Development and characterization of advanced 9Cr ferritic/martensitic steels for ?ssion and fusion reactors
S.Saroja a ,?,A.Dasgupta a ,R.Divakar a ,S.Raju a ,E.Mohandas a ,M.Vijayalakshmi a ,K.Bhanu Sankara Rao b ,Baldev Raj a
a Indira Gandhi Centre for Atomic Research,Kalpakkam 603102,Tamil Nadu,India b
School of Engineering Sciences,University of Hyderabad,Hyderabad,India
a r t i c l e i n f o Article history:
Available online 12October 2010
a b s t r a c t
This paper presents the results on the physical metallurgy studies in 9Cr Oxide Dispersion Strengthened (ODS)and Reduced Activation Ferritic/Martensitic (RAFM)steels.Yttria strengthened ODS alloy was syn-thesized through several stages,like mechanical milling of alloy powders and yttria,canning and consol-idation by hot extrusion.During characterization of the ODS alloy,it was observed that yttria particles possessed an af?nity for Ti,a small amount of which was also helpful in re?ning the dispersoid particles containing mixed Y and Ti oxides.The particle size and their distribution in the ferrite matrix,were studied using Analytical and High Resolution Electron Microscopy at various stages.The results showed a distribution of Y 2O 3particles predominantly in the size range of 5–20nm.A Reduced Activation Ferritic/Martensitic steel has also been developed with the replacement of Mo and Nb by W and Ta with strict control on the tramp and trace elements (Mo,Nb,B,Cu,Ni,Al,Co,Ti).The transformation temperatures (A c1,A c3and M s )for this steel have been determined and the transformation behavior of the high temper-ature austenite phase has been studied.The complete phase domain diagram has been generated which is required for optimization of the processing and fabrication schedules for the steel.
ó2010Elsevier B.V.All rights reserved.
1.Introduction
The sodium cooled fast breeder reactors require development of high temperature radiation resistant materials for achieving high fuel burn-up of the order of 200GWd/t or higher (>200dpa).Such high burn-up is the key factor for the ef?cient and economical operation of the reactors.The performance of core structural mate-rials namely clad and wrapper is limited by high temperatures and intense neutron irradiation,where the neutron ?ux levels are high-er (1015ncm à2s à1),than that in thermal reactors [1].The clad tubes should possess excellent resistance to void swelling and irra-diation embrittlement and must be capable of operating under stress at operating temperatures for long periods of time [2,3].Ferritic/Martensitic steels (9–12%Cr)although exhibit higher void swelling resistance than austenitic steels,have poor high temper-ature creep strength,which limits the operating temperatures to $823K.One of the methods to overcome this limitation is to intro-duce thermally stable oxide dispersoids in the ferrite or ferrite/martensite matrix.The dispersed ?ne Y 2O 3oxide particles block the mobile dislocations and act as trapping sites for point defects
induced by irradiation.Thus,they not only retard irradiation swell-ing but also improve the high temperature strength.
In the last two decades or so there have been considerable ef-forts the world over on development of ODS steels with high tem-perature strength [4–6].Most of the efforts are focused on synthesis of the alloy by powder metallurgical processing methods.The addition of a small amount of Ti promoted the formation of very ?ne coherent mixed oxide particles of yttrium and titanium,thus signi?cantly improving creep rupture strength.The fabrica-tion of clad tubes has been taken up in Europe,USA and Japan by employing mechanical milling of powders,hot extrusion and cold rolling [7–9].The alloys that have received considerable attention for production of clad tubes are of the 9Cr martensitic steel such as Fe–9Cr–0.13C–2W–0.21Ti–0.36Y 2O 3and the 12Cr ferritic variety such as Fe–12Cr–0.025C–2W–0.3Ti–0.25Y 2O 3.The cladding tubes of 12Cr exhibited strength anisotropy,with superior strength in the longitudinal direction to rolling but inferior strength in the per-pendicular direction.On the other hand,the 9%Cr martensitic ODS steels exhibited nearly isotropic mechanical properties after ?nal heat treatment [10,11].Further,the 9Cr-ODS alloys show the lowest DBTT shift on irradiation as compared to other Cr composi-tions and also better formability characteristics over 12Cr ODS steel that is bene?cial in producing them by cold-pilgering [12].In view of the above,efforts have been made to develop ODS alloy
0022-3115/$-see front matter ó2010Elsevier B.V.All rights reserved.doi:10.1016/j.jnucmat.2010.09.022
Corresponding author.
E-mail address:saroja@https://www.doczj.com/doc/2310512306.html,.in (S.Saroja).
of composition9Cr–2W–0.11C–0.2Ti–0.35Y2O3.This paper pre-sents the results on characterization of microstructural and micro-chemical features which include particle morphology and size distribution of the yttria phase at various stages of synthesis.
It is well known that the conventional9Cr–1Mo ferritic–mar-tensitic steels and its variants meet most of the requirements for applications below773K in liquid metal cooled fast reactors.How-ever,the emerging scenario to develop fusion power has put for-ward a demand to develop newer grades of steels,which are not only resistant to neutron irradiation of higher energy but also pose a reduced hazard from the induced radioactivity arising due to transmutation reactions.This has led to an extensive research and development program worldwide on fusion reactor materials leading to the development of low or Reduced Activation Fer-ritic–Martensitic steels(RAFM)[13–20].The design principles for
the RAFM steels are essentially same as that of conventional mid-dle or high chromium low carbon alloys such as V,Nb modi?ed 9Cr–1Mo steel,but with a major variation with respect to elements producing long half-life transmutants like Ni,Mo,Nb,Cu,Co,Al,N, etc.These elements are largely substituted by their comparatively lower activation counterparts,such as Mn,W,V,Ta,and C,so that the RAFM steel at the end of its useful lifespan of about100years can be safely handled after a shorter cooling period,rather than a cooling time of over1000years,which would have been the case with conventional Cr–Mo–Ni–Nb containing ferritic/martensitic steels[13,20].
The indigenous development of Reduced Activation Ferritic–Martensitic(RAFM)steel,considered as viable structural materials for the Test Blanket Module(TBM)of International Thermonuclear Experimental Reactor(ITER)has also been taken up in co-operation with Indian industries.By appropriate selection of raw materials and melting routes the9Cr–1W–Ta–V–0.1C with controlled chem-istry has been achieved.A systematic physical metallurgy study to characterize the phase diagram,transformation temperatures and microstructural evolution during thermal aging has been carried out[21].Some of the results are presented in this paper.
2.Experimental
The ODS alloy and RAFM steel have been characterized at vari-ous stages using a variety of techniques.The ODS alloy in the pow-der and intermediate stages including a model Fe–Ti–Y2O3alloy have been studied using X-ray diffraction(XRD),Scanning Electron Microscopy(SEM)and Transmission Electron Microscopy(TEM)to determine structure,morphology and phase distribution.XRD was carried out in an Inel XRG-3000diffractometer with Cu K a1radia-tion.The powder was dispersed on conductive tape for the SEM analysis.SEM studies were performed in a FEI ESEM model XL-30 operated at30kV in secondary electron and back-scattered elec-tron imaging modes for morphological and topological character-ization and energy dispersive X-ray spectroscopy for chemical analysis.
The powder for TEM studies was prepared by mixing it thor-oughly in a thermal-set,low vapour pressure epoxy,?lling within a3mm inner diameter Te?on ring and curing at403K for $60min.This was subsequently ground,polished,dimpled and ion-milled as for conventional TEM samples.TEM studies were car-ried out in a FEI CM200Analytical Transmission Electron Micro-scope and a JOEL2000High Resolution TEM.A few select specimens were studied in FEI Tecnai G2F30operated at300kV TEM at Indian Institute of Science,Bengaluru,India.
The composition of the RAFM steel used in this study is listed in Table1.The steel was produced by vacuum induction melting fol-lowed by vacuum arc re?ning and supplied in the form of25mm plates in the normalized and tempered condition by MIDHANI,Hyderabad,India.Specimen of dimensions10?10?5mm were sliced from the plate for heat treatments and subsequent evalua-tion of thermal stability and thermal property measurement stud-ies.The steel has been normalized at1253K for30min to ensure the complete dissolution of M23C6type carbides and the possible growth of austenite grains at high temperatures,which would pro-mote the formation of martensite upon subsequent air cooling.The steel was tempered at1033K for60min.In order to simulate the microstructure of the heat affected zone in a weldment the steel was solution treated at different temperatures in the range of 1253–1553K and cooled at different rates.The heat treated RAFM steel specimens were prepared metallographically to study the microstructural features.The nature of phases was studied using Transmission Electron Microscopy on thin foils and carbon extrac-tion replica.
For the Differential Scanning Calorimetry(DSC)experiments cut pieces of nearly identical shapes of mass varying from50to 100±0.1mg were used.The DSC experiments were performed in a Setaram Setsys16òheat-?ux type high temperature differential scanning calorimeter,employing recrystallised alumina crucibles of about100l L volume under a constant?ow(50mL minà1)of high pure argon.The description of the equipment and the calibra-tion procedure has been given elsewhere[22].For the estimation of various on-heating transformation temperatures,slow heating rates of1–5K minà1were adopted,since higher heating rates re-sulted in appreciable nonlinear shift of the measured thermal ar-rest points from their equilibrium values.In order to evaluate the martensite start M s and martensite?nish M f temperatures a few samples were cooled from1253K using high cooling rates ($100K minà1).The basic output of a DSC thermogram,namely the heat?ow is calibrated in terms of a reference signal obtained using pure iron reference.
3.Results and discussion
3.1.Characterization of microstructure during mechanical alloying and consolidation of ODS alloy
The performance of ODS alloy is crucially dictated by size,spac-ing and distribution of the oxide dispersoids in the matrix,which in turn depend on the chemical composition and temperature of processing.Yttria,being a ceramic oxide,is hard and stable phase at high temperatures(up to about1473K),and is the most widely used dispersoid in steels.Excessive contents of yttria will make the steel very brittle.A small amount of yttria($0.3wt%)is considered essential to optimize mechanical properties.Such small amount demands that the particles possess very?ne size($few10nm) in order to achieve a uniform distribution.Under this condition, the steel exhibits excellent creep strength by pinning mobile dislo-cations at elevated temperatures of operation.The presence of Ti in Table1
Comparison of chemical compositions(wt%)of RAFM steel with Eurofer97.
Element/steel Indian
RAFM
Eurofer97Element/steel Indian
RAFM
Eurofer
97 Cr9.048.50–9.50B0.0005<0.001 C0.080.09–0.12Ti<0.005<0.01 Mn0.550.20–0.60Nb0.001<0.001 V0.220.15–0.25Mo0.001<0.005 W 1.00 1.00–1.20Ni0.005<0.005 Ta0.060.05–0.09Cu0.001<0.005 N0.02260.015–0.045Al0.004<0.01 O0.0057<0.01Si0.09<0.05 P0.002<0.005Co0.004<0.005 S0.002<0.005As+Sb+Sn+Zr<0.03<0.05
132S.Saroja et al./Journal of Nuclear Materials409(2011)131–139
this class of steels improves the strength even better.The superior properties are achieved by re?nement of the Y2O3to?ner(3–5nm) sized particles by formation of coherent yttria-titania complex oxi-des.These oxide particles are produced through the interaction of yttria with Ti in presence of excess oxygen during the processing of the steel.The yttria dispersed($0.3wt%)ferritic–martensitic steels (Fe–9Cr–0.1C–2W–0.2Ti)exhibit better mechanical properties than the matrix material even up to973K.The above facts indicate that the uniformity of microstructure and size distribution of dis-persoid is the decisive factor in realizing the performance of the steel,the studies on which are described below.
Fig.1a shows the secondary electron micrograph of the pre-al-loyed steel powder(9Cr–2W–0.2Ti–0.1C).It represents globular particles of size<200l m which also corresponds to the sieve size. Fig.1b shows the microstructure of the attritor milled powder of the pre-alloyed steel with0.3wt%of yttria(initial size$40nm). The duration of milling was4h.A plate like morphology of the milled powder is observed.The plate morphology is indicative of the high ductility of the pre-alloyed powder,while the structure in general can be referred to as a‘well blend lamellar’structure. The difference between a well blend structure and a convoluted lamellar structure are brought out in Fig.1c and d,respectively. Fig.1c corresponds to an attritor milled powder(milled for4h) while Fig.1d corresponds to an incompletely ball milled powder (milled for48h).A well blend structure represents a situation where the lamellar spacing approaches the dispersion size.In this way,it was found that attritor milling is more ef?cient than ball milling.Optimization of the attritor milling time of the powders was also carried out in order to assess the minimum milling time that is required for obtaining the desired microstructure.It was ob-served that3h of milling is suf?cient to obtain a well blend micro-structure and a uniform distribution of yttria particles in the matrix.
The milled powders was then degassed and canned and com-pacted by upset forging,after which they have been consolidated by hot extrusion.The extrusion was carried out in the austenite phase domain in the range of1323–1423K and the optimum tem-perature at which the yttria particle size was?nest and distribu-tion was uniform was found to be1423K.Fig.2shows the TEM dark?eld images from a rod extruded at this temperature. Fig.2a shows martensite laths measuring about100nm in width, and Fig.2b shows the distribution of the yttria dispersoid in the martensite matrix.The distribution of yttria particle sizes was found to be skewed at5nm,and most particles were within 20nm.The presence of a large fraction of?ne particles is consid-ered as a good size distribution of the dispersoid for achieving the desired high temperature mechanical properties.
Tubes were fabricated from the extruded rods followed by a normalization and tempering treatment.Fig.3shows the bright ?eld TEM micrograph of a tube section.It reveals a typical tem-pered martensitic structure with carbide precipitates decorating the lath and prior austenite grain boundaries,which con?rms that the normalizing and tempering heat treatments have resulted in the expected microstructures typical of9Cr ferritic–martensitic steels[23–26].
3.2.Study of dispersoid characteristics in a model alloy
The characteristics of the dispersoid have been studied in a Fe–0.3Y2O3–0.2Ti model alloy.Fig.4a shows the radially averaged intensity of the Selected Area Electron Diffraction Pattern shown in the inset.The intensity pattern was obtained by arbitrary back-ground correction which may have resulted from various mecha-nisms of inelastic scattering.It was concluded that the yttria phase is bcc due to the presence of diffraction from(211)plane. Fig.4b shows the TEM lattice image of the matrix with an
(a) (b)
(c)(d)
S.Saroja et al./Journal of Nuclear Materials409(2011)131–139133
embedded dispersoid nano-particle in this model alloy.The matrix–particle interface is seen to be nearly defect free.The ele-mental distribution of Fe,Ti and Y from a typical region of the sam-ple using Fe-K a ,Ti-K a and Y-K a +Y-K b lines are shown in Fig.4c–e ,respectively.The ?gures indicate that Y is almost always accompa-nied by Ti although a signi?cant amount of Ti is left out.The result implies that the dispersoid contains both Y and Ti and is an yttria-titania complex although oxygen map could not be obtained with suf?cient reliability owing to the presence of residual oxygen in the matrix.Formation of Y–Ti–O type of nanoclusters have been re-ported in the literature [27].
The salient features of the studies on the ODS alloys at different stages of synthesis and fabrication are as follows:
Yttria has a bcc crystal structure and no lattice parameter change is observed due to addition of Ti.
The matrix–dispersoid interface does not show signatures of lattice strain.
By EDS X-ray mapping an association of Y with Ti has been observed.
A uniform distribution of dispersoid was achieved at higher extrusion temperature.
60min respectively based on the variation of prior austenite grain size and hardness of the steel with temperature and time.
In order to study the nature of phases at high temperature,the steel was solution treated at different temperatures in the range of 1253–1553K and water quenched.Examination of the microstruc-ture showed prior austenite grain boundaries and a fully martens-ite product,when quenched from 1253K to 1323K (Fig.5a and b).The hardness of the product was $420VHN,con?rming the mar-tensitic nature of the product.Prior austenite grain size increased from an average value of about 20l m at 1253K to about 37l m at 1323K.The microstructure at 1453K (Fig.5c)did not show a high increase in prior austenite grain size,contrary to the expecta-tion.This suggests that the high temperature d -ferrite is present at 1453K.Presence of ?ne d -ferrite in the form of stringers is ob-served,which explains the absence of grain growth at high temper-ature.The presence of coarse d -ferrite grains is observed at 1553K (Fig.5d).The chemical composition determined using energy dis-persive spectroscopy of X-rays in the d -ferrite and austenite (mar-tensite)showed a higher amount of Cr and W in d -ferrite,suggesting the repartitioning of solutes during solution treatment [28].By quantitative metallography,about 20%d -ferrite was eval-uated at 1553K.The above studies suggest that the onset of d tran-sition is below 1453K.This can be understood based on the presence of ferrite stabilizers like W and Ta and low amount of C
and N in the steel [29].This study has provided a microstructural data base which would help to understand the phase transforma-tions that the steel would undergo due to thermal cycling in the heat affected zone of a weldment.
Slow cooling from the austenite phase domain resulted in a structure of pro-eutectoid ferrite and martensite.No variation in microchemistry between ferrite and austenite (martensite)was observed,due to the slow kinetics of solute repartitioning at low temperatures during cooling.
The steel after normalizing and tempering showed a typical tempered martensite structure.TEM micrograph of the normalized and tempered steel showed the formation of subgrains along with inter and intra lath carbides (Fig.6a and b).The presence of coarse globular carbides of sizes <250nm on the boundaries and ?ne len-ticular precipitates of size ranging from 20to 25nm within the laths was observed.The coarse precipitates (Fig.6c)have been
identi?ed as M 23C 6by electron diffraction (inset in Fig.6c.These carbides are rich with Cr with a large solubility for W (Fig.6d).Using a combination of electron diffraction and energy dispersive analysis of X-rays,the ?ne carbides on the dislocations inside the laths were identi?ed as Ta and V rich MX type of precipitates (Fig.6e and f).The evaluation of high temperature mechanical properties is in progress.
3.4.Determination of transformation temperatures in RAFM steel This section describes the determination of temperature re-gimes for different phase domains using differential scanning cal-orimetry technique.A typical DSC pro?le obtained during heating and cooling of a normalized sample,employing a rate of 5K min à1is shown in Fig.7a and b.An initial microstructure of martensite obtained by normalizing the steel at 1253K was used in this
study.
averaged intensity of the Selected Area Electron Diffraction Pattern shown in the inset;(b)lattice image of the matrix with an embedded elemental map from a representative area corresponding to Fe-K a ,Ti-K a and Y-K a +Y-K b ,respectively.It is observed that the yttria dispersoid homogeneous boundary with matrix and that it is almost present together with Ti inferring that the dispersoid is a yttria-titania complex.
The presence of various thermal arrest events which include solid state transformations and a very prominent event namely melting, is clearly brought out in this?gure.The temperature values corre-sponding to the different transformation arrest points are listed in Table2.Since the enthalpy effects associated with samples of low mass(<50mg)for various solid state phase changes are small, these thermal arrests are determined by taking the derivative sig-nal of the heat?ow,in which the occurrence of the in?ection points are more clearly delineated.The on-heating scans are per-formed suf?ciently slowly so that the equilibrium phase stability domains prevail,and are used for indexing the DSC pro?le.It should be noted that simulation of the high temperature phase at room temperature by interrupted quenching is not possible in such steels,where the M s temperature is quite high and almost invariably a martensite product is obtained.Analysis of Fig.7sug-gests the sequence of phase changes that takes place during slow heating of the RAFM steel.The thermodynamic calculations on phase stability in related steels[30,31]are used to arrive at the transformation temperatures(Table2)in the present experiment.
It is seen from Table2that the A c1and A c3temperatures are 1104and1144K respectively.The phase transformation tempera-tures that are obtained in this study are in general agreement with the values reported for9Cr–1Mo based low carbon steels[32]. These temperatures have also been obtained for various initial microstructures[22]such as normalized,normalized and tem-pered for different durations.It is observed that the values of A c1 and A c3temperatures for different initial tempered microstructures of RAFM steel do not vary appreciably.In other words this steel does not exhibit microstructure sensitivity for transformation tem-peratures.However,they are found to be very sensitive to the chemical composition of the steel.The transformation tempera-tures obtained for the RAFM steel are slightly higher than the cor-responding values obtained for the plain9Cr–1Mo steels[22,32]. This is attributed to the addition of tungsten and tantalum,which enhance the stability of the a-ferrite phase thus increasing the A c1 and A c3temperatures.In a similar fashion,it is also expected that the presence of tungsten will exert a signi?cant in?uence on the kinetics of austenite formation during the solution treatment. The velocity of the austenite front during austenitisation requires the diffusion of W and Ta from the austenite into the ferrite,which is sluggish[32].
Similar to the diffusional transformation onset temperature,the martensite start(M s)and?nish(M f)temperatures measured dur-ing cooling for the RAFM steel are found to be714and614K for an initially normalized and tempered condition.These values are slightly higher than that for the9Cr–1Mo grades[33].Additionally, these temperatures were sensitive to the initial microstructure, more importantly the presence of carbides at the austenitising temperature.It is expected that during solution treatment at 1253K the carbide dissolution is not complete due to the sluggish dissolution kinetics of the W containing M23C6precipitates,result-ing in a carbon lean austenite phase.This results in higher M s and M f temperatures as compared to plain9Cr–1Mo steels or RAFM steel austenitised at higher temperature.The presence of some coarse M23C6and?ne MX particles exert a pinning effect on the movement of the martensitic transformation front.Although,a higher austenitising temperature such as1323K would increase the dissolution of M23C6carbides,which in turn would reduce
(a) 1253K
(c)
1453K
(b) 1323K
(d)
1553K
136S.Saroja et al./Journal of Nuclear Materials409(2011)131–139
the M s temperature;however,a temperature of1253K is consid-ered adequate in order to control the prior austenite grain size as observed in Section3.3.Further,systematic studies on the effect of processing variables on the martensitic transformation are un-der progress.
Although the9Cr steel variety is very well studied with respect to microstructural evolution during thermal treatment,there are few calorimetry based studies focused on the energetics and kinet-ics related aspects of phase changes[22,32].In this context,the present study provides an attempt to establish the phase transfor-mation sequence that follows during continuous heating of nor-malized RAFM steel.It is expected that the?ndings of this study, especially the phase transformation temperatures would enable the identi?cation of temperature windows for different processing and heat treatment schedules for the RAFM steel.Further,interpre-tation of the phase changes during a complex process such as welding is simpli?ed due to the uni?ed understanding of the phase changes with temperature cycles.The study would also offer data that are useful in a simulation study of diffusional phase transfor-mation kinetics.
The salient features of the study on transformation behavior of the high temperature austenite in9Cr RAFM steel are as follows: The temperature regimes for the different phase domain have been determined.
The decomposition modes of high temperature as a function of solution treatment temperatures and cooling rate have been established.
4.Conclusions
The physical metallurgy of the 9Cr Oxide Dispersion Strength-ened and Reduced Activation Ferritic/Martensitic steel at different
stages of alloy synthesis and fabrication has been studied.The main results of the study are listed below:
Attritor milling for 3–4h and an extrusion temperature of 1423K are considered optimum.
The dispersoid has been identi?ed as Y–Ti–O complex. The interface of dispersoid with matrix is defect free.
The A c1,A c3,M s and M f temperatures for RAFM steel are mea-sured as 1104,1144,714and 614K respectively using differen-tial scanning calorimetry.
The in?uence of solution treatment temperature and cooling rate on the microstructural evolution during transformation of austenite in RAFM steel has been established.
Acknowledgements
The authors thank Dr.T.Jayakumar Director,Metallurgy and Materials Group for the useful discussions.They also thank the International Advanced Research Centre for Powder Metallurgy and New Materials (ARCI)Hyderabad,India and the Nuclear Fuel Complex,Hyderabad,India for their support during various stages of this research work.References
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Table 2
Transformation temperatures measured by DSC at a heating rate of 5K min à1in a normalized RAFM steel.Description of phase change
On-heating transformation temperature/K a 0(martensite)?a (ferrite)+M 23C 6
(precipitation of M 23C 6on continuous heating)937T C (Curie temperature)
1018a +MX +M 23C 6?a +c +MX +M 23C 6(austenite start (Ae 1)
1104a +c +MX +M 23C 6?c +MX +M 23C 6(austenite ?nish (Ae 3)
1144c +MX +M 23C 6?c +MX (completion of M 23C 6
dissolution)
1550c +MX ?d +c +MX (onset of d -ferrite formation)
1575d +c +MX ?d +c (completion of MX dissolution)
1681d +c ?Liquid +c +d (appearance of liquid)1730Liquid +c +d ?Liquid +d (dissolution of austenite)
1753Liquid +d ?Liquid (melting)
1805
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