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Effects of Long-Term Thermal Exposure on the Microstructure and Properties


Effects of Long-Term Thermal Exposure on the Microstructure and Properties of a Cast Ni-Base Superalloy
X.Z. QIN, J.T. GUO, C. YUAN, C.L. CHEN, and H.Q. YE Microstructural stability and microstructure-property relationship during long-term thermal exposure in K452 alloy (a new Ni-base cast superalloy with ~21 pct Cr, 11 pct Co, 3.5 pct W, 2.5 pct Al, 3.5 pct Ti, and others) are investigated. It is found that exposure temperature and time have signi?cant e?ects on the microstructure and properties of the alloy. During exposure, the microstructure is degraded by c? coarsening, MC carbide (M mainly represents Ti, W, and Nb) degeneration, precipitation and evolution of grain interior (GI) M23C6 carbide, evolution of grain boundary (GB) microstructure, and precipitation of g phase. Among them, the c? coarsening is the leading reason for the decrease of strength of the alloy. The GI M23C6 and the g phase have negligible in?uence on the properties due to their relatively small populations. Blocky, closely spaced GB M23C6 particles engulfed in c? increase the stress-rupture life, whereas the formation of a continuous GB M23C6 chain has an opposite e?ect. A life peak occurs when the M23C6/c? structure at the GBs is in an optimal form. The degenerated MC is the preferred initiation site of microcracks. Its presence at the GBs promotes the onset of intergranular fracture, and leads to the decrease in mechanical properties. DOI: 10.1007/s11661-007-9381-5 ? The Minerals, Metals & Materials Society and ASM International 2007

I.

INTRODUCTION

THE K452 alloy is a newly developed high Cr content cast Ni-base superalloy for advanced gas turbine vane applications in marine and industrial ?elds. It performs well under laboratory conditions with good fatigue resistance, hot-corrosion resistance, and tensileand stress-rupture properties, in addition to being completely oxidation resistant up to 900 °C.[1] Compared to aircraft engines, gas turbines for industrial applications, which are exposed to corrosive environment for prolonged periods of time, have more stringent requirements for the hot-end superalloy components to accommodate for their formidable operating conditions. Accordingly, for Ni-base superalloys such as K452, microstructural stability is a very important consideration, which must be carefully examined and assured before putting them to use. Superalloys generally experience various microstructural changes during their service life, including c? coarsening, formation of a continuous GB carbide network, topologically close-packed (TCP) phase formation, and MC carbide degeneration.[2–6] These processes remove much of strengthening elements from the c matrix and signi?cantly degrade the properties of the alloys, such as mechanical properties, corrosion resisX.Z. QIN, Student, J.T. GUO, Professor, and C. YUAN, Associate Professor, Superalloys Division, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, P.R. China and C.L. CHEN, Student, and H.Q. YE, Academician, are with Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, Shenyang 110016, P.R. China. Contact e-mail: xzqin@imr.ac.cn Manuscript submitted December 18, 2006. Article published online November 7, 2007
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tance and service life. In this article, the microstructural evolution during long-term thermal exposure in the K452 alloy is examined in detail, and the microstructure-property relationships are explored carefully.

II.

EXPERIMENTAL PROCEDURE

K452 has the composition (wt pct) of 0.105C, 20.9Cr, 11.15Co, 3.5W, 0.6Mo, 0.25Nb, 2.5Al, 3.5Ti, 0.04Zr, 0.015B, and rest Ni. Specimens were subjected to a homogenization for 4 hours at 1170 °C followed by furnace cooling to 900 °C and then air cooling to room temperature; subsequently, two annealing treatments (4 hours at 1050 °C and 16 hours at 850 °C), producing the secondary and tertiary c? precipitates, were carried out. After the heat treatment, specimens were exposed at temperatures of 800 °C, 850 °C, and 900 °C for times of 1000, 3000, 5000, and up to 10,000 hours, respectively. For each exposure condition at least four specimens were prepared for tensile- or stress-rupture tests (two for tensile-rupture tests and two for stress-rupture tests). Tensile- and stress-rupture tests (gage length 50 mm, gage diameter 5 mm) were performed at 900 °C and 900 °C/201 MPa, respectively. Each rupture value, including strength and elongation, represents an average of at least two test results. The microstructures were examined using optical microscope (OM) and scanning electron microscope (SEM). Chemical etching was employed for the general microstructural observation using a solution containing 20 g CuSO4, 50 mL HCl, and 100 mL H2O, which dissolves c?. Deep etching with an electrolyte of 200 g KCl + 50 g citric acid + 200 mL HCl + 1000 mL H2O, which removes the c matrix, was conducted for
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the three-dimensional observation of c?. The specimens worked as the anode while the cathode was a stainless steel plate; the current density was ~0.5 A/cm2. Transmission electron microscopy (TEM) equipped with an energy dispersive X-ray spectroscope (EDS) was used for phase identi?cation. Foils for TEM were prepared on a twin-jet electropolisher with a solution of 10 pct perchloric acid and 90 pct ethanol at -25 °C. The room-temperature lattice parameters of the c? and c phases were determined by X-ray di?raction methods. The c? phase was extracted using the electrolyte of 10 g (NH4)2SO4 + 15 g tartaric acid + 1200 mL H2O. Slices (12 mm in diameter, 4 mm in thickness) of heattreated and thermally exposed specimens were chemically etched for about half a hour in order to remove the c? phase on the surface of the slices; then, the lattice parameter of the c phase was approximately estimated by di?racting the etched slices. The (311)c and (331)c? re?ection peaks were used for the calculation of lattice parameters of the c and c? phases, respectively. Silicon was used as a standard for calibration. Lattice mis?t d is de?ned by d ? 2 (ac0 ? ac )=(ac0 ? ac ) ? 1?

thermal exposure, these constitutional phases undergo signi?cant microstructural changes, resulting in severe modi?cation of mechanical properties. A. Microstructural Evolutions 1. c? Coarsening As demonstrated in Table I, the lattice parameter (ac) of the c matrix declines during long-term thermal exposure though scatter exists. Especially a signi?cant reduction occurs within the exposure time of 1000 hours. In contrast, the parameter (ac?) of the c? precipitates basically remains unchanged in the whole period of exposure. As a result, the magnitude of the c-c? mis?t, according to Eq. [1], ?rst presents a sharp drop within 1000 hours exposure and then a mild decrease when exposure time is longer than 1000 hours. The precipitation of secondary phases (e.g., M23C6-c? cell and g, as depicted later) from the c matrix may be responsible for the reduction of ac. That is, as the secondary phases precipitate, alloying elements escape from the c lattice and result in its shrinkage. In Ni-base superalloys, the morphology of c? particles depends both upon the magnitude of the c-c? mis?t and the size of the particles.[7–10] Under the heat-treatment condition, the investigated alloy has a large c-c? mis?t (-0.35 pct), which makes the secondary c? precipitates having a considerable average size (305 nm) tend to be cuboidal in shape for minimizing the elastic energy (Figure 1). However, during long-term thermal exposure, the secondary c? precipitates always assume a spherical appearance, because of the extremely small mis?ts (Table I), in order to minimize the c/c? interfacial energy (Figure 2). For Ni-base alloys, precipitate alignment is understood in terms of the minimization of the elastic interaction energy, which is a function of not only the

where ac? and ac are the unconstrained lattice parameters of the c? and c phases, respectively. The size and volume fraction of precipitates were measured using image analysis software. Each measured value was an average of 50 readings obtained from 50 di?erent images.

III.

RESULTS AND DISCUSSION

The microstructure of heat-treated K452 alloy consists of c matrix, c? precipitate, c-c? eutectic, primary MC carbide, and GB M23C6 carbide. During long-term
Table I.

Lattice Parameters of the c and c? Phases and the c-c? Mis?ts under the Conditions of Heat Treatment and 900 °C Thermal Exposure Heat Treatment 900 °C/1000 h 3.5902 3.5891 0.03 900 °C/5000 h 3.5901 3.5923 -0.06 900 °C/10,000 h 3.5901 3.5876 0.07

?) ac? (A ?) ac (A d (pct)

3.5904 3.6029 -0.35

Fig. 1—SEM images of c? precipitates in the heat-treated specimen: (a) chemically etched and (b) electrolytically etched. The secondary and tertiary c? precipitates are clearly visible in the insert.
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Fig. 2—Secondary c? precipitates in the specimens exposed at 900 °C for (a) 1000 h, (b) 5000 h, and (c) 10,000 h. Deep etching technique is used for the three-dimensional observation of the c? precipitates.

lattice mis?t between the precipitates and the matrix but also the size and volume fraction of the precipitates.[10–12] In this alloy, the alignment does not occur under the heat-treatment condition, because of the relatively minor volume fraction (13 pct) of the secondary c? precipitates (Figure 3(b)). Similarly, neither the alignment is present during long-term thermal exposure, due to the sharp drop of magnitude of the c-c? mis?t (Table I). In these two cases, the elastic interaction energy is low, which may be the essential reason why the secondary c? precipitates do not align in the present alloy, unlike many other superalloys.[8,13,14] The size and volume fraction changes of the secondary c? precipitates at the exposure temperatures of 800 °C, 850 °C, and 900 °C are shown as functions of exposure time in Figures 3(a) and (b), respectively. It is clear that the coarsening rate of the secondary c? precipitates increases with exposure temperature, and that their volume fraction goes up with exposure temperature or time. Also, one can understand that both the dissolution of tertiary c? precipitates and the redistribution of alloying elements under nonequilibrium conditions are in fact responsible for the increase in the size and volume percentage of the secondary c? precipitates. 2. Primary MC Degeneration In heat-treated specimens, most primary MC carbides (M represents Ti, W, and Nb; Table II) with an average size of 6.92 lm have an irregular blocky morphology and are located both at the GBs and in the interdendritic regions. During long-term thermal exposure, these carbides deteriorates severely and various transformation products are sequentially present on their peripheries with exposure, as illustrated in Figure 4.[15] Initially, a thin layer of c? decorated with small and discrete Cr-rich M23C6 particles emerges at the MC/c
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interface, suggesting that a classical decomposition reaction, that is MC + c ? M23C6 + c?, is operative there (Figure 4(a)). As the reaction proceeds, the MC acts as the source of C and Ti, while the c matrix serves as the source of Ni, Al, and Cr. With the accumulation of c? and M23C6, the exchange of alloying elements between the primary MC and the c matrix becomes increasingly di?cult. Nevertheless, because Ni di?uses through the obstacle of the c?M23C6 ?lm more easily than Al and Ti,[16,17] it combines with Ti to form the g (or Ni3Ti) phase at the interface between the ?lm and the MC (Figure 4(b)). The g phase depletes the c? forming elements (principally Ni and Ti), consequently suppressing c?. Therefore, the form of primary MC degeneration can now be written as MC + c ? M23C6 + g at the intermediate stage of thermal exposure, which was once reported in Reference 6. Finally, during the late stage of exposure, the di?usion of alloying elements, especially the high atomic number W and Mo atoms, is further blocked by the ordered g phase. As a result, W and Mo are enriched locally within the reaction region, facilitating the formation of the a-(W, Mo) phase that is sporadically embedded in the slight-gray g phase (Figure 4(c)). The presence of the a-(W, Mo) phase establishes the operation of another MC decomposition reaction that can be summarized into MC + c ? M23C6 + a-(W, Mo) + g. It should be pointed out that M23C6 is always one of the transformation products in the whole process of primary MC degeneration, for C di?uses much faster than any other element and can relatively easily combine with Cr at any stage of thermal exposure mentioned previously. The EDS measurements of various phases involved in the process of primary MC degeneration are shown in Table II. The various phase interfaces, such as M23C6/c?, g/MC, and g/c?, are formed in the process of primary MC
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of degenerated, coarse, and irregular MC carbides distributed at the GBs are especially more detrimental than when they are in the GIs. 3. Precipitation and Evolution of M23C6 Carbides within GIs M23C6 carbides dendritically precipitate from the supersaturated c matrix and coarsen during thermal exposure. Their average size is both temperature and time dependent: For a given exposure temperature, the size increases and gradually approaches a critical value with time; and, the higher the exposure temperature, the larger the critical size.[21] The precipitation of M23C6 carbides rich in Cr, W, and Mo leads to the local enrichment of c? forming elements, such as Ni, Al, and Ti and then the presence of the c? precipitates, forming an M23C6-c? eutectoid cell.[21] In the cell, initially the M23C6 carbide is dendrite-shaped and the c? precipitates are embedded in the carbide dendrite. Under di?erent thermal exposure conditions, the M23C6 carbide undergoes various degrees of morphological changes: At 800 °C, it basically exhibits a dendritical appearance through the entire period of exposure; while at 850 °C, its morphology tends to transform from ?owerlike dendrite to irregular block. However, the most notable morphological transition occurs at 900 °C, as illustrated in Figure 5, where the M23C6 carbide evolves in shape from ?owerlike dendrite to irregular block to regular polyhedron; meanwhile, the c? precipitates embedded in the M23C6 dendrites migrate outward and agglomerate gradually into a spherical c? envelope encircling the central carbide. It should be pointed out that the minimization of the M23C6/c? interphase boundary energy is actually the driving force of the morphological evolution of M23C6 carbides, as described in Reference 21. Regardless of size and shape, M23C6 can act as a strong barrier to mobile dislocations and stacking faults that have cut through the c? envelope around the carbide.[21] However, the following aspects determine the inappreciable in?uence of M23C6 on mechanical properties: (1) the population of M23C6 is small, less than 1.3 pct in volume fraction; (2) the precipitation and coarsening of M23C6-c? cells remove much of solid solution strengthening elements, that is, Cr, W, Mo, Al, and Ti, from the c matrix and weaken the alloy;[5,22] and (3) M23C6 particles mainly distribute in the interdendritic regions and often crowd in the form of loops, chains, or clusters, as demonstrated in Figure 6, so that their resistance to dislocation motion is discounted severely.

Fig. 3—Increases in the (a) diameter and (b) volume fraction of the secondary c? precipitates with exposure time at various exposure temperatures.

deterioration. Under the applied load, these interfaces, because of high stress concentration and interfacial energy, tend to become the favored initiation sites and propagation paths of microcracks.[15] Because microcracks promote the fracture of superalloys and contribute to the reduction of service life,[5,18–20] MC degeneration, which results in more frequent cracking than the intact MC, is undoubtedly deleterious to the thermally exposed K452 alloy. A considerable amount
Table II.

EDS Results of Various Phases Involved during the Primary MC Degeneration (Weight Percent) Co Cr 26.60 1.27 4.64 65.08 2.44 1.72 Al 1.41 — 10.47 0.75 1.53 0.24 Ti 2.23 46.75 10.42 1.97 13.35 3.21 W 3.64 26.37 1.22 10.70 4.20 72.97 Nb — 16.55 — — 3.74 1.01 Mo — 1.85 — 2.23 — 4.25 Ni 52.12 6.57 69.03 17.04 68.91 14.84

c MC c? M23C6 g-(Ni3Ti) a-(W, Mo)

14.00 0.64 4.22 2.23 5.83 1.76

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Fig. 4—Backscattered electron images of degenerated MC carbides under the di?erent exposure conditions: (a) through (c) thermal exposure of 800 °C/1000 h, 850 °C/5000 h, and 900 °C/10000 h, respectively.[15]

Fig. 5—Morphologies of M23C6 carbides in the specimens exposed at 900 °C for (a) 1000 h, (b) 5000 h, and (c) 10,000 h.

4. Evolution of GB Microstructure The GB phases observed in this investigation are c? precipitate, M23C6 carbide, and infrequently primary MC carbide. In the heat-treated specimen, the GBs are ?ne, as there is only a small quantity of M23C6 particles and c? precipitates present (Figure 7(a)). However, during long-term thermal exposure the GBs coarsen through the following possible processes (Figure 7(b) through 7(e)): (1) the small discrete M23C6 particles at the GBs grow and gradually interlink into a thin discontinuous, and then a coarse continuous chain; and (2) a continuous c? ?lm forms on either side of the
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chain and thickens with the rise of exposure temperature or time. It is well accepted that blocky, closely spaced M23C6 particles engulfed in c? improve rupture life and ductility by hindering GB sliding, thereby reducing nucleation and growth of cavities;[4,23–25] however, the presence of a wide and continuous M23C6 chain sandwiched by two layers of c? may facilitate crack propagation and lead to tensile and notch brittleness.[4,24] Hence, an optimum GB M23C6/c? structure where blocky M23C6 particles are engulfed in c? as densely as possible is desirable.[26–29]
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Fig. 6—Inhomogeneous distributions of M23C6 carbides in the 900 °C/10,000 h thermal exposure specimen: (a) loops, (b) chains, and (c) clusters of M23C6 carbides.

Fig. 7—GBs under the (a) heat treatment and (b) thermal exposure conditions of 800 °C/1000 h, (c) 850 °C/5000 h, (d) 900 °C/5000 h, and (e) 900 °C/10,000 h.
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The c?-free zones do not occur along the GBs in the present alloy. The reason is probably that the alloy is subjected to a correct heat-treatment regimen, that is, a solution anneal (4 hours at 1170 °C) followed by a hightemperature aging (4 hours at 1050 °C) and then a lowtemperature aging (16 hours at 850 °C). The triple heat treatment makes Cr distribute uniformly across the GBs, suppressing the Cr-depletion zones, and accordingly suppresses the c?-free zones.[30–32] Generally, the c?free zones are present only when the alloys are subjected to a dual heat treatment (a solution anneal followed by lower temperature aging), which results in a Cr depletion zone with a steep Cr concentration gradient along the GBs. The Cr lowers the solubility of c? forming elements, so its depletion adjacent to the GBs leads to the dissolution of the c? precipitates forming the c?-free zones,[33–35] which are generally thought to be deleterious to the properties of the alloys.[36,37] 5. g Phase At the exposure temperatures of 800 °C and 850 °C, a small quantity of needle-shaped g phase precipitates from the c matrix, whereas at 900 °C, there is almost no g phase present. Cracks that initiated from the g phase and extended into the matrix were not observed. Considering the small population of the g phase, its in?uence on mechanical properties is negligible. B. Degradation of Mechanical Properties The properties of c?-strengthened alloys depend on many factors, such as volume fraction and size of c?, solid-solution strengthening of c, and presence of hyper?ne c?.[26] However, these factors are modi?ed during thermal exposure in the present experiment, for example, leading to the coarsening of the secondary c? and the absence of the tertiary c?, which cause a signi?cant loss of tensile strength of the alloy (Figure 8(a)). It should be emphasized that although the volume fraction of the secondary c? increases with exposure (Figure 3(b)), it does not provide an essential improvement to the strength, for it not only results in virtually no new c? nucleus, but also withdraws solid-solution strengthening elements from the c matrix. Almost all the tensile or stress-rupture specimens failed exclusively along the GBs perpendicular to the applied load in the present study. Therefore, the GBs with a M23C6/c? structure are crucial to the rupture properties, namely rupture life and strain. During thermal exposure the stress-rupture lives of specimens rise ?rst with the density of the GB M23C6 particles and then decrease due to the development of a continuous carbide chain as well as c? coarsening, as presented in Figure 9(a). A life peak always occurs at a certain exposure time, which may correspond to an optimal M23C6/c? structure at the GBs; the higher the exposure temperature, the earlier the occurrence of the peak. The c? coarsening widens the c channels and facilitates the movement of mobile dislocations. Therefore, it is bene?cial to the elongation of the alloys. However, the formation of a continuous M23C6 chain at the GBs
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Fig. 8—Tensile-rupture properties of thermal exposure specimens: (a) tensile strengths and (b) tensile elongation.

tends to cause the premature onset of intergranular fracture and therefore reduces the elongation. This is especially true when a considerable amount of degenerated MC carbides that facilitate the initiation of microcracks distribute themselves at the GBs (Figures 8(b) and 9(b)).

IV.

CONCLUSIONS

1. The secondary c? precipitates are cuboidal in shape in the heat-treated specimens, whereas they assume a spherical appearance during long-term thermal exposure. As they coarsen, the secondary c? precipitates increase in volume fraction. The c? coarsening is the principal reason for the reduction of strength of the alloy. 2. The evolution of GB microstructure has a signi?cant e?ect on the stress-rupture life, which ?rst rises with the density of GB M23C6 particles and then decreases due to the formation and growth of continuous carbide chains. A peak occurs when the
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ACKNOWLEDGMENTS The authors thank Professor X.X. Jiang and Dr. J.S. Hou, Institute of Metal Research, Chinese Academy of Sciences, for their helpful discussions.

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Fig. 9—Stress-rupture properties of thermal exposure specimens: (a) stress-rupture life and (b) stress-rupture elongation.

M23C6/c? structure at the GBs is in an optimal form. 3. Primary MC degenerates via three reactions, that is, MC + c ? M23C6 + c?, MC + c ? M23C6 + g, and MC + c ? M23C6 + a-(W, Mo) + g. The MC degeneration zone is the preferred initiation site of microcracks. Its presence at the GBs promotes the onset of intergranular fracture and leads to the deterioration of properties. 4. The size and morphology of GI M23C6 carbides experience a signi?cant transition during long-term thermal exposure. As they coarsen, the carbides tend to change in shape from ?owerlike dendrites to irregular blocks to regular polyhedral particles, depending on the reduction in the total free energy of the alloy system. 5. The population of g phase is so small that its influence on mechanical properties is negligible.

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