Microstructural Evolution in Directionally Solidified Ni-Base Superalloy IN792+HfS.M.Seo1)†,I.S.Kim1),J.H.Lee2),C.Y.Jo1),H.Miyahara3)and K.Ogi4)1)High Temperature Materials Research Group,KIMS,Changwon,641-010,Korea2)Dept.of Metall.&Mater.Eng.,Changwon Univ.,Changwon,641-773,Korea3)Dept.of Mater.Sci.&Eng.,Kyushu Univ.,Fukuoka,819-0395,Japan4)Oita National College of Technology,Oita,870-0152,Japan[Manuscript received September3,2007]Microstructural evolution during directional solidification(DS)of Ni-base superalloy IN792+Hf has been in-vestigated with an emphasis on theγ precipitates and M C-type carbides.The quantitative image analyses revealed that the increase in the solidification rate up to100µm/s at constant thermal gradient of178K/cm resulted in afine and uniform distribution ofγ precipitates.The relationship between the as-castγ size and cooling rate was also determined for DS IN792+Hf.In the mean time,the M C carbide size was found to be dependent both on the solidification rate and the S/L interface morphology while the area fraction of M C carbide was significantly influenced by the S/L interface morphology.KEY WORDS:IN792;M C corbide;Gamma prime;Solidification rate1.IntroductionNi-base superalloys are extensively used for tur-bine blades and vanes in aero-and industrial gas tur-bine engines.The mechanical properties of these al-loys depend on grain structure,dendrite arm spacing,γ precipitates,γ/γ eutectic and various types of sec-ondary phase such as carbide and boride[1].IN792+Hf is a Ni-base superalloy that contains about12wt pct Cr with a high Ti/Al ratio in its al-loy chemistry.A high Ti/Al compositional ratio in this alloy,however,promotes the formation of low strength eta(Ni3Ti)phase in preference to the de-sirableγ (Ni3Al)phase due to the strong segrega-tion propensity of Ti during solidification.Recently, Seo et al.[2]have shown that microsegregation also re-sulted in the formation of Cr-rich boride phases in the vicinity of eta phase in the as-cast directional solid-ification(DS)IN792+Hf.In addition,the formation of these phases was found to be deactivated with de-creasing the solidification rate because the solid-state diffusion at lower solidification rates decreased the mi-crosegregation.Even though the effect of solidifica-tion rate on the eta and boride formation in IN792+Hf has been well established,limited information is still available regarding the effect of solidification rate on the microstructural features such as M C carbide and γ precipitates that formed during the solidification and the subsequent cooling.In the present study,a series of directional solidi-fication experiments were carried out over a range of solidification rates and the influence of solidification rate on the microstructural evolution,especially onγ precipitates and M C carbides,was investigated.2.ExperimentalThe material used in the present study is the Ni-base superalloy IN792+Hf whose chemical comp-†Senior researcher,to whom correspondence should be ad-dressed,E-mail:castme@kims.re.kr.osition is listed in Table1.Specimens of5.0mm in diameter and80mm in length were directionally solid-ified under Ar atmosphere with various solidification rates(R),R=0.5–100µm/s and constant thermal gra-dient(G)at the S/L interface,178K/cm.The S/L interface was preserved by quenching the specimens after a desired volume fraction of original liquid was solidified.The DS specimens for microstructural observation were prepared by the standard metallographic proce-dures and examined by using an optical microscope and a scanning electron microscope(SEM).Comput-erized image analysis was also performed to quantita-tively analyze the size distribution ofγ precipitates and M C carbides.3.Results and Discussion3.1As-cast microstructureDS experiments were carried out with the solidifi-cation rates of R=0.5–100µm/s under constant ther-mal gradient G=178K/cm.For these DS conditions, the S/L interface morphology of the alloy developed from planar(R=0.5µm/s),to cellular(R=1.0µm/s), and to coarse andfine dendritic(R≥5.0µm/s)with gradually increasing R.Figure1(a)shows the typical as-cast microstructure of DS IN792+Hf solidified at R=50µm/s.The as-cast microstructure was charac-terized by the dendrite core and interdendritic region, which was composed of rosette shapedγ/γ eutectic, M C carbide,eta and Cr-rich boride phases(Fig.1(b)). Thefineγ particles also precipitated in the entireγmatrix during subsequent cooling after solidification.Most of the M C carbides existed near interden-dritic area.Since the major M C forming elements Ti, Ta and Hf exhibited a partitioning tendency to liquid, these elements would be rejected into interdendritic liquid during solidification[2].Therefore,as the solid-ification proceeds,M C carbide forming elements are enriched in interdendritic liquid,which results in the facilitation of nucleation and growth of M C carbide in interdendritic area.Table 1Chemical composition of Ni-base superalloy IN792+Hf (wt pct)Al Co Cr Hf Mo Ta Ti W C B Zr Ni 3.478.712.10.891.84.23.984.30.0720.0160.03Bal.Fig.1Typical as-cast microstructure of DS IN792+Hf solidified at R =50m/s:(a)optical micrograph and (b)de-tailed SEM micrograph near interdendriticregionFig.2Solidification paths of IN792+Hf predicted byThermo-Calc equilibrium and Scheil modelThe eta and boride phases always appeared in front of the coarse γ/γ eutectic as shown in Fig.1(b).Considering that the solidification of γ/γ eutectic proceeds toward the coarse γ [3],eta and boride phases expected to be developed from the residual liquid just after the completion of γ/γ eutectic reac-tion.3.2Solidification pathIn order to examine the solidification sequence of IN792+Hf during solidification,thermodynamic cal-culations were performed using Thermo-Calc soft-ware with Ni-Data developed by Thermo Tech Ltd.(UK).Figure 2presents the solidification paths of IN792+Hf alloy calculated by Thermo-Calc equilib-rium and Scheil model.In the equilibrium model,solidification products were γphase,M C and small amount of M 3B 2.However,the Scheil calculation accounting for the non-equilibrium solidification fea-tures predicted the solidification of γ/γ eutectic and eta phase in addition to primary γ,M C and M 3B 2.The predicted solidification path by Scheil model is asfollows:liquid (L)→primary γ(1608K)→M C car-bide (1596K)→γ/γ eutectic (1467K)→M 3B 2boride (1448K)→eta phase (1261K).Both equilibrium and Scheil model predicted that M C carbide formed at 1596K,which is about 12K lower than the crystallization temperature of primary γphase.This result is comparable to the microstruc-tural observation of Sun et al.[4]who reported that M C carbide forms just below the liquidus tempera-ture of IN792+Hf.In addition,the solidification se-quence of eta and boride phases predicted by Scheil calculation corresponds to the microstructural obser-vation result.However,the Scheil model predicted a very low crystallization temperature of eta phase (1261K)compared with the reported value of about 1402K [2].This discrepancy might be caused by the Hf solubility in eta phase.The Thermo-Calc Scheil model predicted little solubility of Hf in eta phase while the experimental result reported by Seo et al.[2]clearly showed that more than 12wt pct of Hf is dis-solved in eta phase.Therefore,considering that Hf is one of the major elements comprising the eta phase,the Scheil calculation with little Hf solubility in eta phase might delay the eta phase formation to lower temperature.3.3γ precipitatesThe γ precipitates are the primary strengthen-ing phase for Ni-base superalloys.A fine and uni-formly distributed γ size results in desirable mechan-ical properties.The relationship between the mor-phology/size of γ precipitates and the solidification rate is presented in Fig.3,where the microstructure was observed at the similar position of the DS sam-ple,such as the solidification fraction (f s )is about 0.15.The γ precipitates were very large,and showed an irregular and split shape when the solidification rate is very low (Fig.3(a)and (b)).However,the γ particles became obviously fine and their morphology developed from irregular to cuboidal with increasing the solidification rate.From the SEM micrographs shown in Fig.3,the size distribution of γ precipitates and their averageFig.3Effect of solidification rate on the morphology and size ofγ particles in the dendrite core region:(a)R=0.5µm/s,(b)1.0µm/s,(c)5.0µm/s,(d)10µm/s,(e)25µm/s and(f)50µm/sFig.4Size distribution of particles in the dendrite core(a)–(f)and their average size as a function of cooling rate (G·R)Fig.5M C carbide morphology developed from various solidification rates:(a)R=1.0µm/s,(b)5.0µm/s,(c)10µm/s,(d)25µm/s,(e)50µm/s and(f)100µm/sFig.6Effect of the solidification rate (and the S/L in-terface morphology)on the average size and area fraction of M C carbidesize were determined and were summarized in Fig.4.In the case of lower solidification rates,the γ particle size distributed over a wide range (from about 0.5µm to over 1.5µm for R =0.5µm/s)while it showed uni-form size distribution as the solidification rate gradu-ally increased.Figure 4(g)shows the effect of cooling rate (G ·R )on the average γ sizes in comparison with the reported values [5,6].The average γ particle size decreased with increasing the cooling rate in the double logarithmic plot.The linear regression on the basis of the data was derived as follows:d γ =0.33(G ·R )−0.334(1)where d γ is the average size of the γ precipitates.Although the as-cast γ size linear-logarithmically de-creased with increasing the cooling rate in the present study,significant differences still remained compared with the previous studies (Fig.4(g)).This result in-dicates that the as-cast γ size might be an alloy de-pendent,i.e .chemical composition,total amount of γ forming elements and segregation behavior of al-loying elements may influence on the as-cast γ size in Ni-base superalloys.In addition,the volume fraction of γ precipitates appears obviously low at R =0.5,1.0and 5.0µm/s,where the interface morphologies are planar,cellu-lar and coarse dendritic,respectively (Fig.3(a)–(c)).This result is expected to be related with the interface morphology.Macro-segregation occurs in the planar and cellular interface formed at relatively low solidi-fication rates of 0.5and 1.0µm/s.This appears to occur some in the coarse dendritic interface morphol-ogy at 5.0µm/s in the presence of convection which forms inevitably in the Bridgman type directional solidification [9].The γ forming elements (Al,Ti,Ta,Hf)must be lack at the low solidification fraction of DS samples (f s =0.15)due to macro-segregation.Macro-segregation due to the interface morphologies,such as the planar,cellular,and coarse dendritic in-terfaces,is expected to change the volume fraction of γ in the γmatrix.3.4MC carbideM C-type carbide that formed during directional solidification of Ni-base superalloy strengthens longi-tudinal grain boundaries at elevated temperatures.Italso has a significant effect on the solidification behav-ior of Ni-base superalloys [7,8].Figure 5shows the mor-phology evolution of M C carbide during directional solidification of IN792+Hf under various solidification rates.In the lower solidification rates of R =1.0and 5.0µm/s,the morphology of MC carbide exhibited a faceted blocky shape (Fig.5(a)–(b)).As the solidifi-cation rate increased,small script type M C carbides started to form together with large blocky shaped M C carbides (Fig.5(c)),and finally most of the M C car-bide morphologies changed to dendritic script type when the solidification rate is higher than 25µm/s (Fig.5(d)–(f)).Figure 6shows the variation of average M C car-bide size and area fraction as a function of solidifica-tion rate.The average M C carbide size was found to be dependent on the S/L interface morphology as well as the solidification rate.As the S/L inter-face morphology changes from cellular to dendritic (R =1.0µm/s to 5.0µm/s),the average M C carbide size slightly increased.However,in dendritic solid-ification conditions (R ≥5.0µm/s),the M C carbide size rapidly decreased with increasing the solidifica-tion rate at slower rates of R =5.0–25µm/s,and this tendency became sluggish at relatively high solidifica-tion rate range (R ≥25µm/s).The steep decrease in M C carbide size in the solidification range of R =5.0–25µm/s,appears to be due to the evolution of script type M C carbide and the coarse inter-dendritic spac-ing may provide the larger growth of M C carbide.The area fraction of M C carbide appears to be rather dependent on the S/L interface morphology than the solidification rate.The area fraction of M C carbide increased when the S/L interface changes from cellular to dendritic morphology.However,the increase in the solidification rate did not have a significant effect on the fraction of M C carbide,about 0.76%,where the S/L interface was dendritic morphology.The area fraction of M C carbide at 1.0µm/s,showing the cellular interface,is expected to be low due to the macro-segregation of carbide form-ing elements (Ti,Ta and Hf)at the low solidification fraction of f s =0.15.The lower area fraction of γ ,which contains Ti,Ta and Al shown in Fig.3(a),may also reveal the reducing of these elements due to the micro-and macro-segregation.4.Conclusions(1)The solidification path of IN792+Hf alloy pre-dicted by Thermo-Calc Scheil model is as follows:L →primary γ→M C →γ/γ eutectic →M 3B 2→eta phase.(2)The increase in the solidification rate up to 100µm/s at a constant thermal gradient of 178K/cm resulted in a fine and uniform distribution of γ pre-cipitates within supersaturated γmatrix.The as-cast γ size appeared to be alloy dependent and a following relationship between the γ size and cooling rate was established for DS IN792+Hf:d γ =0.33(G ·R )−0.334(3)The M C carbide size was found to be depen-dent on the S/L interface morphology as well as thesolidification rate while the area faction of M C car-bide was strongly related to the S/L interface mor-phology during directional solidification.The average M C size decreased with increasing the solidification rate,but its area fraction was nearly constant where the S/L interface exhibits dendritic morphology. AcknowledgementThis work was supported by the National Research Laboratory Project of Korean Ministry of Science and Technology.The authors also acknowledge thefinancial support of Japan Society for the Promotion of Science (JSPS)through the RONPAKU fellowship.REFERENCES[1]C.T.Sims,N.S.Stoloffand W.C.Hagel:Superalloys II,Wiley,New York,NY,1986,97.[2]S.M.Seo,I.S.Kim,J.H.Lee,C.Y.Jo,H.Miyahara andK.Ogi:Metall.Mater.Trans.A,2007,38A,883. [3]Y.Zhu,S.Zhang,L.Xu,J.Bi,Z.Q.Hu and C.X.Shi:Su-peralloys1988,eds.D.N.Duhl,TMS,Warrendale,PA, 1988,703.[4]W.R.Sun,J.H.Lee,S.M.Seo,S.J.Choe and Z.Q.Hu:Mater.Sci.Eng.A,1999,271,143.[5]K.O.Yu,J.J.Nichols and M.Robinson:J.Metals,1992,21.[6]X.Guo,H.Fu and J.Sun:Metall.Mater.Trans.A,1997,28A,997.[7]S.Tin and T.M.Pollock:Meter.Sci.Eng.A,2003,A348,111.[8]S.Tin,T.M.Pollock and W.Murphy:Metall.Mater.Trans.A,2001,32A,1743.[9]J.H.Lee,S.Liu and R.Trivedi:Metall.Mater.Trans.A,2005,36A,3111.。
