Evolution of microstructure and changes of mechanical properties

Evolution of microstructure and changes of mechanical propertiesof CLAM steel after long-term agingXue Hu a,Lixin Huang a,b,Wei Yan a,Wei Wang a,Wei Sha c,Yiyin Shan a,n,Ke Yang aa Institute of Metal Research,Chinese Academy of Sciences,72Wenhua Road,Shenyang110016,Chinab State Key Laboratory of Metastable Materials Science and Technology,Yanshan University,Qinhuangdao066004,Chinac School of Planning,Architecture&Civil Engineering,Queen's University Belfast,Belfast BT95AG,UKa r t i c l e i n f oArticle history:Received21June2013Received in revised form1August2013Accepted9August2013Available online19August2013Keywords:SteelAgingMechanical characterizationElectron microscopyLight microscopyMicroanalysisa b s t r a c tThe China Low Activation Martensitic(CLAM)steel has been developed as a candidate structural materialfor future fusion reactors.It is essential to investigate the evolution of microstructure and changes ofmechanical properties of CLAM steel during thermal exposure.In this study,the long-term thermal agingof the CLAM steel has been carried out in air at6001C and6501C for1100h,3000h and5000h.The microstructural evolution with aging time was studied,including characteristics of the growth ofM23C6carbides and the formation of Laves-phase precipitates as well as the evolved subgrains.Themicrostructural evolution leads to the changes of mechanical properties of the CLAM steel.The Ductile–Brittle Transition Temperature(DBTT)increases significantly during the thermal aging,which is relatedto the formation of Laves-phase in the steel matrix.The possible mechanism of stabilizing microstructureduring the thermal exposure has been analyzed based on the interaction between M23C6carbides andsubgrain boundaries.&2013Elsevier B.V.All rights reserved.1.IntroductionReduced Activation Ferritic/Martensitic(RAFM)steels have mod-ified compositions of conventional ferritic–martensitic8–12%CrMoVNb steels,mainly by replacing Mo,Nb and Ni with W and Ta in orderto obtain low activation capability[1].The RAFM steels,such as F82H[2,3],Eurofer97[1,3,4]and CLAM steel[5,6],are considered as candi-date structural materials for future nuclear fusion reactors.The CLAMsteel(9Cr1.5WVTa)has been developed in China and shows com-paratively excellent mechanical properties[5].All breeding blankets have been conceptually designed to servicein the temperature windows of250–5501C for RAFM steels.Moreadvanced versions such as the Oxide Dispersion Strengthened(ODS)steels will be for above550–6501C[1,7].In addition to the reducedactivation performance,the outstanding high-temperature creepstrength is also required for these structural steels.High micro-structural stability is crucial to obtaining excellent creep strength.Inthe CLAM steel,the Cr-rich M23C6carbides with relatively large sizeand the MX carbonitrides rich in Ta or V with comparatively smallsize are employed to stabilize the microstructure so as to achievehigher creep strength at high temperature.However,during long-term creep exposure at elevated temperatures above5001C,themicrostructure will inevitably degrade due to the combined effect ofthermal activation and loading stress,which consequently influencesthe creep strength.It has been reported that the microstructuraldegradation generally consists of three processes[8]:(1)the recov-ery of martensitic laths,(2)the agglomeration of carbides,and(3)the growth of subgrains,causing deterioration of creep strengthat6001C for a9Cr–2W steel.The CLAM steel cannot be an exception.Therefore,for the safety in the future service,it is important toinvestigate the microstructural evolution and the correspondingchanges of mechanical properties of these steels during long-termexposure at high pared with the loading stress,thethermal activation plays a much more effective role on the micro-structural evolution.Thus,the present work was focused on theeffect of long-term thermal aging on evolution of both microstruc-ture and mechanical properties of the CLAM steel.2.Experimental procedureThe chemical composition of the CLAM steel chosen for thisstudy was Fe–0.093C–0.49Mn–8.96Cr–0.16V–1.51W–0.14Ta–0.05Si–0.0048P–0.0024S–0.0073N–0.004O(in wt%).The hot-forgedCLAM steel plates were heat treated with normalizing(9801C/30min/air cooling)and tempering(7601C/90min/air cooling).Finally,the tempered steel was subjected to exposure in air at6001C and6501C for1100h,3000h and5000h.Contents lists available at ScienceDirectjournal homepage:/locate/mseaMaterials Science&Engineering A0921-5093/$-see front matter&2013Elsevier B.V.All rights reserved./10.1016/j.msea.2013.08.025n Corresponding author.Tel./fax:þ862423971517.E-mail address:yyshan@(Y.Shan).Materials Science&Engineering A586(2013)253–258Both the cylindrical tensile specimens(Φ5M10)and the Charpy-V-notch specimens(10Â10Â55mm3with a451and2mm deep notch)were machined parallel to the longitudinal direction of plates. All the tensile tests were performed in air at room temperature and 6001C.Two tensile test specimens were used in each condition for the heat-treated steel and the steel after aging for1100h at6001C. The Charpy-V-notch impact tests were conducted at a temperature range fromÀ1001C to room temperature.Two samples were used in each condition for impact tests.The surfaces of all the samples were carefully polished and etched by Villella's reagent(1g of picric acid,5ml of hydrochloric acid and100ml of ethyl alcohol).The etched surfaces were observed by using an optical microscope(LEICA MEF4A)and a scanning electron microscope(HITACHI S-3400N).TheΦ3mm discs used for the transmission electron microscope(FEI-TECNAI20)observa-tion were thinned with a solution of10%perchloric acid and90% acetic acid by an electron polisher(Struers TenuPol-5),at the temperature of11–151C and the voltage of28–30V.3.Results3.1.Microstructure3.1.1.Heat treatment,martensitic lath and subgrain structure,M23C6Micrographs of the heat-treated CLAM steel are presented in Fig.1,showing a tempered martensitic microstructure.The grain size was10–30μm(Fig.1a),and the width of martensitic laths ranged0.2–0.5μm.M23C6carbides were found to distribute along the martensitic lath boundaries(Fig.1b).The tempered martensitic microstructure of the CLAM steel was relatively stable during aging at6001C,as shown in pared with the heat-treated sample(Fig.1b),the martensitic lath in Fig.2a did not coarsen obviously and the width of laths was0.2–0.4μm.Till aging for5000h,the martensitic lath boundary still could be observed clearly,and the width of martensitic lath was increased to about 0.55μm(Fig.2c).Only a few martensitic laths were decomposed into subgrains.However,after increasing the aging temperature from 6001C to6501C,the stability of the microstructure was reduced.After 1100h,the martensitic lath boundaries could be observed clearly,and only a few subgrains appeared in the steel matrix(Fig.3a).After thermal exposure at6501C for3000h,although shorter than5000h at6001C,martensitic lath boundaries disappeared and a large number of subgrains with the size of0.85–1μm were formed in the steel matrix(Fig.3b).The M23C6type(Fe,Cr)23C6carbides play an important role on the microstructure stability for the heat-resistant steel at high temperature.There were two different morphologies of M23C6 carbides,including the rod-shaped M23C6carbides and spherical M23C6carbides(Fig.3c).The M23C6carbides(60–200nm)usually precipitated at the prior austenite grain boundaries,martensitic lath boundaries(Fig.4)and subgrain boundaries with high density dislocations around them.The electron diffraction pattern of the M23C6carbides is shown in Fig.4b.In the TEM(FEI-TECNAI20),there are four different diffraction aperture sizes in total,including800m m,200m m,40m m and 10m m.The10m m diffraction aperture size was used for the selected-area diffraction work.Corresponding to this aperture size, the actual selected area on the specimen is about170nm,because of the magnification of the objective lens between the specimen and the aperture.So,the matrix contribution in the electron dif-fraction patterns of large M23C6precipitates,approaching200nm, can be avoided.Of course,if we want to acquire electron diffrac-tion patterns of the M23C6with a smaller size,e.g.60nm,the matrix contribution in the electron diffraction pattern will be harder to avoid,because the size of diffraction aperture is larger than that of precipitate.The small size MX type(Ta,V)(C,N)carbonitrides(10–40nm) pinning dislocations mainly dispersed within martensitic laths or subgrains(Fig.4).Subgrain boundaries were actually aligned walls of dislocations.Only the small size M23C6carbides can provide the pinning effect on subgrain boundaries.As the size of M23C6 carbides was similar to the thickness of subgrain boundaries,it was hard for subgrain boundaries to cross these M23C6carbides (Fig.3b).It could be found that both the number and the size of carbides were increased during the thermal aging(Figs.2and3).ves-phaseThe Back Scattered Electron(BSE)images were used to distin-guish the(Fe,Cr)2W Laves-phase from M23C6type(Fe,Cr)23C6 precipitates,arising from the difference in the mean atomic weight for different particles.Furthermore,the atom weight of Ta in the MX type(Ta,V)(C,N)carbonitrides was also high.However,the size of MX carbonitrides(10–40nm)was much smaller than that of Laves-phase(4200nm),so the MX type carbonitrides couldnotFig.1.Micrographs of the CLAM steel after heat treatment.(a)Optical and(b)TEM.X.Hu et al./Materials Science&Engineering A586(2013)253–258254be observed in the low magni fication BSE images.Since the (Fe,Cr)2W Laves-phase was observed as a bright particle in these images (Fig.5),it was possible to examine the change in volume fraction and size distribution of Laves-phase during long-term thermal aging.The Laves-phase usually precipitated on the martensitic lath boundaries and the prior austenite grain boundaries.During the thermal aging at 6001C,the number of Laves-phase was increased signi ficantly (Fig.5a –c).After increasing the aging temperature from 6001C to 6501C,the number of Laves-phase particles was decreased signi ficantly.No Laves-phase appeared in the CLAM steel matrix after aging for 1100h at 6501C (Fig.5d).Laves-phase appeared as block-shaped particles that engulfed M 23C 6carbides along grain and lath/subgrain bound-aries (Fig.6a).The (Fe,Cr)2W Laves-phase had a fine fringe contrast.The chemical composition of (Fe,Cr)2W Laves-phase analyzed by the EDAX was 50Fe –12Cr –4Ta –1Mn –32W (in at%),as shown in Fig.6b.3.2.Strength and impact propertiesStrength of the CLAM steel was changed during the long-term thermal exposure.Two tensile test specimens were used in each condition for the heat-treated steel and the steel after aging forspherical rodFig.3.TEM micrographs of the CLAM steel after aging at 6501C.(a)1100h;(b)3000hand (c)5000h.Fig.4.TEM micrograph (a)of precipitates and subgrain boundary and electron diffraction pattern (b)of M 23C 6carbides in the CLAM steel after aging at 6501C for 1100h.Fig.2.TEM micrographs of the CLAM steel after aging at 6001C.(a)1100h;(b)3000h and (c)5000h.X.Hu et al./Materials Science &Engineering A 586(2013)253–2582551100h at 6001C.The standard deviations of tensile strength (s b )and yield strength (s s )data at room temperature and 6001C were very small.The small-size error bars even cannot be seen clearly,as shown in Fig.7.After aging at 6501C,all the samples exhibited softening (Fig.7b).However,the samples exhibited hardening slightly after aging for more than 3000h at both 6001C and 6501C.Changes of Upper Shelf Energy (USE)and DBTT of the CLAM steel after thermal aging are shown in Fig.8.The standard deviation of impact absorbed energy data in some conditions was relatively large.The USE of the steel changed somewhat irregularly.The toughness of the thermally-aged CLAM steel decreased during the thermal exposure,especially after aging for 5000h.The DBTT of the thermally-aged CLAM steel was increased to about À101C and À201C,respectively,after aging for 5000h at 6001C and 6501C.4.Discussion4.1.Martensitic lath recovery and its effect on strengthThe tempered martensite is not a thermodynamically equili-brium phase.Hence,it evolves gradually during the thermal aging at high temperature.The most signi ficant and obvious microstruc-tural evolution is the recovery of martensitic lath,which has thestrongest effect on strength.The martensitic lath recovery is a process of dislocation movement and dislocation annihilation,resulting in the migration of martensitic lath boundaries and the formation of subgrains [8].The formation of subgrains was progressively developed by increasing temperature and the plastic strains [9].Both plastic deformation and elevated temperature can accelerate the martensitic lath recovery process due to their promoting dislocations mobility [9,10].It was found in the present study that the martensitic laths in the CLAM steel evolved to subgrains after aging at 6501C for 3000h (Fig.3b),but they were still stable after aging at 6001C for 5000h,which indicated that the elevated temperature played an effective role in the formation of subgrain.During the thermal exposure at 6501C,the dislocation mobility was high owing to the thermal activation.Lots of dislocations were annihilated,resulting in a great decrease in dislocations density.A low dislocations density would lead to an obvious decrease in strength of the CLAM steel after long-term aging,as shown in Fig.7.However,compared with the CLAM steel thermally aged at 6501C,the mobility of dislocations in the CLAM steel was relatively weaker at 6001C,as indicated by the slightly decreased dislocation density shown in Fig.2.Even after much longer time aging at 6001C,subgrains could not be formed in the steel matrix.Therefore,the strength of the CLAM steel after aging at 6001C did not reduce as much as that after aging at 6501C.The precipitation of Laves-phase after aging for 3000h at both 6001C and 6501C,asFig.5.Low magni fication BSE images of the CLAM steel after aging.(a)6001C/1100h;(b)6001C/3000h;(c)6001C/5000h;(d)6501C/1100h;(e)6501C/3000h and (f)6501C/5000h.Fig.6.TEM micrograph of precipitates (a)and EDAX analysis of Laves-phase and (b)in the CLAM steel after aging at 6001C for 5000h.X.Hu et al./Materials Science &Engineering A 586(2013)253–258256shown in Fig.5,could provide the precipitation strengthening,compensating for the loss of strength.This could lead to a slight increase in strength after aging over 3000h,as shown in ves-phase precipitation behavior and its effects on toughness During thermal exposure at high-temperature,the precipita-tion behavior of Laves-phase is crucial to mechanical properties and thermal stability of the heat-resistant steel.The Temperature –Time –Precipitation (TTP)curve of Laves-phase exhibits the C shape.Both the precipitation and dissolution of Laves-phase occur simultaneously in the CLAM steel at high temperature.Different volume fractions of Laves-phase dissolve in the matrix at different temperatures.For the 9Cr –W steel,the nose point of the Fe 2W phase TTP curve is located at 6501C and in 20–30ks [11].At the nose temperature,the Laves-phase precipitates in the shortest time.It can be observed in the present study that after aging at 6501C,no Laves-phase appeared in the CLAM steel after aging for 1100h (Fig.5d)and the quantity of Laves-phase was small in the CLAM steel after aging for 3000h or longer time (Fig.5e and f).All these phenomena indicate that the nose temperature of Laves-phase for CLAM steel should be lower than 6501C.Two reasons can be mainly used to explain these phenomena.Firstly,compared with the 9Cr –W steel [11],the W content is 1%higher in the CLAMsteel.The higher W content forces the Laves-phase to precipitate at lower temperature.Secondly,since 6501C is quite close to the dissolution temperature of Fe 2W,Laves-phase may be more prone to dissolving than precipitating at 6501C.The small size Laves-phase could provide precipitation hardening to a certain degree,but it grows so fast to reduce the mechanical properties of heat-resistant steel,especially the creep properties [10,12]and the impact toughness [13,14].The Laves-phase with an average size beyond 0.13μm can trigger the fracture mode transition from ductile (transgranular fracture)to brittle (intergranular fracture)[12].Cavities seem to nucleate at boundaries next to large particles,such as Laves phases [10,12].The large size Laves-phase can act as the cavity trigger.These cavities induce the initiation of microscopic cracks during the plastic deformation.Moreover,the formation of Laves-phase is at the expense of the dissolved W and Mo atoms near grain boundaries,which reduces the solid solution strengthening.Consequently,the appearance of Laves-phase should be detrimental to the mechanical properties of CLAM steel.After aging for 5000h at both 6001C and 6501C of the CLAM steel,a large number of Laves-phase precipitates appeared in the steel matrix (Fig.5c and f),and the size of Laves-phase precipitates was larger than in other samples.So it was understandable that the DBTT of samples aged for 5000h at both 6001C and 6501C was increased to around À201C as shown in Fig.8,the most signi ficant increase.In addition,the formation of Laves-phase swallows many M 23C 6carbides on the prior austenite grain boundaries.M 23C 6200300400500600700S t r e n g t h / M P aAging time/ h200300400500600700S t r e n g t h / M P aAging time/ hFig.7.Strength of the CLAM steel after long-term thermal aging.(a)Aging at6001C and (b)aging at 6501C.050100150200250I m p a c t E n e r g y / JTemperatrure/ oCTemperatrure/ o C050100150200250I m p a c t E n e r g y / JFig.8.Impact energy of the CLAM steel after long-term thermal aging.(a)Aging at 6001Cand (b)aging at 6501C.X.Hu et al./Materials Science &Engineering A 586(2013)253–258257carbides can mainly provide the nucleation site for Laves-phase on the grain boundaries.Without enough M23C6carbides pinning grain boundaries,as will be discussed below,the stability of microstructure would decrease obviously.4.3.Precipitates and subgrain boundariesFor9–12%Cr ferritic/martensitic heat-resistant steel,the sub-grain boundaries are the main obstacles against the motion of dislocations.The migration of subgrain boundaries,causing the coarsening of subgrains,is closely correlated to an acceleration of the creep[11,15].Hence,the subgrain boundary hardening is an important thermal stability mechanism.However,only the proper size particles can prevent the motion of subgrain boundaries so as to achieve thermal stability for the microstructure.During a long-term thermal aging,there were mainly three kinds of particles formed in the CLAM steel,M23C6carbides,MX carbonitrides and Laves-phase particles.The MX carbonitrides (10–40nm)were generally formed inside laths(Fig.2a and3a) and subgrains(Fig.3b).However,the M23C6carbides(60–200nm) and Laves-phase(mainly4200nm)mainly precipitated along martensitic lath boundaries and prior austenite grain boundaries (Fig.2c).The different types and sizes of particles affected the thermal stability of microstructure differently.The M23C6carbides played a more important role than the MX carbonitrides on controlling the migration of subgrain boundaries due to the correlation between the spacing of M23C6carbides and subgrain size[16].The MX carbonitrides mainly pinned the single disloca-tion in subgrains,but could not prevent the motion of subgrain boundaries and lath boundaries.If there are only MX carbonitrides on martensitic lath boundaries,the lath boundaries still bulge out and enter into the adjacent laths[17].However,the recovery of laths could be stopped by M23C6carbides,as shown in Fig.2c. The migration of subgrain boundaries was inhibited by the pinning force provided by M23C6carbides,as shown in Fig.3b and c. Therefore,it can be concluded that the thermal stability of subgrains depends on the M23C6carbides[18].In order to stabilize the subgrain boundaries it is important to prevent the coarsening of M23C6carbides.5.ConclusionsIn this study,the microstructural evolution and the mechanical property changes of the CLAM steel after long-term thermal aging were investigated.The following conclusions can be reached. (1)After increasing the aging temperature from6001C to6501C,the tempered martensite microstructure of CLAM steel evolved into subgrains after a long-term exposure,which indicates that the thermal activation plays a very important role in the microstructural evolution.(2)During long-term thermal aging at both6001C and6501C,theprecipitates played an important role on pinning the lath boundaries and subgrain boundaries in the CLAM steel.The M23C6carbides(60–200nm)could stabilize the subgrain boundaries,but thefine MX carbonitrides(10–40nm)could have little contribution.(3)During thermal aging at6501C,Laves-phase showed slowerprecipitation than at6001C,indicating that the nose tempera-ture of Laves-phase formation in the CLAM steel possibly is lower than6501C.At the beginning of precipitation,Lave-phase could make a small contribution to the strength.Mean-while,the precipitation of Laves-phase was detrimental to the impact toughness,and promoted the DBTT even up toÀ201C after aging for5000h at both6501C and6001C. AcknowledgmentsThis work wasfinancially supported by funds from the ITER Program(No.2009GB109002)and Chinese Academy of Sciences Knowledge Innovation Program(KJCX2-YW-N35).This work was in cooperation with the FDS Team of Institute of Plasma Physics, Chinese Academy of Sciences.References[1]R.Lindau,A.Möslang,M.Rieth,M.Klimiankou,E.Materna-Morris,A.Alamo,A.A.F.Tavassoli,C.Cayron,ncha,P.Fernandez,N.Baluc,R.Schäublin,E.Diegele,G.Filacchioni,J.W.Rensman,B v.d.Schaaf,E.Lucon,W.Dietz,Fusion Eng.Des.75–79(2005)989–996.[2]P.Marmy,B.M.Oliver,J.Nucl.Mater.318(2003)132–142.[3]K.Ehrlich,Fusion Eng.Des.56–57(2001)71–82.[4]E.Lucon,W.Vandermeulen,J.Nucl.Mater.386–388(2009)254–256.[5]Y.Li,Q.Huang,Y.Wu,T.Nagasaka,T.Muroga,J.Nucl.Mater.367–370(2007)117–121.[6]Z.X.Xia,C.Zhang,Q.Y.Huang,S.J.Liu,Z.D.Li,Z.G.Yang,J.Iron Steel Res.Int.18(2011)47–52.[7]R.Lindau,A.Möslang,M.Schirra,P.Schlossmacher,M.Klimenkov,J.Nucl.Mater.307–311(2002)769–772.[8]F.Abe,S.Nakazawa,H.Araki,T.Noda,Metall.Trans.A23(1992)469–477.[9]H.Cerjak,P.Hofer,B.Schaffernak,ISIJ Int.39(1999)874–888.[10]C.G.Panait,W.Bendick,A.Fuchsmann,A.F.Gourgues-Lorenzon,J.Besson,Int.J.Pres.Ves.Pip.87(2010)326–335.[11]F.Abe,H.Araki,T.Noda,Metall.Trans.A22(1991)2225–2235.[12]J.S.Lee,H.Ghassemi-Armaki,K.Maruyama,T.Muraki,H.Asahi,Mater.Sci.Eng.A428(2006)270–275.[13]L.Schäfer,J.Nucl.Mater.283–287(2000)707–710.[14]P.Fernández,M.Hernández-Mayoral,peña,ncha,G.De Diego,Mater.Sci.Technol.18(2002)1353–1362.[15]F.Abe,Metall.Mater.Trans.A34(2003)913–925.[16]H.Ghassemi-Armaki,R.P.Chen,K.Maruyama,M.Igarashi,Metall.Mater.Trans.A42(2011)3084–3094.[17]K.Sawada,M.Taneike,K.Kimura, F.Abe,Mater.Sci.Technol.19(2003)739–742.[18]H.Ghassemi-Armaki,R.P.Chen,K.Maruyama,M.Yoshizawa,M.Igarashi,Mater.Lett.63(2009)2423–2425.X.Hu et al./Materials Science&Engineering A586(2013)253–258 258。