CHAPTER 2 LITERATURE REVIEW

Transcription

1 51 CHAPTER 2 LITERATURE REVIEW 2.1 THE DEVELOPMENT OF ODS 9Cr STEELS FOR HIGH TEMPERATURE APPLICATIONS Ukai et al (2002) studied the hollow capsule extrusion, HIP-extrusion and CIP-extrusion processes to produce a large scale annular billet for mass production. A hollow capsule was filled with mechanically milled (MM) powders and extruded by a 2000 ton press at 1423 K. The hotextrusion experiment was conducted by using 400 ton hydraulic press. The test condition consists of extrusion speed of 30 mm/s and extrusion pressure of 616 MPa. The extrusion ratio which is the ratio of cross-section before and after extrusion is 7. Hot isostatic pressing was applied as a consolidation process of MM powders at 1423 K and 190 MPa for s. These powders were manufactured by means of an attrition type ball mill with milling period of s. A consolidation process made by hot isostatic pressing has advantage, since its hot extrusion is easier and capsule needs less tight structure. The extrusion experiment by 2000 tons press was carried out. The extrusion ratio was 6:4. Based on these results, property improvement of HIPed product by hot-extrusion should be indispensable in the course of cladding manufacturing. However, HIP-extrusion is an economical process from the practical viewpoint.

2 52 The CIP processing was conducted using rubber mould filled with MM powders along with 2 % paraffin by the application of an isostatic pressure of 600 MPa. The compacted CIPed products were sintered at 1423 K for s in vacuum to get higher true density. The MM powders were compacted in the columnar shape, but sintering occurred only in regions of direct contact between the powder particles. The true density ratio of the sintered compacted product was only 70%. It is considered that removal of pores between the particles and improvement of true density ratio can be hardly attained, even if the sintering temperature increases over 1423 K. For hot-extrusion, a true density ratio of over 90% is usually required. Therefore, it is judged from these tests that the hot extrusion of CIPed and sintered products to produce the billet is not feasible. Japan Nuclear Cycle Development Institute has successfully produced cladding tubes with 9Cr-ODS martensitic steel. It was also revealed that increase of excess oxygen (defined as total oxygen contents in steel minus oxygen contents in Y 2 O 3 ) and lack of titanium lead to particle coarsening and high-temperature strength deterioration (Ohtsuka 2004). In yet another report ODS EUROFER 97 was produced employing gas atomized powder containing 0.3 wt. % of Y 2 O 3 nanosized particles. The milled powder, canned and out gassed, was consolidated by HIP at 1373 K and 200 MPa for 2 h under Ar followed by cooling at a rate of 30 K/min keeping constant the pressure of 200 MPa during cooling (Leguey 2005, Ramar et al 2005).

3 STUDIES ON MECHANICAL PROPERTY EVALUATION OF 9Cr-ODS STEELS Extensive investigation of the continuous cooling transformation diagram of 9Cr 0.13C W 0.2Ti 0.35Y 2 O 3 reveals that a critical cooling rate of about 423 K/h gives rise to the ferrite phase at room temperature without martensite, which induces lower hardness. The hardened cladding due to accumulation of cold-rolled deformation can be successfully softened by the heat treatment with a slow cooling rate of 423 K/h after normalizing at 1323 K. The final heat-treatment consists of normalizing at 1323 K for 1 h, followed by tempering at 1023 K for 30 minutes. Applying this heat treatment technique, cold-rolling by pilger milling, was conducted under the softened ferrite structure. The martensitic 9Cr-ODS cladding showed superior ultimate tensile strength (UTS) and uniform elongation (UE) over the entire temperature range when compared with the conventional ferritic martensitic stainless steel, PNC-FMS. The ratios of UTS and UE to those of PNC-FMS are also represented in Figure 2.1. From these comparisons, the strength and ductility improvement in the martensitic 9Cr-ODS claddings are prominent, especially above temperatures around 873 K; this advantage could arise from the retardation of recovery and continuing work hardening due to pinning of the dislocations by finely distributed oxide particles. The creep rupture strength of the manufactured martensitic 9Cr-ODS cladding at 923, 973 and 1023 K emphasizes that the strength anisotropy perfectly disappears in the hoop and longitudinal directions, and the internal creep strength level approaches the target of 120 MPa for h at 973 K which is far beyond that of PNC-FMS, PNC316 beyond 1000 h at 1023 K. 9Cr-ODS cladding has a typical tempered martensitic structure with a lath size of less than 1 µm.

4 54 Figure 2.1 Tensile properties of manufactured martensitic 9Cr-ODS steel cladding as a function of temperature The fine and homogeneous grains are formed in the prior austenite grains that are finer due to grain growth retardation by the existence of Y 2 O 3 particles at the normalizing condition. Making a coarser grain formation would be a unique solution to approach the dispersion strength level on the basis of findings that accelerated deformation could arise from a grainboundary sliding among extremely fine and homogeneous grains less than 1 µm in martensitic 9Cr-ODS claddings. The irradiation performance of 9Cr-ODS and ferritic 12Cr-ODS up to 15 dpa and at a temperature range of K, the strength and ductility levels are adequately maintained. For the first wall and breeder blanket structural materials of the fusion reactor, the tentative design requires UTS of MPa at 923 K after 15 MW/m 2 to

5 55 maintain their integrity for the thermal stress, when reduced-activation ODS ferritic steels are applied. Figure 2.1 compares the tensile strength behavior of as-manufactured F82H (8Cr 0.1C 2W 0.2V 0.04Ta) and martensitic 9Cr-ODS steel tubes. The design requirement for the tensile strength at 923 K is just located at the strength lines of martensitic 9Cr-ODS steels (Ukai 2002). In a collaborative work between Japan Nuclear Cycle Development Institute and CEA-Saclay, extensive efforts were made to study the tube manufacturing processes in 9/12Cr-ODS ferritic martensitic steels. Argon gas atomized powder of Fe 0.1C 9.0Cr 2W 0.1Ti was mechanically alloyed with Y 2 O 3 powder by using the attrition type ball mill for 48 h in an argon gas atmosphere at a rotational speed of 220 rpm. The particle size of the atomized powder is less than 150 µm. The mechanically alloyed powder was then canned and degassed at a temperature of 623 K in 0.1 Pa vacuum for 2 h. Using a 400 ton press, the cans filled with the mechanically alloyed powder were hot extruded at 1323 K. Two kinds of cold-forming processes (cold drawing and HPTR rolling) were tested in the first trial to determine the suitable process to be applied and the allowable maximum cold-working level. HPTR cold rolling was consequently adopted to manufacture ODS martensitic steel cladding, since the cold-drawing process induced early cracking. The HPTR cold rolling by CEA route induces relatively low increase of hardness. The intermediate heat treatment, performed at 1073 K for 2 h in argon gas atmosphere, decreases hardness approximately by only about 20 Hv, which is lower than by furnace cooling after normalization. The final heat treatment consists of normalization at 1373 K for 15 minutes under vacuum atmosphere followed by fast cooling to induce the martensite transformation and the subsequent tempering at 1023 K for 1 h. The microstructure of both cladding tubes consists of a tempered martensite structure with homogeneous grain

6 56 morphology, indicating that the elongated grain structure produced during hot-extrusion and cold-rolling processes was destroyed and rearranged by the alpha-to-gamma phase transformations. Tensile tests in the longitudinal and transverse direction of the manufactured cladding tubes were conducted in the temperature range K. Equivalent tensile properties were found in both directions, in particular for the temperature range K. Similar strength levels in tensile and creep rupture properties were attained for the cladding tubes manufactured by JNC and CEA using HPTR rolling processes, not depending on different fabrication routes including different cross-section reduction ratios, number of passes and intermediate heat treatments. The JNC program has been in existence since 1987, and the early work concentrated on ferritic 12Cr ODS steels (Ukai 2004). In Europe, much of the development effort is on EUROFER ODS composition martensitic steel, mechanically alloyed with Y 2 O 3 powder. Tensile properties indicate that EUROFER ODS has a much higher 0.2% yield stress and excellent uniform elongation relative to conventionally produced EUROFER and F82H (Figure 2.2) (Ehrlich 2001, Klueh 2002). Preliminary room-temperature impact results are also encouraging. More recently, work has been conducted on a martensitic (9Cr 0.12C 2W 0.20Ti 0.35Y 2 O 3 ) ODS steel. 12Cr 0.3C 2W 0.30Ti 0.23 Y 2 O 3 steel produced extremely high creep strength in the longitudinal (axial) direction, but with a considerably lower strength in the hoop direction. The martensitic steel, on the other hand, had a lower creep strength, but without anisotropy. Although the room temperature yields stress is 100% greater, the impact properties degraded with transition temperature K higher than that of F82H (Klueh 2002).

7 57 Figure 2.2 The 0.2% yield strength and uniform elongation of the EUROFER ODS steel compared to other reduced-activation and conventional ferritic/martensitic steels (Klueh 2002) Mechanically alloyed EUROFER - 0.3wt% Y 2 O 3 processed via cross-roll technique, followed by appropriate thermal treatment resulted in a 50% raise in yield and ultimate tensile strength compared to the non-ods RAFM steels. The total and uniform elongation of the ODS steel are superior over the whole temperature range to that of common RAFM steels. The DBTT of ODS EUROFER lies presently in the range K (Baluc 2006). Impact tests show improved DBTT which could be shifted from +393 K for hipped ODS-Eurofer of the first generation to values well below 273 K (Möslang et al 2004, Lindau 2005, Ramar et al 2005).

8 58 The strain-controlled low cycle fatigue tests were conducted at 873 K, 923 K, 973 K and 1023 K. It was shown that 9Cr ODS martensite and 12Cr ODS ferrite exhibited similar fatigue behaviour and both were demonstrated to be superior to the conventional 12Cr ferritic-martensitic steels (Ukai and Ohtsuka 2005, Ramar et al 2005). 2.3 EFFECT OF OXYGEN AND TITANIUM/SOLUTE ADDITIONS ON MECHANICAL PROPERTIES OF 9 Cr ODS STEELS Fe 0.13C 9Cr 2W 0.2Ti 0.35 Y 2 O 3 alloy manufactured by fourpass cold rolling using a pilger mill cladding has superior 0.2% yield stress, associated with larger amount of 0.35 mass% Y 2 O 3 and 0.2 mass% Ti. The final heat treatments consisted of normalizing at 1323 K for 1 h and tempering at 1073 K for 1 h. Stepwise increase of creep rupture strength could be attributed to the increase of Y 2 O 3 and Ti addition, which results in the reduced size of Y Ti O complex oxide and concomitantly the reduced spacing between oxide particles (Ukai 2004). The Fe 9Cr 2W 0.2Ti 0.13C alloy should transform to gamma during the 1423 K heat treatment, referring to the phase diagram; however, the 9Cr-ODS steel shows a peculiar transformation behavior that the alphaphase remains even at 1423 K. The carbon depletion in matrix due to titanium carbide formation should decrease the driving force for the alpha to gamma transformation because titanium is added in 9Cr-ODS. The pulling force against gamma grain boundary migration by oxide particle dispersion (Zener effects) consumes the driving force while the grain boundaries are moving over oxide particles. These two factors would restrict the alpha gamma transformation and cause the residual alpha generation at 1423 K in 9Cr-ODS (Ohtsuka 2004). Ohtsuka et al s (2004) report on the effects of different

9 59 percentages of excess oxygen and titanium on the mechanical properties and microstructure of 9Cr-ODS steel can be summarized as follows: 1) Controlling the atomic ratio between excess oxygen and titanium (x in TiO x ) to around 1.0 produces elongated residual alpha-grains, which possess ultra-fine and close oxide particle distribution. 2) It is considered that the number density of oxide particles in equiaxed grains would be maximized when x in TiO x is around 1.0, where oxide particles are most finely distributed and the required amount of excess oxygen for Y Ti complex oxide is maintained in the steel. 3) It is indispensable for achieving the superior high temperature strength to control x around 1.0 and promote the formation of residual alpha-grains and high number density of oxide particles in equiaxed grains (Ohtsuka 2004). The creep strength of 9Cr ODS can be improved by appropriately controlling hot-extrusion temperature and chemical compositions (oxygen and titanium concentrations) and that the creep strength improvement should be correlated with number density of oxide particles and volume fraction of elongated grains containing ultrafine oxide particles. 9Cr ODS bars containing different concentrations of oxygen and titanium were produced by MA and hot-extrusion process. The hot extrusion temperature was set at 1423 K and 1473 K. Oxygen and titanium concentrations were controlled to 0.10~0.24 wt% and 0.20~0.46 wt%, respectively for attaining optimum properties (Ohtsuka 2005, Ramar et al 2005).

10 60 1) In the case of Si or Al addition, oxides disappeared after MA process, which means Y 2 O 3 and other elements should be in solution at non-equilibrium condition. Two types of oxides of Y 2 O 3 and Al 2 O 3 or SiO 2 developed after the annealing at 1123 K, but only complex oxides were developed after the annealing at 1423 K. This result suggests that the oxide formation is independent process for Y and Si or Al. 2) In the case of Ti addition, oxides disappeared after MA process, but developed after annealing at 1423 K. This means that Ti can stabilize complex oxides of Y and Ti, and enhance the fine distribution of the oxides comparing with simple Fe-9Cr- Y 2 O 3 alloy (Ohnuki 2005, Ramar et al 2005). 3) The addition of Ti is an effective method of achieving uniformly distributed ultra-fine oxide particles (Sakuma 2004). This result suggested that the MA process can induce some super-saturation of oxide-forming elements, which is termed a mechanically forced super-saturation, which can produce precipitation following the HIP process. In addition, W is introduced as a solution hardening element (Kaito 2004). When Ti-added steel is mechanically alloyed with both yttrium oxide and excess oxygen, complex oxide of Y, O, -TiO, or Y, O, -2Ti0, is formed and is finely dispersed as dispersoids. The diameter of dispersoids is concomitantly reduced from 18 nm in the raw yttrium oxide powder to a few nm, which extensively enhances the effect of dispersion strengthening. The manufactured claddings within the specification limit exhibited a superior high temperature strength and sufficient Charpy impact properties (Ukai 1993).

11 61 The effects of the partial recrystallization and the cooling rate during the normalizing treatment on the mechanical properties were investigated on 9% Cr ODS steel samples prepared by a mechanical alloying of the elemental powders of Fe, Cr, V, Ti together with Y 2 O 3 powder, a hipping at 1423 K/10.3MPa (1500 psi) for 4 hr and subsequent hot rolling at 1323 K. The one-half recrystallized sample showed better mechanical behavior than the fully recrystallized (normalized) one at the intermediate temperature range (773 to 973 K). Equiaxed grains were observed in the HIPed samples, but the mean grain size of the 0.2% V added sample was two orders of magnitude smaller than that of the other one which did not contain V. On the other hand, the hot rolled samples showed similar grain sizes irrespective of the V content. M 23 C 6 carbides were present on the grain boundaries of both samples after any process (Jinsung Jang 2005, Ramar et al 2005). The effect of dispersoid content on the properties of ODS-Eurofer 97 is as follows: Scanning electron microscopy of the surface appears smooth in the case of the 0.3% yttria, while the 0.5% yttria reveals pores that are about 10 m in size. Both ODS steels present interlocked ferrite grains with chromium carbides at the boundaries. In the case of the 0.5% yttria, the carbides at the grain boundaries with sizes up to about 1 m are larger than in the case of the 0.3% yttria, where they attain about 0.2 m maximum. The grain size of about 5 m in both ODS steels is smaller than in the case of the base tempered martensite EUROFER 97, which presents grains of about 15 m in size that are decomposed in a martensite lath structure. No clear microstructural difference could be identified between the transverse and the longitudinal cuts, as for the tensile test response. Yttria particles appeared in the material in either groups of about nm round particles or groups of 1 5 nm round particles. They appear heterogeneously distributed, with

12 62 regions free of them. EUROFER 97 steel containing 0.3% yttria has superior tensile resistance and similar, if not better, uniform elongation behaviour than the base material; in the case of the 0.5% yttria presents weak mechanical properties. This fact is certainly related to the presence of the porosity observed in SEM and to the large carbides observed in TEM. The concentration of yttria seems to be of importance for the mechanical properties of the ODS alloy. The case of the 0.5% yttria is not so promising, with a critical stress similar or lower than the one of the base material, and similar or lower elongations (Schaeublin 2002). The critical hydrogen concentration required to induce inter granular cracking in ODS steels was in the range of wppm that is almost one order of magnitude higher value than that of 9Cr martensitic steels. The high critical concentration of hydrogen in the ODS steel was interpreted in terms of high capacity of hydrogen trapping caused by very fine grain size and introduction of yttria particles in the steel matrix (Yun Chen 2006, Bloom et al 1999, Kimura, 2007). Uniform strain does not disappear at 293 K in the high strength ODS-steels below 1500 MPa as observed for irradiation-strengthened 7 9CrW (Ge)V Ta steels as shown in Figure 2.3. Pronounced precipitation (from increasing N and C contents) or fine Y 2 O 3 - dispersions in ODS-(9 13) Cr steels strongly increase work-hardening and uniform ductility but decrease fracture strain and USE of impact tests (Preininger 2002, Kasahara 2002).

13 63 Figure 2.3 Creep rupture stress Vs Rupture time for ODS alloy 2.4 OXIDATION AND CORROSION RESISTANCE OF 9Cr-ODS STEELS The formation of YFeO 3 from metallic iron, yttrium oxide and oxygen is thermodynamically possible, especially in environments at relatively low temperatures and low oxygen content. This suggestion is also supported by the investigation carried out by Zeng et al who studied the oxidation behavior of Fe Y alloys containing 2 5 at % Y at K in air and reported the existence of YFeO 3 phase in the inner mixed oxide layer which significantly enhanced inward oxygen diffusion and decreased outward iron diffusion. Most authors conclude that alloys containing yttrium are oxidized by inward diffusion of anions at elevated temperatures.

14 64 The 9Cr ODS steel provides large amount of short-circuit transport paths for both anion and cation diffusion. Initially, the Y Cr-rich oxides preferentially nucleate and grow as an interconnected network at the grain boundaries of the base steel, owing to the high oxygen affinities of chromium and yttrium compared with that of iron. These grain boundaries act as the major path for cation diffusion for subsequent oxidation process. The formation of these ribbons reduced the cation flux at these regions, while not decreasing the inward oxygen diffusion, consistent with the work of Zeng. Therefore, the species that diffused at the highest rate in the system changed from cations to anions, and the location of the reaction to form new oxide changed from the scale/oxygen interface to the steel/scale interface. As a result, the overall corrosion rate of the 9Cr ODS decreased and the scales formed on this steel are thinner. Primary oxide growth occurred in the internal oxidation layer and was controlled by slow oxygen diffusion. This differs from ferritic martensitic steels with similar Cr contents where growth is primarily due to faster iron cation diffusion that expands the magnetite and spinel layers. 9% Cr ODS ferritic steel with yttrium present as Y Ti O dispersions exhibited lower oxide weight gain compared to conventional non-ods 9% Cr ferritic/martensitic steels when exposed to supercritical water at 773 K and 25 MPa with a 25 ppb dissolved oxygen concentration. The improved oxidation performance is caused by the modification of ionic transport due to the segregation of the rare earth elements to the oxide grain boundaries and to the oxide metal interface, although some investigations reported no such segregation. Rare earth elements, such as yttrium added either elementally or as oxide dispersoids in alloys influence oxide growth rate in Fe-base alloys (Yun Chen 2006).

15 65 For 9Cr-ODS martensitic steel and 12Cr-ODS ferritic steel with fine grains, the weight gain due to oxidation under a controlled dry air atmosphere is quite small and comparable to that of PNC316 containing 17 mass % Cr. Their weight gain is limited to below 0.1 mg/mm 2, with almost no dependence on temperature and time (Figure 2.4). It is known that addition of rare earth elements, i.e., Y, La, Ce, suppresses the oxidation of steels. The internal oxide of rare earth elements, that are the most thermodynamically stable, may adhere to the protective a-cr 2 O 3 (keying-on effect). The superior oxidation resistance of ODS steel is attributed to early formation of the protective a-cr 2 O 3 at the matrix and scale interface, which may be stabilized by finely dispersed Y 2 O 3 particles (Kaito 2004). Figure 2.4 Weight gain by oxidation Vs testing time for ODS steels and FM steels ODS steels (9 12 wt% Cr) showed almost similar corrosion behavior with ferritic/ martensitic steel in the SCPW (25 MPa, 793 K). Passive current density of the ODS steel is one order of magnitude lower than that of ferritic steel at room temperature. The weight gain of the 12Cr-ODS steels is almost the same as that of a ferritic/martensitic steel and about 1 order of magnitude larger than for SUS316 (Cho 2004).

16 66 The 9Cr ODS alloy showed the lowest weight gain due to oxidation, even though the 9Cr ODS alloy had less bulk Cr (9 wt% Cr) than HCM12A (12 wt% Cr), (see Figure 2.5) after exposure to supercritical water at 773 K with 25 ppb dissolved oxygen showed a typical dual layer, which is composed of an outer Fe-O magnetite layer and an inner Fe-Cr-O spinel layer and superior corrosion resistance of the 9 Cr ODS alloy relative to other ferritic-martensitic steels (Was and Matsui 2005). Figure 2.5 Weight gain due to oxidation The alloys developed showed resistance to corrosion at temperatures up to 923 K and above for both static and dynamic-isothermal experiments. In oxygen containing environments, a two layer film forms consisting of an outer chromia layer and internal silica based layer that effectively terminates the corrosion process (Lim 2003). 2.5 IRRADIATION EXPERIMENTS REPORTED ON 9Cr-ODS STEEL A ferritic steel T91 and a 9Cr-ODS martensitic steel were irradiated with 5 MeV Ni ions at 773 K at a dose rate of dpa/s to doses of 5,

17 67 50 and 150 dpa. For T91, the irradiated microstructure was dominated by tangled dislocation and precipitates, similar to the unirradiated condition except the presence of large dislocation loops of type a<100>. The microstructure of alloy 9Cr-ODS for both the unirradiated and irradiated cases were dominated by dense dislocations, precipitates and yttrium oxide particles and no dislocation loops were observed. The average size of yttrium oxide particles slightly decreased with dose from 11.8 nm for the unirradiated to 9.1 nm at 150 dpa. No voids were detected for both alloys up to a dose of 150 dpa (Gan et al 2003). ODS steels were fabricated as follows with a changing gas environment during MA, heat-treatment condition and chemical composition and electron-irradiated to 12 dpa at K. 1) MA in an Ar gas environment and recrystallization at 1423 K for 1.8 ks (F-Ar0). 2) MA in an Ar gas environment and cold rolled 20% after recrystallization at 1423 K for 1.8 ks (F-Ar20). 3) MA in a He gas environment and recrystallization at 1423 K for 1.8 ks (F-He0). 4) MA in a He gas environment and cold rolled 20% after recrystallization at 1423 K for 1.8 ks (F-He20). 5) MA in an Ar gas environment and normalization at 1323 K for 3.6 ks (M-Ar).] An ODS martensitic steel (M-Ar) with high dislocation density showed very good swelling resistance. Swelling levels of ODS ferritic steels depended on the gas environment during MA and the recrystallization

18 68 condition. These indicated that a helium gas environment during MA was more effective to suppress swelling than an argon gas environment and that cold working after recrystallization reduced void formation and swelling. The chemical composition for the ODS ferritic steels was 0.06C-11.8Cr-0.26Ti- 0.23Y 2 O 3 (wt%) and that for the ODS martensitic steel (M-Ar) was 0.13C-9Cr-0.2Ti-0.3Y 2 O 3 (wt%). There were obvious differences in the size of oxide particles among F-Ar0, F-He0 and M-Ar; the particle sizes in F-He0 were lower than those in other alloys that were MA processed in an argon gas atmosphere. Voids in F-Ar0 had begun to nucleate at a low-dose level (<4.8 dpa), and grew continuously with increasing irradiation dose. Void formation in F-He0 was suppressed compared with that in F-Ar0. Except for F-Ar0, these swelling levels were lower than that of typical austenitic steels at the same conditions. Voids formed at the oxide particles indicated that the oxide particles had acted as point defect sinks and nucleation sites for defect clusters during electron-irradiation. The oxide particles might contribute to enhanced recombination to decrease point defect concentration in the matrix during electron-irradiation. This might cause the lower swelling in F-He0. Application of helium gas as MA environment and the cold working after recrystallization seem to be effective to reduce void swelling. Swelling was also affected by the size and the density of the oxide particles (Yamashita, 2000). EUROFER97 ODS steels were irradiated with 590 MeV protons for doses ranging from 0.3 to 2 dpa at room temperature and at 623 K. After irradiation ODS EUROFER97 showed a slight change in the mechanical properties. The presence of voids, dislocation loops are observed in the samples irradiated at 623 K, whereas very small sized defects with sizes from

19 69 2 to 3 nm are observed on the samples irradiated at room temperature. The dispersed yttria particles are found to be stable and undissolved on irradiation over all the temperature range and irradiated dose (Ramar et al 2005). The scope of the present study is the indigenous synthesis and development of ODS 9Cr RAFM steel by MA capable of operation as a structural material for fast fission and future fusion reactors.