferritic/martensitic ODS materials. MATERIALS AND EXPERIMENTAL SET-UP The chemical compositions of the steels used in this study are listed in the following table. Heat treatment applied on each materials are descried in [1].

The neutron scattering experiments were performed at the Laboratoire Léon Brillouin (Laboratoire Commun CEA-CNRS), Saclay, on PAXY and PAXE small-angle instruments. The experimental set-up and references on the data treatment (correction, normalisation, and calibration) can be found in [1]. THERMAL AGEING AND IRRADIATION CONDI-TIONS Samples were thermally aged under vacuum for various temperatures ranging between 250 and 550°C and for durations up to 22000 hours. Four materials (F82H, LA12LC, LA4Ta and HT9) were irradiated in the OSIRIS reactor at Saclay in PWR conditions, in the frame of the “Alexandre” irradiation [2]. The irradiation conditions were : temperature 325 ± 10 °C; fast neutron flux 2 1014 n/cm2.s (E > 1 MeV); duration 43 days, corresponding to an average displacement dose of 0.9 d.p.a. (displacement per atom). Thermal ageing EM10 steel : The scattered intensities after treatment at 400 and 450°C are similar and markedly lower compared to those observed after ageing at 500°C. The data treatment shows that the SANS profiles can be interpreted with a distribution of particles with a mean radius close to 3 nm. The A ratios (the ratio between scattered intensities measured perpendicular and parallel to the sample magnetisation that gives information about the chemical composition) are constant on the whole measured q domain and equal to 3.5 ± 0.6. This value could correspond to several kinds of particles : M2C carbides with the following chemical content M(at%)= 83Cr, 12Fe, 5Mo, or VC carbides.

Table 1 : Chemical composition of martensitic steels (wt. %)

Steel Cr W Mo Ni Mn Ta V Si C N

F82H 7.47 1.96 0.02 0.21 0.023 0.15 0.10 0.087 0.006

JLF-1 8.70 1.91 0.52 0.08 0.18 0.05 0.106 0.028

LA12TaLC 8.80 0.73 1.13 0.10 0.30 0.03 0.090 0.019

LA12LC 8.92 0.73 1.13 0.01 0.30 0.03 0.089 0.035

LA12Ta 9.86 0.84 0.88 0.10 0.28 0.03 0.155 0.043

LA13Ta 8.39 2.79 0.79 0.09 0.24 0.04 0.179 0.048

LA4Ta 11.08 0.72 0.78 0.07 0.23 0.03 0.142 0.041

EM10 8.76 <0.05 1.05 0.18 0.48 0.03 0.37 0.11 0.024

HT9 11.8 0.51 0.99 0.48 0.50 0.29 0.37 0.21 NA*

* Not Analysed

- 368 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials At 500°C, the intensity is higher and follows a Porod law (I ∝ A/q4) characteristic of large size particles. This kind of profile can be due to Laves phases at grain boundaries. The A ratio is equal to 6 ± 1, which is consistent with Fe and Mo-rich particles such as the Laves phase Fe2Mo observed by TEM. HT9 steel : At 450°C and 500°C SANS also detect the Laves phase formation confirmed by TEM observations. At 400°C, the study of the form factor shows that the radius of the particles is around 1 nm. The associated A ratio is equal to 2 ± 0.4, in agreement with the value expected for α’clusters. With the assumption of α’clusters containing 90%Cr and 10%Fe, the precipitated volume fraction is around 0.25%. For this thermal treatment, no new phase has been observed by TEM. F82H alloy : SANS and TEM results show that, whatever the ageing treatment, there is no evolution of the microstructure in comparison with the as-received sample. All the SANS intensities are identical taking into account the error bars. The scattered intensities are high at low q and correspond to large size precipitates, i.e. the M23C6 carbides initially present in the material. JLF-1 alloy : As in the case of the F82H alloy, the variation of scattered intensity in comparison with the reference sample is too weak to be considered as representative of a microstructure evolution. At 550°C, the SANS intensity seems to increase at low q with ageing, but this evolution is too weak to be interpreted. Furthermore, the A ratio measured at low q, does not change with ageing and is about 3-3.5. This value is slightly lower than for F82H alloy, although the chemical composition of the M23C6 carbides is not significantly different in these two materials. The difference of the A ratio between the steels is significant and probably representative of the presence of additional or different phases. TEM observations have shown that this material contains Ta-rich MX and V-rich M4X3 carbonitrides in addition to M23C6. Qualitatively, only the VC (A = 2.2) or V(C,N) particles can induce a decrease of the mean A ratio. Indeed, the presence of TaC carbides (A = 10) should sharply increase the A factor. So, to explain the results obtained on JLF-1, we have to assume an important contribution of vanadium carbide particles in the matrix. LA4Ta alloy : For this alloy, which presents the highest Cr content among the studied RA materials, SANS data showed that, whatever the ageing duration and temperature, the scattered intensity in the aged materials, in all the q range, is always clearly increased compared to the as-received material. This behaviour could be explained by a modification of the carbonitride precipitation, observed directly by TEM : a dissolution of the M2X precipitates, located preferentially near M23C6 (before thermal ageing), which could enhance the precipitation of the M4X3 V-rich particles. On the other hand, ageing for 10000 hours at 400°C, shows a further increase of the scattered intensity for the high q values (q ≥ 0.08 Å -1) : this must be due to the appearance of some new small size particles, which have not precipitated yet after 2000 h at 400°C.

The increment of SANS intensity from 2000 to 10000 hours ageing at 400°C is consistent with spherical nanoparticles of mean radius 1.5 nm. The corresponding A ratio is equal to 2, value near the theoretical one for the chromium-rich b.c.c. α’ particles. The volume fraction is around 0.16% if one assumes for α’ the composition 80%Cr - 20%Fe. LA12Ta alloy : This alloy is similar to the preceding, but with a lower Cr content. No change of the precipitation state was observed by TEM; consistently, no evolution of the SANS intensity was observed after 2000 h ageing, whatever the temperature, or after 10000 h at 550°C. However, a similar increase as for LA4Ta, was observed after thermal ageing at T ≤ 400°C for times ≥ 10000 hours, for the high q values (q ≥ 0.08 Å -1). This has been analysed in the same way as for LA4Ta, as a precipitation of the α’ Cr-enriched b.c.c. phase. LA13Ta : This steel has the highest W content among all the studied alloys. The comparison between the scattered intensities of thermally aged and as-received samples shows that thermal ageing does not induce significant differences in the range 250 to 450°C up to 10000 hours, in agreement with TEM observations. This is not the case for thermal ageing performed at 500°C. Indeed, in all the q domain, the scattered intensities of samples aged during 10000 hours, but also for shorter time (2000 hours), are significantly higher compared to the reference sample one. TEM observations, performed on LA13Ta steel aged during 10000 hours at 550°C, show that the precipitation of Laves phase (Fe2W type) occurs as a very thin film on grain and lath boundaries. This precipitation explains the signal observed by SANS. The chemical composition of precipitates, from microanalysis, is close to 48Fe-25Cr-2V-25W (at %), in agreement with the value of the A ratio (A ≈ 5, after subtracting the signal of the as-received sample). LA12TaLC and LA12LC alloys : As for JLF-1 and F82H materials, we did not observe by SANS any significant variation of the scattered intensity after ageing. The A ratio measured at low q is about 2.5 for the LA12LC alloy and 3 for the LA12TaLC alloy. These values are low compared to the other alloys. The A ratio is higher for the Ta rich alloy which is compatible with the presence of TaC in the matrix of LA12TaLC and not in the matrix of LA12LC. Irradiated steels F82H alloy :The SANS signal is unchanged by the irradiation, this shows that, as for thermal ageing reported above, the precipitation state is unchanged by irradiation in this material. LA4Ta alloy : In this RA material, contrary to F82H, neutron irradiation induces a large increase of the SANS intensity at high q values (q ≥ 0.05 Å -1) (see Figure 6). The increment of SANS signal between irradiated and as-received samples shows a maximum around q ~ 0.1 Å -1 : this behaviour, typical of concentrated alloys, indicates a spatial periodicity in the composition of the material, with a characteristic length 2π/q ~ 2-3 nm. The A value is in agreement with α-α’ phase separation of the ferritic matrix.

- 369 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials Two interpretations are possible : (i) isolated Cr-enriched α’ precipitates, surrounded by a spherical exclusion volume depleted in Cr, or (ii) bipercolating α-α’ two-phase system, close to the image of a spinodal decomposition. Modelling of the present SANS data is in progress. For the considered Cr concentrations (< 20%), the exclusion-sphere model seems more adequate. LA12LC alloy : This RA material shows a small increment of SANS signal after irradiation, indicative of some α’ precipitation. HT9 alloy : This conventional material shows a large increase of scattered intensity after irradiation. Taking into account its high Cr content (11.8%), this must include some α’ contribution. But, contrary to LA4Ta, the SANS increment induced by irradiation is a monotonically decreasing function of q and shows no maximum; moreover, the A ratio is not compatible with simple α-α’ phase separation. Therefore, the material contains probably another precipitated phase, which could be M6X; indeed, this phase has been observed in this material, only after irradiation, but at high temperature [3]. Effect of the Cr content on αα - αα’ phase separation For the martensitic steels containing 7.5 to 12% Cr, the major precipitated phase after tempering is M23C6 carbides, which dissolve a large amount of Cr : indeed, TEM observations show that in first approximation these carbides have the chemical formula Cr15M’8C6, with M’ = Fe, V, W or Mo. According to “thermocalc” software associated with the SGTE solution data base, the carbon concentration remaining in solid solution in the matrix after tempering ranges between 50 and 130 at. ppm. If one takes the average value of 100 at. ppm, and assumes that all the precipitated carbon is in the form of M23C6 carbides, approximate values of the Cr concentrations in the ferritic matrix of the various materials can be calculated : these, given in column b of Table 2, are found to be very close to those calculated from the “thermocalc” software (given in column a of Table 2).

The SANS results show that the Cr-rich α’ phase appears only around 400°C in the materials with a Cr concentration in the ferrite matrix close to or larger than 8.8% (LA12Ta, LA4Ta and HT9 alloys, see Table 2). For LA4Ta aged 10000 hours at 400°C, the volume fraction of α’ phase estimated by SANS is 0.16%, with a mean precipitate radius of 1.5 nm. This corresponds to an average distance between precipitates of approximately 50 nm, which is of the same order as the diffusion distance covered by Cr at this temperature. CONCLUSIONS The SANS technique has been used in order to study the microstructural evolution of martensitic steels under neutron irradiation at 325°C up to a dose of 0.9 d.p.a. and after thermal ageing between 250°C and 550°C up to 22000 hours. The investigated materials were Fe-7.5/11Cr WVTa Reduced Activation and Fe-9/12Cr MoWV conventional martensitic steels. SANS experiments were coupled with Transmission Electron Microscopy (TEM) observations. The structural informations obtained by SANS and TEM are correlated to the evolution of mechanical properties. Three main results have been obtained : (i) above 8.8%Cr in solid solution, the ferritic matrix is unstable at the ageing temperature of 400°C, and separates out into a Cr-depleted α and a Cr-rich α’ b.c.c. phases; (ii) moreover, a strong acceleration of radiation-induced α’ precipitation is observed in alloys irradiated at 325°C; (iii) the precipitation of Laves phases Fe2(Mo,W) is strongly dependent on the Mo and W content in the alloys.

Table 2 : Conditions of appearance of α-α’ phase separation in the ferrite matrix of several martensitic steels after thermal ageing at 400°C or fast neutron irradiation at 325°C : correlation with the total Cr content

and the Cr concentration in the ferritic matrix. The latter is obtained from a “thermocalc” calculation (column a) or from an approximation described in the text (column b)

Cr in matrix (at%) α/α’ phase separation

Alloy Total Cr content (at%) a) b)

Thermal ageing at 400°C

Irradiation at 325°C 0.9 dpa - 1000 hours

LA13Ta 9.1 7.5 7.0 No after 10000h F82H 8.1 7.2 7.1 No after 13500 h No EM10 9.3 8.2 No after 15000 h JLF-1 9.4 8.4 8.2 No after 13500 h

LA12TaLC 9.4 8.5 8.4 No after 5000 h LA12LC 9.5 8.7 8.5 No after 5000 h Yes LA12Ta 10.5 8.9 8.8 Yes after 10000 h LA4Ta 11.8 10.3 10.2 Yes after 10000 h Yes

HT9 12.5 10.7 10.2 Yes after 22000 h Yes *

* In irradiated HT9 steel, another phase is also present


REFERENCES [1] Mathon M. H., de Carlan Y., Geoffroy G., Averty X.,

de Novion C. H. and Alamo A. “Microstructural Evolution of Reduced Activation and Conventional Martensitic Steels after Thermal Ageing and Neutron Irradiation”, Effects of Radiation on Materials: 20th International Symposium, ASTM STP 1405, to be published.

[2] Brachet, J.C., Averty, X., Lamagnère, P., Alamo, A.,

Rozenblum, F., Raquet, O., and Bertin, J.L., “Behavior of different Austenitic Stainless Steels and Conventional, Reduced Activation (RA) and ODS Chromium Rich Ferritic-Martensitic Steels under Neutron Irradiation at 325°C for PWR environment”, Effects of Radiation on Materials : 20th. Int. Symposium, ASTM STP 1405, to be published.

[3] Dubuisson, P., Gilbon, D. and Séran, J.L.,

“Microstructural evolution of ferritic-martensitic steels irradiated in the fast breeder reactor Phénix”, J. Nucl. Mater. 205, 1993, pp 178-189.

REPORTS Mathon M. H., de Carlan Y., Geoffroy G., Averty X., de Novion C. H. and Alamo A. “Microstructural Evolution of Reduced Activation and Conventional Martensitic Steels after Thermal Ageing and Neutron Irradiation”, Progress Report UT-SM&C-LAM3, CEA report, NT SRMA 00-2396, décembre 2000.



UT-SM&C-ODS Task Title : DEVELOPMENT OF FORMING AND JOINING TECHNOLOGIES

FOR ODS STEELS INTRODUCTION The joining of Oxide Dispersion Strengthened (ODS) steels is a very important topic that has to be taken into consideration in order to develop the use of this kind of material for structural parts applications in fusion reactors. Different joining techniques can be looked into, like brazing or friction welding, but regarding the potential needs, diffusion welding by Hot Isostatic Pressing (HIP) seems to be the best solution, because of the ability of this technique to join parts of large dimensions with complex shapes. The work has started last year [1], with the elaboration of heterogeneous joints between a commercial ODS ferritic steel, MA956, and different grades of Reduced Activation Ferritic Martensitic ODS steels, ODS EUROFER, produced in our laboratory. The workprogramme of this year has been focussed on the joining of ODS steels to themselves and on their microstructural and mechanical characterization [2]. In the same time, further characterizations have been performed on heterogeneous joints by Energy Dispersion Spectrometry (EDS) and TEM observations. 2000 ACTIVITIES BASE MATERIALS A bar of ferritic MA 956 has been supplied by INCO. This bar is a special product, which has been delivered directly after extrusion, without any recrystallization heat treatment, in order to keep a small grain size (1 µm). The hardness in both the transversal and longitudinal sections is of 345 HV1. Four grades of ODS EUROFER powders have been produced in our laboratory by mechanical alloying: - EF5A : reinforced with 0,5 wt % Y2O3 in attritor; - EF10A : reinforced with 1 wt % Y2O3 in attritor; - EF5B : reinforced with 0,5 wt % Y2O3 in ball mill; - EF10B : reinforced with 1 wt % Y2O3 in ball mill; These ODS EUROFER steels have been consolidated by HIP with the following cycle: 1020°C/2h/1000bar under argon and furnace cooled (-15°C/min). The microstructure in the as-HIPed state is a milling microstructure (figure 1), which is characterized by a stacking of thin upset alloy layers (in grey) separated by prior surface oxide layers (in black).

These stackings are surrounded by the Prior Milled Particle Boundaries (PMPB), which are richer in oxides and filled with coarser reinforcement particles. The hardness measurements give values of 475 HV1 for 0,5% reinforced steels, and 505 HV1 for 1% reinforced ones, whatever the kind of mill was used.

Figure 1 : Microstructure of 1% Y2O3 reinforced EUROFER sample in the as-HIPped state

EXPERIMENTAL PROCEDURES Cylinders of Ø 30 x 30 mm have been machined from the base materials. The surfaces to be joined were machined to obtain a roughness of Ra 0,6 µm. Few samples were polished with paper grade 600, in order to remove surface oxides. All the samples have been degreased and dried with a neutral gas. No quenching or recrystallization heat treatments have been performed before the joining. However, few samples were outgassed at 1200°C for 1h under vacuum and furnace cooled. Samples were fitted in 304L stainless steel canisters, which have been previously degreased and dried like samples. Canisters were outgassed at 250°C until a vacuum of 5.10-6 mbar has been reached, and then sealed. Different couples of materials have been HIP bonded at different temperatures, for 2 hours under 1000 bar of argon pressure, followed by furnace cooling. Some characterization of the joints produced last year have been performed. Their conditions of elaboration are summarized in Table 1. New joints have been made with EF5B, EF10B and MA 956 to themselves following conditions summarized in Table 2.

20 µm


Table 1 : Joints elaborated in 1999

Joint n°

Material 1 Material 2 Surface

state HIP bonding temperature

99-E1 EF5A MA 956 Ra 0.6 1100°C

99-E2 EF10A MA 956 Ra 0.6 1100°C

99-E3 EF5B MA 956 Ra 0.6 1100°C

99-E4 EF10B MA 956 Ra 0.6 1100°C

99-E5 MA 956 MA 956 Gritt 600 1100°C

99-E6 MA 956 MA 956 Ra 0.6 1100°C

Table 2 : Joints elaborated in 2000

Joint n°

Material (to itself)

Surface state

Outgassed at 1200°C

HIP bonding temperature

00-E3 EF5B Ra 0.6 1250°C

00-E6 EF5B Ra 0.6 X 1250°C

00-E4 EF10B Ra 0.6 1350°C

00-E1 MA 956 Ra 0.6 1350°C

00-E2 MA 956 Ra 0.6 X 1350°C

RESULTS ODS EUROFER / MA 956 joints The four ODS EUROFER grades have been HIP diffusion bonded to MA 956 at 1100°C (Table 1). On all these joints, a continuous thin line made of particles is observed at the interface. EDS analysis has revealed that these particles were constituted with Al and N. Away from the interface, no differences can be seen in the microstructure of the MA 956. The hardness of the EUROFER side has fallen from 475 to 355 HV1, although the milling microstructure seems to have been not tampered. ODS EUROFER homogeneous joints The interface of the EF5B joint bonded at 1250°C without hot outgassing is filled with particles, which size ranges from 1 to 10 µm, and which form a 20 µm thick layer. These particles are believed to be oxides, since EDS analyses show that they are constituted mainly with Cr, V, Mn, and Fe. Away from the interface, the hardness has decreased from 475 to 325 HV1, and the milling microstructure has begun to disappear, but not completely. The outgassing treatment at 1200°C under vacuum has considerably decreased the amount of oxide particles at the interface, but a lot still remain. Away from the interface, the hardness and the microstructure are not modified by the treatment. The interface of the EF10B joint bonded at 1350°C without hot outgassing is also filled with oxide particles, but the layer seems less dense. Away from the interface, the hardness has decreased from 505 to 390 HV1, and the milling microstructure seems to have been completely suppressed (figure 2).

Figure 2 : Microstructure away from the interface of the EF10B homogeneous joint, bonded

at 1350°C without hot outgassing MA 956 homogeneous joints The interface of the MA 956 joint, bonded at 1100°C without previous surface grinding, looks good as quite no particles are filling it. In order to check rapidly the quality of the joint, qualitative bending tests have been performed by hand on 1 mm thick sheets. Unfortunately, the specimen breaks very easily at the interface. The joint, which surfaces were previously ground to remove surface oxides, shows a less clear interface but is slightly less brittle. The interface of the MA 956 joint, bonded at 1350°C without previous outgassing, looks good as quite no particles are filling it (figure 3). On the other joint, the treatment of outgassing at 1200°C has promoted the nucleation and coarsening of particles at the interface. Away from the interface, the hardness in the as-joined state has decreased to 250 HV1, for both cases (hot outgassed or not).

Figure 3 : Interface of MA 956 homogeneous joints bonded at 1350°C without hot outgassing

Observation made without etching

20 µm

20 µm

- 373 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials Heat treatments The specifications [3] for regular (e.g. non-ODS) EUROFER steels request a normalization treatment at 950°C followed by quenching in order to obtain a fully martensitic structure, which morphological grain structure ensures good impact properties, and a tempering treatment at 750°C in order to reduce the hardness by transforming the martensitic phase into ferrite, while keeping the same grain texture. The joint 99-E4, bonded at 1100°C, has been heat treated at 950°C for 1 hour and water quenched. Unfortunately, it seems, from TEM observations (see below), that no martensite has been formed, although the hardness of the EUROFER side, away from the interface has increased from 505 to 550 HV1. TEM Observations TEM observations have been performed on the three following states: - as-HIPped: on EF10B; - as-joined: on the EUROFER side of joint 99-E4; - as-quenched: on the EUROFER side of Joint 99-E4. On all these states a bimodal microstructure is observed (figure 4) with “large” ferritic grains with size ranging from 1 to 5 µm, and small ferritic grains with a nanometric size (< 200 nm). The main difference lies in the volume fraction of these two kinds of grains. In the as-HIPped state, the ratio is quite balanced ( 40 % of small grains). After joining at 1100°C, more than 80 % of large grains are observed. When the same sample is quenched from 950°C, the ratio is inversed as approximately 80 % of small grains are formed. Indeed, the amount of large grains seems to increase with increasing temperature of the final heat treatment.

Figure 4 : TEM observations showing the bimodal microstructure of 1% Y2O3 reinforced EUROFER

in the as-HIPped state The bimodal microstructure could be explained by the heterogeneous distribution of large reinforcement oxide particles that would refine the grains through pinning effect. Indeed, while small dispersoïds (< 30 nm) are rather homogeneously distributed, large ones (50-200 nm) are distributed along lines with a higher density of small dispersoïds. These oxide-rich lines are assumed to be the PMPB. The promotion of large grains with increasing temperature could be related to the fact that the size of austenitic grains formed during each of these heat treatments (consolidation, joining, quenching), increases with increasing temperature, thus reducing the number of nucleation sites for ferritic grains. Despite the high value of hardness reached on the quenched sample, no martensite has been observed. The absence of martensitic transformation has been confirmed by dilatometric analyses [4], and can be explained by the enhancement of diffusion kinetics, thanks to the high density of grain boundaries, as grain size never exceeds few microns. As a matter of fact, the ratio of small and large grains can explain the variations of the hardness, since it increases from 355 to 550 HV1 with decreasing temperature, and thus with increasing amount of small grains.

Table 3 : Room temperature tensile properties of homogeneous joints of ODS EUROFER

and MA 956 steels compared with the specifications for regular EUROFER steels

0.2% yield

stress (MPa)

Ultimate tensile stress

(MPa)

Elongation to rupture

(%)

Striction (%)

Fracture at the

interface

Specifications for EUROFER steels [3] 500 600 15

00-E1 (MA 956, 1350°C) 700 720 11.5 nd No

00-E6 (EF5B, 1200°C+1250°C) / 760 0.1 0 Yes

00-E3 (EF5B, 1250°C) 855 880 0.3 0 Yes

00-E4 (EF10B, 1350°C) 990 1015 0.3 0 Yes

nd: not determined, but not nil.

700 nm

- 374 - Underlying technology / Tritium Breeding and Materials / Materials development / Structural materials Tensile tests The best looking ODS EUROFER and MA 956 homogeneous joints have been tensile tested at room temperature at a strain rate of 8.10-4 s-1. As expected due to the layer of oxide particles at their interface, the ODS EUROFER joints are brittle in term of ductility (Table 3). However, their yield and ultimate tensile stresses are high compared to the MA 956 joint or to the specifications for regular EUROFER steels. Although the interface looks better after outgassing at 1200°C, the joint is more brittle, as can be observed with joints 00-E3 and 00-E6. Good properties have been obtained on MA 956 with joining at 1350°C. Moreover, the striction and the fracture have arisen away from the interface. CONCLUSION ODS EUROFER and MA 956 steels have been joined by diffusion welding to themselves. Good room temperature tensile properties have been obtained on MA 956 joints. ODS EUROFER joints have shown very good yield and ultimate tensile stresses, but quite no ductility, and failed at the interface, embrittled by the formation of a layer of oxide particles resulting from an insufficient outgassing treatment. Contrary to the joints elaborated last year at a lower temperature (1100°C), the milling microstructure of ODS EUROFER sides, characterized by the Prior Milled Particle Boundaries (PMPB), seems to have been (partly) suppressed. TEM observations have revealed a bimodal microstructure on the ODS EUROFER steels, with “large” grains with size of few microns, and small grains of nanometric size (<200 nm). Even after water quenching, the structure is fully ferritic. Depending on the austenitisation temperature of the last heat treatment, the volume fraction of small grains strongly increases with decreasing temperature (up to 80% after water quenching from 950°C). The next steps will be to investigate better outgassing treatments to improve the quality of joints, and to look into appropriate heat treatments in order to promote the martensitic transformation and suppression of the PMPB.

REPORTS AND PUBLICATIONS [1] S. Revol, C. Labonne, “Task UT-SM&C-ODS:

Development of forming and joining technologies for ODS steels”, CEA report NT DEM N°79/99, nov. 1999.

[2] S. Launois, S. Revol, C. Labonne, “Task UT-SM&C-

ODS: Development of forming and joining technologies for ODS steels”, CEA report NT DEM N°117/2000, dec. 2000.

[3] W. Dietz, A. Alamo, R. Lindau, “Technical

specification : supply of reduced activation ferritic martensitic steel type 9CrWTaV”, NET 97/917, European Blanket project – Structural Material Task 5.3, 23/07/97.

[4] S. Launois, S. Revol, C. Labonne, “Task TTMS-4.2 :

Mechanical alloying including ODS”, CEA report NT DEM N°106/2000, dec. 2000.

TASK LEADER Stéphane REVOL DRT/DTEN/SMP CEA Grenoble 17, rue des Martyrs 38054 Grenoble Cedex 9 Tél. : 33 4 76 88 98 42 Fax : 33 4 76 88 54 79 E-mail : [email protected]

annual report of the association euratom/cea 2000annual report of the association euratom/cea 2000...

Documents