Materials and Design 39 (2012) 55–62

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Materials and Design

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The microstructure evolution and its effect on the mechanical propertiesof a hot-corrosion resistant Ni-based superalloy during long-term thermalexposure

Jian Wang a, Lanzhang Zhou a,⇑, Liyuan Sheng b,c, Jianting Guo a

a Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, Chinab College of Engineering, Peking University, Beijing 100871, Chinac PKU-HKUST, ShenZhen-HongKong Institution, Shenzhen 518057, China

a r t i c l e i n f o

Article history:Received 23 November 2011Accepted 13 February 2012Available online 21 February 2012

Keywords:A. Ferrous metals and alloysC. Heat treatmentsF. Microstructure

0261-3069/$ - see front matter � 2012 Elsevier Ltd. Adoi:10.1016/j.matdes.2012.02.020

⇑ Corresponding author. Tel.: +86 24 83971911; faxE-mail address: lzz@imr.ac.cn (L. Zhou).

a b s t r a c t

The microstructure evolution and its influence on the mechanical properties are investigated in a hot-corrosion resistant Ni-based superalloy during long-term thermal exposure. It is found that the tertiaryc0 phase disappears and the secondary c0 phase coarsens and coalesces gradually, which acts as the mainreason for the decreasing of strength at both room temperature and 900 �C. During exposure, the grainboundary coarsens from discontinuous to half-continuous and finally to continuous structure. The opti-mum half-continuous grain boundary structure composed of discrete M23C6 and M3B2 wrapped by c0 filmleads to the elongation peak at room temperature in the thermally exposed specimens. At 900 �C, theincrease in the elongation is attributed to the much softer matrix and the formation of microvoids.The behavior of primary MC decomposition is a diffusion-controlled process. During exposure, variousderivative phases including M23C6, c0, g, M6C and r sequentially form in the decomposed region. PrimaryMC decomposition and the precipitation of r phase have little effect on the mechanical properties due totheir low volume fractions.

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1. Introduction

This Ni-based superalloy is a candidate material as blades ofmodern turbine engines due to its good high temperature strengthand excellent hot-corrosion resistance. Its good mechanical proper-ties at high temperature is derived from the high content of Al andTi elements forming the ordered c0 phase with the volume fractionof about 45% and amounts of solid solution strengthening elements,such as Cr, W and Mo. Its excellent hot-corrosion resistance is en-sured by high concentration of Cr and Co elements and high Ti/Alratio. Indeed, the introduction of these elements confers the alloygood mechanical properties and hot-corrosion resistance. But, italso brings about complex and unstable microstructure, whichwould inevitably degenerate gradually during long term service.Extensive studies [1–5] have shown that prolonged thermal andstress exposure causes overaging of the microstructure, consistingof c0-phase coarsening and coalescence, TCP phase formation, pri-mary MC carbide decomposition and formation of continuous sec-ondary M23C6 carbide chains on the grain boundaries. It is wellknown that the coarsening and coalescence of c0 phase is the most

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principal factor that deteriorates the mechanical properties [2].Also, the precipitation of large amounts of TCP phase plays animportant role in degrading the strength because of its depletingof the solid solution elements [5,6]. In addition, the TCP phase ofneedle morphology could initiate and accelerate the crack propaga-tion [7,8]. Depending on the composition of the primary MC and theconstituent elements present in the alloy, subsequent solid-statetransformation may decompose the MC into a variety of carbidesand g phase [9,10]. Furthermore, the decomposition of primaryMC would release carbon atoms into the matrix, which acts as thecarbon source of the carbides coarsening in the grain boundary[10]. But in general, primary MC decomposition has little effecton the mechanical properties due to its low volume fraction. Therole of grain boundary in influencing the mechanical property isdual depending on its structure. The optimum grain boundarystructure composed of thin c0 film decorated with discrete carbideparticles inhibits sliding and damage accumulation while the con-tinuous grain boundary structure consisting of carbide chainwrapped by coarser c0 film facilitates the propagation of crack dur-ing service [4,11].

Obviously, the microstructure evolution during long-termservice is really critical to the mechanical properties. So it is mean-ingful and necessary to investigate the thermal stability of this

56 J. Wang et al. / Materials and Design 39 (2012) 55–62

hot-corrosion resistant Ni-based superalloy. In this article, the ef-fect of long-term thermal exposure on the microstructure evolu-tion has been examined and correlated with changes in tensileproperties.

2. Experimental procedure

The chemical composition of the investigated alloy is given asfollows (wt.%): 0.08 C, 15.6 Cr, 10.6 Co, 5.5 W, 2.0 Mo, 3.1 Al, 4.3Ti, 0.2 Nb, 0.2 Hf, 0.05 B and balance Ni. The alloy was remeltedin an industrial scale vacuum induction furnace and then cast intobars of 15 mm in diameter and 220 mm in length. Then the as-castbars were subjected to standard heat treatment: 1160 �C/3 h/AC + 1060 �C/4 h/AC + 850 �C/16 h/AC. After standard heat treat-ment, the bars were divided into three groups. The three groupswere exposed at temperatures of 800 �C, 850 �C and 900 �C fortimes of 1000, 3000, 5000 and up to 10,000 h, respectively. Andthe subsequent microstructures were examined on the JEOL 6340

Fig. 1. Microstructure observation of the standard heat treated specimen: (a) SEM microgbimodal distribution of c0 , (c) SEM micrograph of the grain boundary structure, (d) BF micleft insets showing the SAED patterns of M3B2 and M23C6, respectively), (e) BF micrograpMC.

Field Emission Gun Scanning Electron Microscope (FEGSEM)equipped with an energy dispersive spectroscopy (EDS) microanal-ysis and the TECNAI G2 F30 transmission electron microscope(TEM). The micrographs were recorded under bright-field (BF),dark-field (DF), and selected area electron diffraction (SAED).Z-contrast imaging was obtained in the Tecnai G2 F30 transmissionelectron microscope equipped with a high-angle annular dark field(HAADF) detector.

Metallographic samples were grounded, mechanically polished,and then etched by two different methods. Chemical etching meth-od in a reagent containing 20 g CuSO4, 50 ml HCl and 100 ml H2Owas used to reveal the c0, carbides, grain boundary structure andminor phases. Deep etching method in an electrolyte containing5 ml HNO3, 10 ml acetic acid and 85 ml H2O, which stripped awayc matrix, was employed for three-dimension observation of c0

phase. Each datum of the average size and volume fraction of c0

phase was evaluated using ImageJ software by at least 10 micro-graphs. To prepare foils for transmission electron microscopy

raph of c/c0 matrix, c–c0 eutectic, primary MC and M3B2, (b) SEM micrograph of therograph of M23C6, M3B2 particles in the grain boundary (the upper right and bottomh of primary MC (the inset showing its SAED pattern), (f) EDS spectrum of primary

Fig. 2. Microstructure observation of the specimen heat treated at 850 �C for 1000 h: (a) SEM micrograph of c0 phase, (b) SEM micrograph of the discontinuous grainboundary structure, (c) SEM micrograph showing fine M23C6 particles along the MC/matrix interface, (d) BF micrograph of the c0–M23C6 film along the MC interface (the insetshowing the orientation relationship of (100) c0//(100) M23C6 and [100] c0//[100] M23C6).

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(TEM) observation, thin slices were sectioned and grounded. Thefoils were dimpled using gatan model 656 Dimple Grinder, and thenion milled to perforation using precision ion polishing system 691.

Tensile tests at room-temperature and 900 �C were carried outon a Shimadzu AG-250KNE test machine. The cylindrical, threadedtensile test rods with gauge length of U 5 mm � 25 mm machinedfrom the thermally exposed bars were prepared as Standard GB/T4338-2006 [12]. Each datum is an average of at least two tensiletested values.

3. Results and discussion

3.1. Microstructure evolution

Microstructure observation of the specimen after standard heattreatment is presented in Fig. 1. It can be seen that the microstruc-ture consists of c/c0 matrix, c–c0 eutectic (primary c0), primary MCcarbides and M3B2 borides, as shown in Fig. 1a. The c0 phase in thedendrite core displays bimodal distribution: small oval c0 particles(tertiary c0) with small volume fraction distribute in the matrixchannel among larger cuboidal c0 particles (secondary c0) withaverage edge length of 300 nm, as shown in Fig. 1b. After standardheat treatment, the alloy contains much fine grain boundary,which is decorated with c0 precipitates, M23C6 particles and smallquantities of M3B2 particles, as displayed in Fig. 1c. Further trans-mission electron microscopy (TEM) observation reveals that thefine M23C6 and M3B2 particles homogeneously distribute alongthe grain boundary, as shown in Fig. 1d. The upper right and bot-tom left insets show the selected area electron diffraction (SAED)patterns of the M3B2 with the lattice parameters of a � 5.78 Å,c � 3.13 Å and the M23C6 with the lattice parameter ofa � 10.52 Å, respectively. In Fig. 1e, it can be seen that bright-field

(BF) micrograph and the SAED pattern (in the inset) confirm theblocky primary MC as TiC type carbide with the lattice parameterof approximately 4.32 Å. Moreover, the energy dispersive spectros-copy (EDS) microanalysis shows the primary MC is rich in Ti and Welements, as shown in Fig. 1f.

During long-term thermal exposure, the microstructure of theexperimental alloy degenerates gradually, which depends on theaging temperature and time. The microstructure observation ofthe specimen heat treated at 850 �C for 1000 h is shown in Fig. 2.It can be seen that the tertiary c0 particles disappear while the sec-ondary c0 particles coarsen to the size of 410 nm but degenerate tothe cube morphology with rounded corner, as shown in Fig. 2a. It isclear that the larger secondary c0 particles grow at the expense ofthe smaller ones, which is consistent with the Ostwald ripeningprocess. According to LSW theory [13,14], the ripening process isdriven by the reduction in total interfacial energy. After standardheat treatment, the bimodal distribution of c0 possesses high inter-facial energy, which would decrease gradually as thermal exposureproceeds by merging smaller oval c0 particles and degenerating lar-ger edged cuboidal c0 particles into the rounded cube ones. InFig. 2b, it can be seen that the grain boundary is still fine: the c0

film is thin, and the discrete M23C6 and M3B2 particles distributehomogeneously along the grain boundary. However, the primaryMC has already decomposed to some extent, which results in sometiny M23C6 particles along the MC/matrix interface, as shown inFig. 2c. Further observation by TEM reveals that a c0 film decoratedwith M23C6 particles (c0-M23C6 film) forms and surrounds the pri-mary MC, as shown in Fig. 2d. Typical [001] pattern exhibits thatthe M23C6 has a perfect cube-on-cube orientation relationship withthe c0 phase, which can be described as: (100) M23C6//(100) c0 and[100] M23C6//[100] c0, as demonstrated in the inset. The behaviorof primary MC decomposition is a diffusion-controlled process [2].

Fig. 3. Microstructure observation of the specimen heat treated at 850 �C for 3000 h: (a) SEM micrograph of c0 phase, (b) SEM micrograph of the half-continuous grainboundary structure, (c) SEM micrograph of the darker g phase and brighter M6C particles in the c0–M23C6 film, (d) HAADF micrograph of the g and M6C surrounded by the c0–M23C6 film, (e) and (f) the SAED patterns of g and M6C, respectively.

58 J. Wang et al. / Materials and Design 39 (2012) 55–62

As described above, primary MC is mainly enriched with C, Ti, andW elements. During thermal exposure, the primary MC acts as thesource of C and Ti elements, while the matrix serves the source ofNi, Al and Cr. The rapidly diffused C atoms arrest the Cr atoms fromthe matrix and form M23C6 particles along the MC/matrix interface.In the meantime, amounts of Ti element (c0 forming element) areseparated from the MC and diffuse into the matrix, which leadsto the formation of c0 film around the primary MC.

As thermal exposure proceeds to 3000 h at 850 �C, the micro-structure degenerates more severely, as shown in Fig. 3. It can beseen that the secondary c0 particles with rounded corner coarsencontinuously to the size of 450 nm and a few of them coalesce witheach other, as shown in Fig. 3a. Note that the coarsening and coa-lescence of c0 particles are still driven by the reduction in the totalinterfacial energy. At this thermal exposure condition, the grainboundary coarsens noticeably. The c0 film becomes thicker and

the M23C6 and M3B2 particles increase in size and quantity andcoalesce into a half-continuous chain, as shown in Fig. 3b. It mustbe noted that the discrete, closed spaced M23C6 and M3B2 particlesengulfed in c0 film improve the rupture life and ductility by hinder-ing grain boundary sliding, which is widely accepted by previousinvestigators [4,15]. In addition, the primary MC also decomposesseverely. As shown in Fig. 3c, the c0 film becomes thicker and theM23C6 particle grows larger. Further examination reveals that somebrighter M6C particles and darker g phase form in the c0–M23C6

film. TEM observation of the decomposed MC is displayed inFig. 3d–f. The high-angle annular dark field (HAADF) micrographclearly shows that the g phase and M6C surrounded by the c0–M23C6 film form around the decomposed MC, as shown inFig. 3d. The SAED patterns confirm the M6C with the lattice param-eter of 11.16 Å and g phase with the lattice parameters ofa � 5.07 Å and c � 8.22 Å, respectively, as shown in Fig. 3e and f.

Fig. 4. Microstructure observation of the specimen heat treated at 850 �C for 5000 h: (a) SEM micrograph of c0 phase, (b) SEM micrograph of the continuous grain boundarystructure, (c) SEM micrograph showing the darker g phase, brighter M6C particles, gray M23C6 particles and needle r phase in the decomposed region, (d) HAADF micrographof the decomposed MC showing the g phase with M6C and M23C6 inside and c0–M23C6 shell, (e) and (f) BF micrograph and the SAED pattern of the needle r phase,respectively.

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With the thickening of the c0 film during thermal exposure, theexchange of elements between the primary MC and the matrixbecomes difficult. Because the thicker c0 film not only blocks thediffusion the Al element through the c0 film from the matrix, buthinders the Ti element diffusing into the matrix as well [16]. Sothe high Ti/Al ratio is achieved on the MC/c0 interface, whichpromotes the nucleation of the g phase. The formation of g phaseleads to the segregation of W and Mo elements due to their limitedsolubility in the g phase [17]. The segregated W and Mo atomstrap the C atoms released from the remaining MC and form theM6C phase.

Microstructure observation of the specimen heat treated at850 �C for 5000 h is shown in Fig. 4. As it can be seen, the second-ary c0 particles coarsen and coalesce continuously, and its meandiameter increases to 500 nm, as shown in Fig. 4a. Beyond this,the microstructure degeneration displays some distinct features.

First, the coarsening of the grain boundary becomes exacerbated.The c0 film thickens and widens aggressively into the matrix; thediscrete M23C6 and M3B2 particles grow much larger and coalesceinto a continuous chain, as shown in Fig. 4b. A continuous grainboundary structure composed of thicker c0 film decorated withcoarser M23C6 and M3B2 chain forms at last. The formation of thewide and continuous grain boundary structure is detrimental tothe mechanical property since the continuous M23C6 and M3B2

chain wrapped in the c0 film facilitates the propagation of crackand leads to the tensile brittleness [11,18]. Moreover, the primaryMC decomposition becomes more severe. g Phase expands in-wards extensively and some brighter M6C and gray M23C6 particlesprecipitate in its interior, as shown in Fig. 4c. Further observationby TEM of the decomposed MC clearly shows that the g phase dec-orated with the brighter M6C and gray M23C6 particles is the innerfilm while the c0–M23C6 film is the outer film in the decomposed

Fig. 5. The yield strength and elongation of the standard heat treated and thermallyexposed K444 alloy (a) at room temperature and (b) 900 �C.

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region, as displayed in Fig. 4d. The formation of M6C and M23C6

particles in the interior of g phase should be ascribed to the exten-sive formation of g phase which leads to the segregation of W, Moand Cr elements. The segregated W, Mo and Cr atoms absorb the Catoms released from the remaining MC and form the M6C andM23C6 phases in the interior of g phase.

Another distinct feature is the formation of r phase with lim-ited volume fraction around the decomposed MC, as shown inFig. 4c. TEM observation of the r phase is displayed in Fig. 4eand f. It can be seen that the r phase presents needle morphologyand its lattice parameters are determined to be as: a � 8.81 Å andc � 4.55 Å. The preferred precipitation of r phase around decom-posed MC is different from previous reports [19] that the r phaseis prone to precipitate in the dendrite core and interdendritic re-gion. This might be related to the precipitation of M6C and M23C6

phases in the interior of g phase. Their precipitation consumesthe released C atoms from the MC, which would result in the for-mation of carbon-depleted zone around the decomposed MC.According to previous research of Weiss and Stickler [20], the highcarbon content in solid solution can prevent or at least retardappreciably the nucleation of intermetallic phases. Therefore, itcan be understood that the depletion of carbon around the decom-posed MC promotes the formation of r phase.

With the increasing of time to 10,000 h at 850 �C, further micro-structure degeneration occurs. The c0 particles coarsen to the aver-age size of 570 nm and the grain boundary keeps continuous state.Moreover, the long aging time almost completely decomposes theprimary MC to the g phase with some M6C and M23C6 particles in-side. r Phase with limited volume fraction still forms around

decomposed MC. It is necessary to mention that no r phase is ob-served when thermal exposure at 800 or 900 �C.

During long term thermal exposure, the microstructure of the cmatrix also degenerates to some extent. The c matrix channelwidens gradually by the coarsening and coalescence of c0 phase,which is beneficial to the ductility. And moreover, the c matrixbecomes depleted of solid solution elements because of the coars-ening of M23C6 and M3B2 particles in the grain boundary. In addi-tion, the formation of various derivative phases of primary MCand the precipitation of r phase also consume the solid solutionelements. All these microstructure evolutions would decrease thestrength of the alloy [6,10].

3.2. Mechanical properties

The tensile test results of standard heat treated and thermallyexposed specimens at room temperature and 900 �C are illustratedin Fig. 5. In Fig. 5a, it can be seen that the yield strength at roomtemperature decreases as thermal exposure proceeds. The elonga-tion also decreases generally, but it shows a peak at aging time of3000 h. At 900 �C, the yield strength decreases while the elonga-tion increases with the increasing of aging temperature and time,as shown in Fig. 5b.

Fig. 6 shows the micrographs of fracture surface and the longi-tudinal section near the rupture site after tensile tests at roomtemperature. The fractographs of the specimens exposed at850 �C for 3000 h and 5000 h both present intergranular fracturemode. But it displays considerable amounts of dimples in thematrix of the specimens exposed at 850 �C for 3000 h, as shownin Fig. 6a. While in the specimen exposed for 5000 h it shows thedecreased amount of dimples and some faceted-cleavages in thematrix, as shown in Fig. 6c. Further observation on the longitudinalsection shows that the microcracks form in the grain boundaries ofboth specimens exposed at 850 �C for 3000 h and 5000 h, as shownin Fig. 6b and d. However, the microcracks are much larger in sizeand quantity in the specimens exposed for 5000 h.

The micrographs of fracture surface and longitudinal sectionnear the rupture site after tensile tests at 900 �C are illustrated inFig. 7. It can be seen that the fractographs display the intergranularfracture mode and have great amounts of dimples in the matrix ofboth standard heat treated specimens and the specimens exposedat 850 �C for 5000 h, as shown in Fig. 7a and c. And further exam-ination on the longitudinal section shows that some microcracksand small quantities of microvoids form in the grain boundariesof the standard heat treated specimens, as shown in Fig. 7b. Whilein the exposed specimens, it is observed to have a mass of microv-oids in the grain boundaries and interdendritic regions, as shownin Fig. 7d.

For cast nickel based superalloys, the most important strength-ening mechanism is the c0 precipitate strengthening. Both thepresence of hyperfine c0 and the average size and volume fractionof c0 precipitate are the significant factors that affect the strength-ening effect [21]. During long term exposure, the disappearing oftertiary c0 and the coarsening of secondary c0 are mainly responsi-ble for the decreasing of the yield strength at both room tempera-ture and 900 �C.

There are a number of factors affecting the ductility of cast poly-crystalline superalloys, including deformation mechanism, embrit-tlement of grain boundaries by carbide particles, and instability ofc0 particles during exposure [22]. The highest elongation of stan-dard heat treated specimen at room temperature might be attrib-uted to the deformation mechanism. At room temperature, thedislocation is difficult to cut through the c0 phase [23], thus form-ing pile-ups at c/c0 interface and leading to the strain localizationin the thermally exposed specimen. However, the homogeneousdistributed hyperfine c0 precipitates in the matrix channel block

Fig. 6. SEM micrographs of fracture surface (a) and (c) and longitudinal section (b) and (d) near the rupture site in the tensile specimen tested at room temperature: (a) and(b) 850 �C/3000 h, (c) and (d) 850 �C/5000 h.

Fig. 7. SEM micrographs of fracture surfaces (a) and (c) and longitudinal section (b) and (d) near the rupture site in the tensile specimen tested at 900 �C: (a) and (b) SHT, (c)and (d) 850 �C/5000 h.

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the movement of dislocation, which relieves the pile-ups at theinterface between larger cuboidal c0 and the matrix. Therefore,the dislocation distributes more homogeneous and little strain

localization forms in the matrix. As a result, the elongation exhibitsthe highest in the standard heat treated specimen at roomtemperature.

62 J. Wang et al. / Materials and Design 39 (2012) 55–62

The occurrence of elongation peak at aging time of 3000 h is thecombined effect of the widening of the c channel and the grainboundary structure. The widening of the c channel due to thecoarsening of c0 facilitates the movement of mobile dislocations,which contributes to the ductility. The role of grain boundary onthe ductility is twofold. As described in Fig. 6b and d, the half-continuous grain boundary structure consisting of discrete M23C6

and M3B2 particles in the grain boundary hinders the formationof microcracks in the grain boundary, while the continuous grainboundary structure composed of M23C6 and M3B2 chain wrappedby c0 film has the contrary effect. Therefore, the ductility woulddecrease gradually as grain boundary coarsens from the half-continuous to continuous structure during thermal exposure. Sothe widening of the c channel and the structure of grain boundaryco-affect the ductility of the tensile tested specimen at roomtemperature, which leads to a peak of elongation of the thermallyexposed alloy.

At 900 �C, the matrix becomes much softer. And strain localiza-tion is difficult to form in the tensile tested specimens. Moreover,the formation of masses of microvoids in the thermally exposedspecimens would relieve the stress concentration, which is benefi-cial to the ductility. So the elongation increases at 900 �C as ther-mal exposure proceeds.

4. Conclusion

The microstructure evolution and its effect on mechanical prop-erties of a hot-corrosion resistant Ni-based superalloy have beenstudied during long-term exposure at elevated temperature andcan be summarized as follows:

(1) During long-term thermal exposure, the tertiary c0 phasedisappears and the secondary c0 phase coarsens and coa-lesces gradually. The disappearing of tertiary c0 phase andthe coarsening and coalescence of secondary c0 phase arethe principal reason for the reduction in the yield strengthat both room temperature and 900 �C.

(2) The grain boundary coarsens from discontinuous to half-con-tinuous and finally to continuous structure during exposure.The optimum half-continuous grain boundary structurecomposed of discrete M23C6 and M3B2 wrapped by c0 filmcorresponds to the elongation peak of the thermally exposedspecimens at room temperature. The much softer matrix andthe formation of microvoids both result in the rise of elonga-tion at 900 �C.

(3) The behavior of primary MC decomposition is a diffusion-controlled process. During long-term thermal exposure, var-ious derivative products including M23C6, c0, g, M6C and rsequentially form in the decomposed region. Primary MCdecomposition and the precipitation of r phase has littleeffect on the mechanical property due to their low volumefractions.

Acknowledgements

The authors are grateful to the Program of ‘‘863’’ (Grant No.2011AA030104) and the Program of ‘‘973’’ (Grant No. 2009CB930004) for the financial support.

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