Yixun He ,Xudong Liu ,Chenxu Xing,b ,Zhngchi Bin ,Qing Zhou,b ,Jun Wng ,Hifeng Wng,b,Jinshn Li,＊＊
a State Key Laboratory of Solidification Processing,Northwestern Polytechnical University,710072,Xi′an,China
b Center of Advanced Lubrication and Seal Materials,Northwestern Polytechnical University,710072,Xi′an,China
cNational Key Laboratory for Precision Hot Processing of Metals,Harbin Institute of Technology Harbin,150001,China
Keywords:GH 605 superalloy Long-term aging treatment Carbide Strengthening mechanism
ABSTRACT Aging treatment is an effective way to optimize the mechanical properties of Co-based superalloys.In this study,commercial GH 605 superalloy was subjected to aging treatment at 650 °C in a wide time range up to 1000 h.The effects of aging time on the tensile characteristics,microstructure evolution and mechanical properties were systematically investigated at room temperature(RT)and 900 °C.The results showed that the volume fractions of M6C and M23C6 carbide increased with the aging time.After long-term aging treatment,the yield strength(YS)at RT was enhanced from 490.3 MPa to 805.9 MPa,while the alloy still had high tensile ductility (above 20%).Microscopic observations by transmission electron microscopy(TEM)indicated that the strengthening mechanism was related to carbide precipitation inside the grains and the change in the dislocation slipping mode.Moreover,long-term aging treatment can increase the elongation from 24.1% to 47.3% at 900 °C accompanied by a slight increase of YS from 299.3 MPa to 313.9 MPa.Based on detailed microstructure analysis the strengthening mechanism can be attributed to the refined grains as well as carbide precipitation inside the grains and around the grain boundaries.
GH 605 (L-605),nominally Co–20Cr–15W–10Ni,is one kind of typical solid-solution strengthening Cobalt-based superalloy.Because of its exceptional combination of high temperature mechanical properties,corrosion resistance,good workability,and biocompatibility,it has been widely used in aeronautical industry and medical applications [1–4].Afterward,it was found that some of ordered phases segregate in the alloy matrix during aging at a temperature from 600°C to 1000°C,such as the precipitation of Co2W(Laves phase),Co7W6(μ-phase)and various carbides (includingMC,M7C3,M23C6,andM6C) [2,3].MC appears directly in the alloy matrix after solidification and then transforms toM7C3[5].M7C3is a meta-stable carbide that easily degenerates toM6C orM23C6after long-term service[3].Carbides appeared in the alloy matrix sometimes inhibit the movement of dislocation and improve strength[6–10].But the accumulation zone of coarse carbide would appear with the extension of aging time,and it usually acts as crack nucleation sites during thermal deformation at elevated temperature[10,11].Therefore,the variation of different phases in the GH 605 alloy has a great dependence on temperature and aging time,which in turn significantly affects its mechanical properties.
In order to improve the mechanical properties of Co-based superalloys,many researchers have begun to adjust the second phase through a variety of heat treatment methods in recent years.Shun et al.[12]adopted a heat treatment method that combine hot rolling and aging to optimize the grain size of the Co–Cr–Fe–Ni alloy,and successfully improved the tensile strength nearly to 1000 MPa.Ding et al.[13]studied the effect of aging treatment temperature on the surface hardness and wear resistance of a Co-based-VN alloy.They proved that compared with the cases of 550 and 650°C,long-term aging at 750°C can obtain dispersed short rod-shaped dendrites,thereby improving the microhardness and wear resistance.Zhao et al.[14] investigated the effect of aging time at 900°C on the evolution characteristic of γ′precipitates in a Co–Cr–Fe–Ni–Ti alloy.It was noticeable that the morphological transformation of nanoscale γ′precipitates evolved from an initially spheroidal to spheroidal-like shapes,following a cuboidal shape with the increasing aging time.Tang et al.[15]used two aging processes(normal aging and interrupted aging) to control the microstructure and mechanical properties of a Co–Be–Cu–Ni alloy.The results show that disk-shaped γ′′precipitates in the alloy treated by interrupted aging distribute homogenously throughout whole grains with a length of about 3–10 nm,consequently the tensile elongation is significantly improved compared with the results from normal aging.Zangeneh et al.[16]investigated the role of isothermal aging on the microstructure and tribological characteristics of Co–28Cr–5Mo-0.3C alloy.They showed that fine dispersion ofM23C6and lamellar-type carbides obtained by aging at 850°C for 16 h can significantly improve wear resistance as compared with as-cast and heat-treated samples aged for 4 and 24 h.
Motivated by the recent reports,the aging treatment is an effective way to optimize the mechanical properties of Co-based superalloys,but the aging temperature and time should be carefully selected.To date,there are still limited studies on the influence of aging time on the microstructure evolution and mechanical properties of Co-based superalloys,especially long-term aging of hundreds or even thousands of hours.Nevertheless,the thrust-to-weight ratio in advanced aircraft is constantly increasing in recent years,which has led to harsh service environment,including high temperature and long servicing time.Therefore,evaluating the influence of long-term aging time on the microstructure and mechanical properties of Co-based superalloys is of great significance.In this study,commercial GH 605 superalloy was subjected to uniaxial tensile tests at room temperature(RT)and 900°C.The effect of long-term aging treatment on tensile flow characteristics,microstructure evolution and mechanical properties were studied.The results of this research are expected to supply a deeper understanding on the optimization of the service performance and improving the reliability of Co-based superalloys.
Materials used in the present work were produced using vacuum induction melting (VIM) and then refined by the electroslag remelting(ESR) technology.To minimize the casting defects,the casting was forged in multi-directional passes flowed by homogenization treatment at 1150°C for 1 h and water quenching.Aging treatments were performed at 650°C for 100 h,300 h,750 h,and 1000 h in a box electrical resistance furnace,and then air cooling was applied for each one.It should be noted that 650°C is a topical temperature for the precipitation of the strengthening phases.Then,the heat-treated materials were processed into cylindrical tensile samples along the forging direction by the wire electro discharge machining,with the size of 25 mm in gauge length and 5 mm in radius,as shown in Fig.1(a) [17].Tensile tests were performed at RT and 900°C on an Instron 8502 testing machine.The tensile strain rate was 1 × 10-3s-1.The samples were heated to the target temperatures at 5°C s-1and soaked for 10 min to remove any temperature gradient.After the tests,the fractured samples were immediately quenched into water.The fracture and adjacent location of the tensile fracture surface were selected for microstructure characterization,as illustrated in Fig.1(b).
Scanning electron microscope (SEM: FEI Quanta 650 F) equipped with an energy dispersive X-ray spectrometry (EDS) and an electron backscattered diffraction (EBSD) system was used to observe microstructure morphologies,composition,and grain orientations.The samples were ground up to 4000 grit SiC paper followed by mirror polished.For the EBSD observation,the mirror polished samples were subsequently polished by a VibroMet 2 vibratory polisher(Buehler)for 6 h to eliminate the residual stress.The step size and operating voltage of EBSD analysis in this study were 0.4 μm and 20 kV,respectively.All EBSD data were analyzed using Channel 5.0 software (Oxford Instruments).Transmission electron microscopy (TEM) lamellas were prepared by two methods.To analyze the element composition and crystallographic orientation of the specified area,the TEM lamellas were lifted out form the mirror polished surface by using a dual-beam focused ion beam(FIB)workstation (FEI Helios Nanolab 600),and the final thinning was selected as a low Ga+current of 80 pA.The lamellas for the strengthening mechanism research were prepared by hand grinding to a thickness of 60 μm.Then,further thinning was achieved through a twin jet polishing technique with a solution of 10%perchloric acid+90%alcohol at-30°C,20 V condition.TEM observations were conducted under a JEOL JEM-2200FS microscope operating at 200 kV.
3.1.Initial microstructure
Fig.2(a) shows the typical backscattered electron (BSE) micrograph of the received GH 605 alloy.It can be noted that the microstructure consisted of equiaxed recrystallized grains and a few annealing twins.Based on the statistical results by Image Pro software,the average grain size of alloy was about 43 μm (more than 500 grains).As illustrated in Fig.2(b),there were some spherical and island-shaped precipitates sized about 500 nm distributed in intragranular and around the grain boundary.To characterize the precipitate in detail,a TEM lamella was lift-out by FIB from the location indicated in Fig.2(b),and the corresponding bright-field image is shown in Fig.2(d).According to the selected area electron diffraction(SAED)patterns in Fig.2(e)and(f),it is known that both the precipitate and the matrix were of FCC structures.Fig.2(g)and(h) give the elemental composition of the precipitate and the matrix,respectively.For the former,the specific atomic ratio of W/Co/Cr/C was 0.33:0.26:0.2:0.12.Thus,by combining the SAED pattern and elemental composition,these precipitates were identified to beM6C.For the latter,the EDS results (Fig.2(h)) were almost equivalent to nominal composition of GH 605 alloy.
3.2.Microstructure evolution after long-term aging
Fig.1.(a) Shapes and dimensions of the tensile specimen in the present research,(b) specimen morphology before and after tensile test.
Fig.2.Initial microstructure of the received GH 605 alloy:(a)backscattered electron micrograph.(b)Enlarged segment in Fig.2(a)showing the location for preparing the TEM lamella by FIB.(c)FIB cross-section observation of the precipitate in Fig.2(b).(d)TEM image of the precipitate in Fig.2(c).(e)and(f)SAED patterns obtained in the locations indicated by the circles in Fig.2 (d).(g) and (h) Elemental composition of the precipitate and the matrix shown in Fig.2 (d).
Fig.3 shows the BSE micrograph of GH 605 alloy after aging at 650°C for different time.As shown in Fig.3 (a) and (b),the content and distribution ofM6C carbides in the alloy matrix after aging at 650°C for 100 h and 300 h were nearly the same as the initial microstructure(Fig.2(a)and(b)).With the further prolong in aging time to 750 h(Fig.3(c))and 1000 h (Fig.3 (d)),the characteristics of carbides distribution changed significantly compared with the initial microstructure.Fig.3 (c1) and(d1)are the partial enlarged view of Fig.3(c)and(d),respectively.There were a lot of sphericalM6C carbides appearing inside the grains and around grain boundaries as shown by light blue arrows.Moreover,as showing by the yellow arrows,the precipitation ofM23C6carbides at the grain boundaries could also be observed.This means that a typical reaction occurred in the studied GH 605 alloy,that isMC→M23C6,andM23C6→M6C[18].And the precipitation of carbides would need a long time (more than 300 h).
Fig.4.Tensile stress-strain curves of aged GH 605 alloy at (a) RT and (b) 900 °C.
3.3.Tensile properties at different temperatures
Fig.4 depicts the tensile stress-strain curves of GH 605 alloy after aging at 650°C for different time(0 h,100 h,300 h,750 h,1000 h).As shown in Fig.4(a),the tensile flow curves at RT were composed of three distinct stages,i.e.,elastic deformation,work hardening,and necking fracture stages.The yield strength(YS)and the ultimate tensile strength(UTS) of the aged samples increased with the prolonging of aging time.For instance,as shown in Fig.4(a),when the aging treatment time was up to 1000 h,the YS and the corresponding UTS increased from 490.3 MPa to 805.9 MPa and 1010.2 MPa–1130.1 MPa,respectively,compared with the unaged alloy(0 h),with a reduction of the elongation decrease from 41.9%to 20.6%.The specific values are shown in Table 1.
Table 1 Tensile properties of the GH 605 superalloy at RT.
Table 2 Tensile properties of the GH 605 superalloy at 900 °C.
Improving the tensile temperature to 900°C (Fig.4 (b)),it can be noted that the tensile flow curves were also divided into three stages i.e.,elastic deformation,steady stress and flow softening stages.The flow stress firstly increased to a peak value,which is associated with the effect of work hardening.Due to the low stacking fault energy of the studied GH 605 superalloy,the high density of dislocation can accelerate the occurrence of dynamic recrystallization (DRX) once a critical strain is achieved.When the balance between the work hardening and DRX softening is obtained,the tensile flow curves show a transitory steady stress stage,in which the flow stress nearly remains constant with the increasing strain[19].Then,because of the enhancing of DRX softening and the appearance of nucleation-growth process of the voids in the final fracture stage[20,21],the flow stress decreases monotonously till plastic fracture.Interestingly,both the tensile yield strengths and elongations of the samples after long-term aging (750 h and 1000 h) were improved compared with the unaged sample.As shown in Table 2,the YS were close to 299 MPa and 314 MPa when the aging treatment time prolonged from 0 h to 1000 h with an obvious increase in the elongation from 24.1% to 47.2%.This means that the long-term aging treatment is an effective method to improve tensile strength at elevated temperature without sacrificing elongation.Therefore,an in-depth study on the strengthening mechanism was conducted,as discussed in Section 4.2.
3.4.Fracture characteristics
Tensile fracture morphologies of the studied GH 605 alloy tested at RT are shown in Fig.5.Fig.5(a)shows the overview fractography of the unaged sample,and Fig.5 (a1) is the partially enlarged image in Fig.5(a).As illustrated in Fig.5 (a1),the fracture morphology of unaged sample shows a lot of fine dimples and a few tearing ridges,marked by light blue and yellow arrows,respectively.This suggests that a quasicleavage fracture process occurred during the tensile test,which corresponds to a better combination of tensile strength and ductility.
Nevertheless,the fracture morphologies of the aged sample at 650°C for 750 h (Fig.5 (b) and (b1)) had some different characteristics.For example,the number of dimples was greatly reduced in Fig.5(b),and a lot of cleavage surfaces emerged marked by the purple arrows.This means that the transgranular ductile fracture gradually disappeared,and intergranular fracture dominated the tensile deformation.As shown in Fig.5 (c) and (c1),the number of dimples was further reduced and the cleavage surface increased as the aging time was extended to 1000 h.The formation of the intergranular fracture mode is related to the brittle fracture characteristics,which is in agreement with the low tensile elongation and high strength[22].
3.5.Microstructure evolution after tensile test at 900 °C
The deformation microstructures of the studied GH 605 alloy after tensile test at 900°C are shown in Fig.6.As lustrated in Fig.6 (a),the inverse pole figure (IPF) image demonstrates a strong <111>fiber texture and <100>cube texture parallel to the tensile direction (TD).Based on the Taylor model,the <111>fiber texture is one kind of the typical textures in face-centered cubic(FCC)metals,which is developed by polycrystalline homogenous deformation[23].And,the cube texture is the common recrystallization texture during thermal deformation[24].By counting the grain size distribution map in Fig.6 (d),the average grain size of the unaged sample was found to be 8.6 μm,which was significantly reduced compared with the initial microstructure.This is related to the fact that the occurrence of DRX can refine grains during thermal tensile deformation.
Fig.6(b)shows the deformation microstructure of the samples aged at 650°C for 750 h.It can be noted that the strength of <111>fiber texture and<100>cube texture were obviously weakened.With further extension of the aging time to 1000 h,the crystallographic orientation suggests a nearly random distribution characteristic as shown in Fig.6(c).The weakening of texture is related to the occurrence of recrystallization at elevated temperature.Moreover,the average grain size of deformation microstructure after aging for 750 h(Fig.6(e))and 1000 h(Fig.6(f))were reduced to 5.1 μm and 4.7 μm,respectively.This suggests that the long-term aging treatment can not only weaken the texture,but also refine the grain size during tension at elevated temperature.
4.1.Strengthening mechanism at RT
To reveal the effect of long-term aging treatment on the strengthening mechanism of the studied GH 605 alloy,the deformation microstructure such as dislocation characteristics,and the type,size,and distribution of precipitates were observed by TEM.Fig.7(a)shows the TEM images of the unaged sample after tensile tests,and Fig.7(b)is a partial enlarged view of Fig.7 (a).As shown in Fig.7 (b),there were a lot of slip bands appeared in the unaged sample.This suggests that dislocation slipping is the dominant deformation model during tensile tests at RT,which is also observed in other Co-based alloy that subjected to tensile deformation[25].Comparing the orientation of the slip bands with respect to the diffraction vectors,the slip bands were determined to be on the different{1 1 1}planes,and all of the dislocations are slipping in the direction of g=200.
Fig.5.Fracture morphologies of samples tested at RT after aging for (a,a1) 0 h,(b,b1) 750 h,(c,c1) 1000 h.
Fig.7(c)presents the TEM image of the aged sample(650°C/1000 h)after the tensile test.It can be noted that a lot of slip bands existed in the alloy matrix.This indicates that dislocation slipping is still the dominant deformation mechanism in the aged sample.Nevertheless,two typical slipping directions (g=111 and g=200) were observed in the aged sample,which is related to the fact that higher slipping resistance appears compared with the unaged sample during tensile deformation.Fig.7(d)is the detail of a partial enlargement in Fig.7(c).Noted that the dispersive distribution of granular precipitates was observed in the grains as marked by green arrows.The corresponding dark field image in Fig.7(e)confirms that these precipitates belong to the FCC structure,as can be inferred from the SAED patterns in the inset of Fig.7(d).Moreover,the EDS point compositions are illustrated in Fig.7 (f).Based on the SAED analysis and elemental composition,these precipitates are identified asM6C particles,which maintain a coherent relationship with the alloy matrix.
The excellent tensile strength(i.e.,over 1 GPa)of long-term aging GH 605 alloy can be mainly attributed to the introduction of such substantial precipitationM6C particles in the alloy.There would exist abundant lowmisfit coherent interfaces betweenM6C particles and the matrix [26],which resist dislocation motion effectively.As shown in Fig.7 (d),the dispersedM6C particles were surrounded by slipping bands.It can impede dislocation movement by the Orowan by-passing or cutting mechanism.The cutting mechanism would occur when the precipitates are coherent and small,while the Orowan by-pass mechanism can be active when the particles are large or incoherent with the matrix [27].TheM6C particles in Fig.7(d)and(e)were fine and keep coherent with the matrix,therefore,the cutting mechanism is expected to play a dominant role in the present work.In order to coordinate plastic deformation,dislocation slipping occurs accompanied by the cutting through these fine particles.As a result,residual defects generate around the particles.With the increasing strain,the relatively high stresses will be required for dislocation slipping.In this situation,the part of dislocations continues to slip in the original direction,consequently the alloy is strengthened [28].The other part of dislocations would slip along the other close-packed directions.Thus multi-directional slipping directions are activated,which can also be proved in Fig.7(c).
4.2.Strengthening mechanism at 900 °C
Fig.6.Inverse pole figures and grain size distribution maps of the aging sample for (a,d) 0 h,(b,e) 750 h,and (c,f) 1000 h after tensile tests at 900 °C.
Fig.8 shows the SEM-BSE micrographs from near fracture area of studied GH 605 alloy after the tensile test at 900°C.Fig.8 shows thatM6C particles formed inside the alloy matrix and around the grain boundaries after aging treatment.Compared with the unaged sample(Fig.8(a)),increasing aging time can result in a higher number density ofM6C particles (Fig.8(b) and (c)).And the longer aging time,the more dispersed theM6C particles distribution.Therefore,similar to the strengthening mechanism at RT,the dispersedM6C particles can improve the tensile strength of the aged samples by effectively hindering the dislocation motion,including slipping and climbing.In addition,there were a lot of intergranular voids appearing around grain boundaries in the unaged sample as marked by red arrows in Fig.8 (a).Interestingly,the number of intergranular voids was significantly reduced after thermal tensile test when the aging time was extended to 750 h and 1000 h,as illustrated in Fig.8(b)and(c),respectively.The appearance of the voids would accelerate the crack propagation and reduce tensile strength and elongation.Grain boundary sliding(GBS)usually occurs in recrystallized grains and fine grains especially under high temperature and low strain rate.This is because the number of grain boundaries related to sliding is higher and the distance for accommodation by diffusion and/or slip is shorter [29,30].Moreover,GBS can lead to stress concentration around grain boundaries,which is easily to be the initiation zone of intergranular voids [31].Fig.8 (d) is the partial enlarged view of Fig.8 (c).It can be noted that a lot ofM23C6particles (shown by yellow arrows) appeared around the grain boundaries.TheM23C6carbides usually play a main role in preventing the GBS,and provide high stability of grain boundaries[9].Therefore,theM23C6particles precipitated along grain boundaries also contribute to the improvement of tensile strength at elevated temperature.
Recrystallization is a process producing new undistorted and equiaxed grains in the original matrix.The main driving force of recrystallization is the local strain energy[32].Generally,DRX is divided into three categories in FCC alloy: geometric dynamic recrystallization (GDRX),continuous dynamic recrystallization (CDRX),and discontinuous dynamic recrystallization (DDRX).GDRX usually appears at high strain rate,and the size of deformation grains is close to the dimensions of subgrain diameter [33].Since the grain size in the studied alloy was approximately 43 μm(more than 500 grains)(Fig.2(a)),GDRX was hard to happen.Thus,the possible DRX mechanism in this study are CDRX and DDRX.CDRX is characterized by the continuous absorption of dislocations into subgrain boundaries,which would give rise to the appearance of HAGBs and fine grains in the original grains [34].DDRX have obvious grain boundary migration process,and is initiated by the local bulging of deformation grain boundaries [35].Once the grains recrystallize,they would grow up by swallowing the deformed matrix.With the growth of the recrystallized grains,the deformation zone decreases and the recrystallized zone increases,until all deformed matrix transforms to the recrystallized grains.When dislocations pass through theM6C carbides,theM6C particles can exert a pinning force to the dislocation motion in the alloy matrix,consequently the CDRX process is delayed.In addition,the DDRX process can be slowed down because the grain boundaryM23C6particles suppress the growth stage of bulging recrystallized grains.Thus,the carbides,includingM6C andM23C6obtained by the aging treatment can effectively inhibit dynamic recrystallization and refine the grains,which can also be confirmed by grain size distribution map in Fig.6.The fine dynamic recrystallized grains can shorten the slip distance of dislocations and help to release the stress concentration around grain boundaries,as a result,the elongation of aging sample is improved compared with the unaged sample[36].
In addition,the fine grains also play an important role in proving tensile strength at elevated temperature.According to theHall-Petchrelationship,the effect of grain size on the yield strength can be well described as follows[37]:
where σ0is the frictional stress resisting the glide of dislocations,ky isthe Hall-Petch factor,anddis the mean grain size of materials.Moreover,the strength increase caused by refining grain size can be calculated:
Based on the statistical results of the grain size in Fig.6,the strength increment in aged samples for 1000 h is 0.12kycompared with the unaged sample.Roebuck et al.[38] reported that the parameterkyis about 200 MPa μm1/2for the bulk Co-based alloys.Therefore,the strength increment caused by refining grain size was calculated as 24 MPa,which is similar to the test results in Table 2.
Fig.7.TEM images after tensile tests at RT of the studied GH 605 alloy subjected to aging for (a,b)0 h,(c,d)1000 h.(e) TEM dark field image of M6C particles inFig.7 (d).(f) EDS elemental composition pointed by the green arrow in Fig.7 (d).
In the present work,the tensile deformation behaviors of GH 605 alloy have been investigated at different temperatures.The effect of longterm aging treatment on tensile flow characteristics,microstructural evolution and mechanical properties have been studied.
1.With the improving aging treatment time the volume fraction of the carbides(includingM6C andM23C6)increases.The dispersed carbides can effectively inhibit dislocation slipping and change the deformation mechanism from single slipping to multi-directional slipping.After the long-term aging treatment,the YS and UTS increase to 805.9 MPa and 1130.1 MPa,respectively.
2.The Carbides obtained by the aging treatment can suppress the dynamic recrystallization by hinder the dislocation motion and the growth stage of bulging recrystallized grains.As a result,the grain size is refined.The fine dynamic recrystallized grains can shorten the slipping distance of dislocations and help to release the stress concentration around grain boundaries.Consequently the elongation of aged sample after the tensile deformation at 900°C is improved.
3.The carbides precipitating along grain boundaries plays a main role in preventing the grain boundary sliding,and provides high stability of grain boundaries.Therefore,the tensile strength at 900°C of aged sample is improved compared with the unaged sample.
Fig.8.SEM-BSE micrographs from the near fracture area after tensile tests at 900 °C of studied GH 605 alloy subjected to aging for (a) 0 h,(b) 750 h,(c) and (d)1000 h.
Declaration of competing interest
The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.
Acknowledgment
This work was supported by the fund of National Key Laboratory for Precision Hot Processing of Metals (6142909200104),National Science and Technology Major Project(2017-VI-0014-0086 and MJ-2018-G-48),National Training Program of Innovation and Entrepreneurship for Undergraduates (S202010699137) and the Fundamental Research Funds for the Central Universities.We thank Dr.Zheng from ZKKF (Beijing)Science &Technology Company for supporting the characterization of the materials.
