1 Introduction

The materials applied in industrial gas turbine blades are expected to withstand high temperature and stress during long-term service. It is reported that the proper examining and repairing interval of gas turbine blades is 24,000 equivalent operation hours (EOH) and the total service life is beyond 50,000 EOH.[1,2] Extensive studies have shown that prolonged thermal and stress exposure causes multiple microstructural degradations such as γ′ coarsening and rafting, MC carbides decomposition, secondary precipitations along grain boundaries, and formation of TCP phases and cavity or void in the gas turbine blades after long-term service, almost all of which are detrimental to mechanical properties.[3–5]

The morphologies and mechanisms associated with γ′ particle coarsening upon aging have been a topic of significant discussion over the past decades.[6–10] Since the changes in morphology, size, and distribution of γ′ precipitates during long-term thermal exposure mainly cause the degradation of tensile strength and creep resistance, it is very important to monitor the evolution of γ′, be able to predict the kinetics of growth during long-term thermal exposure, and correlate that with mechanical properties. However, the evolution of size/morphology and volume fraction of γ′ precipitates during long-term thermal exposure beyond 20,000 hours and its impact on the creep properties of superalloys have been seldom reported in previous studies.

Compared to single crystal (SC) Ni-based superalloys, directional solidification (DS) Ni-based superalloys have a particular creep deformation characteristic due to MC decomposition in grain boundary (GB) region and grain boundary precipitates coarsening during long-term thermal exposure. Unfortunately, limited information is available for the coarsening kinetics of grain boundaries and the effects on creep property of DS Ni-based superalloys during long-term thermal exposure.

It is interesting to note that most of the above microstructural evolutions are reversible, in the sense that they can be rejuvenated by re-heat treatment and hot isostatic pressing (HIP).[11,12] Unfortunately, primary MC decomposition is irreversible and may affect the microstructure stability and mechanical properties of the rejuvenated blades during the next service cycle.[4,11–13] Three classical reactions, MC + γ → M23C6 + γ′, MC + γ → M23C6 + η, and MC + γ → M23C6 + (MC)deg, are often operative in the alloys, such as GTD-111 and In-738 applied in industrial gas turbine blades after long-term thermal and stress exposure.[3–5,14] The mechanism of primary MC decomposition in gas turbine blades made from both conventionally cast (CC) superalloys and directionally solidified (DS) alloys was reported in these previous studies.[3–5] However, there is not much published information on the time dependence of MC decomposition in terms of morphological and compositional evolution during long-term thermal exposure. Moreover, there is still augment regarding the role of MC carbides decomposition during creep after long-term thermal exposure. Some researchers thought that MC carbides decomposed into fine M23C6 particles, which played a positive role in influencing the long-term creep rupture properties of the DS superalloys,[15–18] whereas others insisted that MC decomposition provided the site of cracks initiation, which was deleterious and contributed to the failure of creep specimens.[5,19,20]

Obviously, during long-term thermal exposure, the mechanical performance of superalloys is closely related to the microstructural degradations. However, the effect of multiple microstructural degradations (e.g., γ′ coarsening, MC decomposition, and secondary precipitations coarsening along GBs) on mechanical properties of DS Ni-based superalloys during different stages of thermal exposure has seldom been reported in previous studies. In this paper, we studied the evolution of microstructure as well as the creep rupture properties during long-term thermal exposure. The relationships between microstructural evolution and stress rupture properties of the DS Ni-based superalloy have been correlated as a first attempt to explore the possible life assessment methodology.

2 Experimental

The nickel-based superalloy DZ411 in our present study is a typical hot corrosion resistant alloy applied in industrial gas turbine blades with 14.2Cr, 9.5Co, 3.1Al, 5.0Ti, 3.0Ta, 3.6W, 1.5Mo, with 0.1C and 0.01B, and balance Ni, in weight percent. The cylindrical specimens with the size of Φ16 mm × 220 mm were DS by the conventional Bridgman process. The specimens received the standard solution and age heat treatment, 1493 K (1220 °C)/2 hours/air cooling (AC), 1393 K (1120 °C)/2 hours/AC, and 1123 K (850 °C)/24 hours/AC. Then the as-heat-treated specimens were subjected to a long-term thermal exposure at 1173 K (900 °C) up to 24,000 hours. The thermally exposed specimens were machined into creep specimens in the longitudinal direction with gage length of 25 mm and a gage diameter of 5 mm to test the creep rupture property at 1253 K (980 °C) and 220 MPa.

Specimens after different periods of thermal exposure were cut and prepared by standard procedure for microstructural observations, i.e., they were polished and either etched electrolytically with a solution of 10 mL H3PO4 + 90 mL H2O under a voltage of 7 to 10 V for 10 seconds mainly for the microstructural examination of γ′ precipitates or etched chemically with a solution of 17 mL HNO3 + 100 mL HCl + 7.5 g CuCl2 + 7.5 g FeCl3 mainly for the microstructural examination of MC and M23C6 carbides. The microstructural examination of γ′ precipitation was performed using a scanning electron microscopy (SEM). Backscattered electron (BSE) imaging was used to observe the microstructure of MC and M23C6 carbides. Chemical composition of carbides was analyzed by electron probe micro-analyzer (EPMA). Size (equivalent diameter) and volume fraction of the γ′ particles were characterized on the SEM micrographs using Image Tool software. The equivalent diameter is defined as the diameter of a circle with an equal area of a feature of arbitrary shape. In order to evaluate average area fraction and equivalent diameter, more than 10 SEM micrographs were recorded from different areas for each sample, and 50 γ′ particles per micrograph were considered. The width of grain boundaries was characterized by measuring the average thickness of GB wrapped by γ′ precipitates and M23C6 carbides in the transverse direction of the specimens.

3 Results

3.1 Microstructure After Heat Treatment

The primary dendrite arm spacing of the as-cast samples is around 250 μm. Microstructure of the alloy after heat treatment was mainly composed of γ matrix, γ′ precipitates, residual γ/γ′ eutectics, and primary MC carbides within grain interiors (GIs) and along grain boundaries (GBs), as shown in Figure 1. The secondary γ′ precipitates were cubic, and the average size and area fraction were about 0.4 μm and 54 pct respectively. Numerous hyperfine tertiary γ′ precipitates forming during the aging treatment could be observed in the matrix channel (Figure 1(b)). The volume fraction of residual γ/γ′ eutectics was less than 1.5 pct. Primary MC carbides as block were scattered within grain interiors as well as along grain boundaries (Figure 1(c)). Chemical composition of primary MC carbides in the heat-treated alloy measured by EPMA is illustrated in Table I. It can be seen that, the main elements forming primary MC carbides are Ti, Ta, and W. γ′ precipitates along grain boundaries were cubic and the size was about 0.6 μm (Figure 1(d)).

Fig. 1
figure 1

Microstructure of the as-heat-treated alloy: (a) residual γ/γ′ eutectics and MC carbides, (b) γ′ precipitates, (c) primary MC carbides within GIs and along GBs, and (d) γ′ precipitates at grain boundary

Full size image
Table I Chemical Composition of Primary MC Carbides in the As-Heat-Treated Alloy
Full size table

3.2 Microstructural Degradation During Long-Term Thermal Exposure

During long-term thermal exposure at 1173 K (900 °C), the microstructure of the experimental alloy degenerates gradually, which depends on the aging time. Figures 2 through 6 show the typical microstructure evolution of γ′, primary MC carbides, and grain boundaries during the long-term thermal exposure.

Fig. 2
figure 2

Microstructure in the thermally exposed alloy for 5000 hours: γ′ precipitates (a) and primary MC carbides within GIs (b) and along GBs (c) and γ′ precipitates at grain boundary (d)

Full size image
Fig. 3
figure 3

Microstructure in the thermally exposed alloy for 9000 h: γ′ precipitates (a); primary MC carbides within GIs (b) and along GBs (c); γ′ precipitates at grain boundary (d); EDX spectra of the discrete gray particles on the periphery of primary MC (e)

Full size image
Fig. 4
figure 4

Microstructure in the thermally exposed alloy for 15,000 h: γ′ precipitates (a) and primary MC carbides within GIs (b) and along GBs (c) and γ′ precipitates at grain boundary (d)

Full size image
Fig. 5
figure 5

Microstructure in the thermally exposed alloy for 20,000 h: γ′ precipitates (a) and primary MC carbides within GIs (b) and along GBs (c) and γ′ precipitates at grain boundary (d)

Full size image
Fig. 6
figure 6

Microstructure in the thermally exposed alloy for 24,000 h: γ′ precipitates (a); primary MC carbides within GIs (b) and along GBs (c); M23C6 particles wrapped by γ′ precipitates (d); γ′ precipitates at grain boundary (e)

Full size image

3.2.1 Evolution of γ′ precipitates

A significant change of morphology, size, and area fraction of γ′ precipitates can be found during long-term thermal exposure at 1173 K (900 °C). In the specimen exposed for 5000 hours, tertiary γ′ precipitates dissolved and the secondary γ′ particles became larger and degenerated to the cube morphology with rounded corner, which widens the matrix spacing between the γ′ precipitates (Figure 2(a)). Coarsening process of secondary γ′ precipitation was going on as the thermal exposure was prolonged to 9000 hours (Figure 3(a)). As the thermal exposure was over 15,000 hours, the large irregular plate morphology of γ′ precipitation was present over the entire sample due to the growth and coalescence of smaller γ′ precipitates (Figures 4(a) through 6(a)). As shown in Figure 7, size (equivalent diameter) of secondary γ′ precipitates increased gradually, and there was a tiny change to be seen for area fraction of secondary γ′ with the increasing aging time.

Fig. 7
figure 7

Evolution of size and area fraction of secondary γ′ precipitates during long-term thermal exposure at 1173 K (900 °C)

Full size image

3.2.2 Decomposition of primary MC carbides

Primary MC carbides remained stable and only dissolved at the edge within the first 5000 hours exposure as seen in Figures 2(b) and (c). Similar to the previous observation,[3–5] a thin well-defined decomposition zone was also observed on the peripheries of primary MC carbides both within GIs and along GBs as seen in Figure 2(b). As is reported,[3–5] the chemical composition at the decomposition zone is consistent with that of MC carbides, with part of the Ti and Ta atoms replaced by Ni, Co, and Cr.

As the thermal exposure was extended to 9000 and 15,000 hours, fine gray particles were found in the vicinity of primary MC within GIs as well as at grain boundaries (Figures 3(b) and (c)). The discrete gray particles were identified by EDS as M23C6 carbides with high Cr content (Figure 3(e)).

As the thermal exposure was prolonged to 20,000 and 24,000 hours, MC decomposition proceeded and more coarsened M23C6 particles precipitated within GIs as well as at GBs, as shown in Figures 5 and 6. In addition, the M23C6 particles were wrapped by γ′ as illustrated in Figure 6(d). Similar to the classical reaction,[5] the decomposition of primary MC can be summarized as MC + γ → M23C6 + γ′. Meanwhile, more decomposed MC began to break down into disintegrated MC fragments (<2 μm) with the extended aging time, as indicated by the black arrow in Figure 5(b), and variations of the percentage of primary MC carbides showing disintegration vs aging time both within GIs and at GBs are shown in Figure 8. Moreover, it should also be noted that primary MC carbides located at or near the GBs exhibited a higher degree of decomposition than those in the GIs.

Fig. 8
figure 8

Variations of the percentage of MC showing disintegration vs aging time both within GIs and at GBs

Full size image

3.2.3 Grain boundary evolution

Discrete M23C6 particles wrapped by γ′ precipitated along GB within the first 5000 hours at 1173 K (900 °C) as shown in Figures 2(c) and (d). As the thermal exposure was prolonged up to 24,000 hours, M23C6 particles wrapped by γ′ in grain boundaries region continued to coarsen and merged with each other, and consequently, resulted in the formation of a continuous layer of M23C6 and γ′ precipitates (Figure 6). The width of the M23C6 and γ′ layer at GBs during thermal exposure is shown in Figure 9. It can be seen that the coarsening rate of GB increases within the first 10,000 hours and then slows down during further exposure.

Fig. 9
figure 9

The width of grain boundaries after thermal exposure at 1173 K (900 °C) for different aging times

Full size image

3.3 Degradation of Stress Rupture Property

The stress-rupture lifetime of the thermally exposed specimens at 1253 K (980 °C) reduces remarkably with the extended thermal exposure, and the elongation increased from 16 to 31 pct within 5000 hours exposure and then decreased to 23 pct up to 24,000 hours, as illustrated in Figures 10(a) through (b). The evolution of stress-rupture lifetime can be divided into three stages. The stress-rupture lifetime decreases sharply by 50 pct within 5000 hours and then remains at a steady level from 5000 to 15,000 hours. Finally, the lifetime of stress rupture decreases drastically again to about 20 hours as the thermal exposure time is over 20,000 hours.

Fig. 10
figure 10

Stress-rupture lifetime (a) and elongation (b) of the thermally exposed alloy

Full size image

The microstructure of longitudinal section near the rupture site of the thermally exposed specimens for 5000 and 24,000 hours is shown in Figure 11. Compared to the thermally exposed specimens within 5000 hours, more microcracks in the thermally exposed specimen for 24,000 hours were observed near the rupture site (Figures 11(a) and (b)). It is clearly shown that numerous microcracks initiated and grew in the position of curve GBs (Figure 11(c)), where decomposed MC and coarsened M23C6 and γ′ precipitated after 24,000 hours of thermal exposure (Figure 11(d)).

Fig. 11
figure 11

Backscattered electron (BSE) imaging of longitudinal section near the rupture site in the creep rupture specimens of the thermally exposed alloy for 5000 h (a) and 24,000 h (b, c and d)

Full size image

4 Discussion

4.1 Microstructural Degradation During Long-Term Thermal Exposure

4.1.1 γ′ evolution

In some alloys, the time-dependent γ′ coarsening process conforms to the LSW model:[6,7]

$$ \bar{D}_{t}^{3} - \bar{D}_{0}^{3} = kt, $$

where \( \bar{D}_{t} \) is the mean particle size at a time t, \( \bar{D}_{0} \) is the mean particle size at the onset of coarsening (t = 0), and k is the coarsening rate constant.

Tertiary γ′ precipitates dissolved during the thermal exposure at 1173 K (900 °C) within 5000 hours aging. With the thermal exposure prolonged, coarsening and coalescence of secondary γ′ precipitates occurred, which is driven by the reduction in total interfacial energy. The variations of \( \bar{D}_{t}^{3} - \bar{D}_{0}^{3} \) vs time for secondary γ′ precipitates are shown in Figure 12, according to our present experimental date shown in Figure 7. It is clear that the coarsening kinetic for the secondary γ′ precipitates follows approximately a linear relation, which conforms well to the LSW theory.

Fig. 12
figure 12

Variations of \( \bar{D}_{t}^{3} - \bar{D}_{0}^{3} \) vs time for secondary γ′ precipitates

Full size image

4.1.2 Carbides decomposition

Compared to the primary MC carbides in the virgin alloy (Figure 1(c)), primary MC carbides in the thermally exposed alloy within 5000 hours were surrounded by a decomposed zone. During thermal exposure within 5000 hours, the primary MC carbides with high carbon content serve as a carbon resource and γ matrix as a Ni, Cr, and Co reservoir. Outward diffusion of carbon from the primary MC carbides into the surrounding matrix and inward diffusion of Ni, Cr, and Co from the γ matrix develop a well-defined transition zone around the primary MC carbides.[5] The replacement of the strong carbide forming atoms such as Ti, Ta, and W with Ni, Co, and Cr atoms resulting from the outward and inward diffusion is known to weaken the interatomic bonds in the primary MC carbides and results in the decrease of their stability and disintegration in the prolonged thermal exposure.[3]

With the thermal exposure extended over 9000 and 15,000 hours, primary MC carbides were subjected to further decomposition due to the enrichment of the elements of C and Cr through the outward and inward diffusion between MC and γ matrix, creating the favorable environment for the formation of M23C6 precipitations in the vicinity of the decomposed primary MC carbides.

With the thermal exposure extended to 20,000 and 24,000 hours, the decomposition of primary MC carbides proceeded and more decomposed MC began to break up into disintegrated MC-type particles, especially in GB region. Moreover, the precipitations of M23C6 carbides along GBs resulted in the local enrichment with γ′ forming elements, Al, Ti, and Ta, allowing γ′ film to form along GBs.[21] It is obvious that primary MC carbides along GBs underwent a more severe decomposition than those within GIs due to faster diffusion and enrichment of elements such as C and Cr through the diffusion path along GBs, creating the favorable environment for severe disintegration of primary MC and the formation of considerable M23C6 precipitations. The disintegrated MC particles have a much higher surface/volume ratio than primary MC, and outward and inward element diffusion rate for the particles is considerably faster as compared to primary MC decomposition. Hence, the decomposition rate of the disintegrated MC particle might be faster than primary MC carbides during the extended thermal and stress exposure, and the related work will be done in the near future.

4.1.3 Grain boundary (GB) evolution

Grain boundary coarsening during long-term thermal exposure is mainly facilitated by M23C6 and γ′ precipitates coarsening in the GBs region. GB energy is the main driving force of elements diffusion along GBs, which significantly influences M23C6 and γ′ precipitates coarsening and MC degeneration along GBs during the long-term thermal exposure.

Different from the coarsening behavior of γ′ precipitates within GIs via Ostwald ripening, GB coarsening is associated with directional coarsening and coalescence of M23C6 and γ′ precipitates via faster GB diffusion, for which LSW theory can’t be applied directly.[22] The Johnson–Mehl–Avrami–Kolmogorov (JMAK) theory has been applied to quantify microstructural evolution in isothermal situation.[22–25] The JMAK equation for an isothermal transformation can be written as[25]

$$ f\left( t \right) = 1 - \exp ( - kt^{n} ), $$

where k is a constant and n is an integer or half integer. The variations of ln[−ln(1−f(t))] vs lnt follow a linear line, according to the equation above. The variations of width of coarsened GB during long-term isothermal exposure can be described as[22]

$$ f\left( t \right) = \frac{{W_{t} - W_{0} }}{{W_{\text{f}} - W_{0} }}, $$

where W t is the mean width of GB at a time t of the thermal exposure, W 0 is the mean width of GB at the onset of the thermal exposure (t = 0 hours) and W f is the mean width of GB at the final state of coarsening (t = 24,000 hours).

The variations of ln[−ln(1− f(t))] vs lnt in our study are shown in Figure 13. The GB coarsening kinetic follows approximately a linear line, which conforms well to the JMAK theory.

Fig. 13
figure 13

Variations of ln[−ln(1−f(t))] vs lnt during long-term isothermal exposure at 1173 K (900 °C)

Full size image

4.2 Influence of Microstructural Degradation on Stress Rupture Property

The DS DZ411 nickel-based superalloy is strengthened primarily by γ′ phase. The mechanical properties depend on the morphology, size, volume fraction, and distribution of the γ′ precipitates.[26–30] Compared to γ′ precipitates in the “new” alloy, hyperfine tertiary γ′ precipitates disappear and secondary γ′ precipitates degenerate to the cube morphology with rounded corner, which enlarges the spacing between the γ′ precipitates after 5000 hours thermal exposure, and it is easier for the dislocations to move in the γ matrix. Moreover, it has been reported that when the size of γ′ precipitates ranges from approximately 0.5 to 0.8 μm (approaching to the evolution of γ′ precipitates between 0 and 5000 hours thermal exposure), the creep life is inversely proportional to the size of γ′ precipitates.[30] Moreover, the initial GB coarsening and stable primary MC carbides within 5000 hours are generally considered to have a negligible influence on the initiation and propagation of microcracks, as illustrated in Figure 11(a). Therefore, a significant decrease of stress rupture life of the aged alloy within 5000 hours was observed.

Although γ′ precipitates continue coarsening and volume fraction decreases, there is no obvious decrease of the stress-rupture lifetime during the prolonged thermal exposure between 5000 and 15,000 hours. Based on previous investigation,[30] lattice mismatch appears to be an important factor in the consideration of the influence of γ′ particle size on the high temperature creep properties of some alloys. As γ′ precipitates are subjected to further coarsening and coalescence, the γ′ particles become a more spherical and irregular morphology and semi-coherent or non-coherent, and a higher magnitude of mismatch may result in a finer spacing of misfit dislocations, which in turn would act as a stronger barrier to mobile dislocations.[30] Moreover, the M23C6 precipitates along GB would improve the resistance to GB slip and have a beneficial effect on stress rupture property. Consequently, the stress-rupture lifetime presents no significant change between 5000 and 15,000 hours thermal exposure.

As the thermal exposure was prolonged from 20,000 up to 24,000 hours, GB coarsening facilitated by coarsened γ′ layers and severe MC carbides decomposition along GBs took place, which may mainly contribute to the secondary decrease of stress-rupture lifetime shown in Figure 10(a). As was reported,[31–33] when the continuous GB structure wrapped by M23C6 and γ′ layer severely coarsen, GB is embrittled and become notch sensitive and prone to cracking as illustrated in the fracture (Figures 11(c) and (d)). Although the aim of producing DS superalloy is to eliminate transverse grain boundaries for improving creep properties, GB in the DS superalloys is not completely straight but an approximate linear line with curves. The GB in the curve position performs like a small scale transverse GB during creep. Hence, the coarsened curve grain boundaries after long-term thermal exposure would have significant detrimental effects on the decreasing of stress-rupture lifetime. Consistent with the conclusions in the previous studies,[5,19,20] severe MC decomposition and continuous M23C6 and γ′ precipitates near the positions of curve GBs provided the site of cracks initiation and propagation, which reduce the force-bearing area and result in a higher stress level and creep rate, but have not decreased the elongation obviously. Meanwhile, primary MC decomposition within GIs might have no significant influence on stress rupture property of the thermally exposed alloy.

5 Conclusions

  1. 1.

    During long-term thermal exposure at 1173 K (900 °C), microstructural degradation of the DS alloy includes γ′ coarsening, primary MC decomposition, and GB coarsening. γ′ precipitate gradually changed its morphology from cubic to spherical and irregular. The coarsening kinetic appears to conform well to the LSW model during long-term thermal exposure. Primary MC decomposed to a tiny degree and was surrounded by a well-defined decomposition zone within 5000 hours. As the thermal exposure was extended over 9000 hours, primary MC decomposed into M23C6 carbides wrapped by γ′, which can be summarized as MC + γ → M23C6 + γ′. As the thermal exposure was prolonged up to 20,000 and 24,000 hours, the MC decomposition proceeded and more decomposed MC broken down into smaller MC-type disintegrated particles. GB coarsening was mainly facilitated by M23C6 carbides and γ′ precipitates coarsening in GBs region, and the GB coarsening kinetics conforms well to the JMAK theory.

  2. 2.

    The stress-rupture lifetime at 1253 K (980 °C) exhibits three stages during thermal exposure. The stress-rupture lifetime decreases sharply by 50 pct within 5000 hours and then remains at a steady level from 5000 to 15,000 hours. Finally, the lifetime of stress rupture decreases drastically again to about 10 pct of that of the virgin alloy as the thermal exposure time is over 15,000 hours.

  3. 3.

    During different stages of the thermal exposure, dominant factors for the decrease of stress-rupture lifetime vary due to the evolution of multiple microstructures. Secondary γ′ coarsening and widening of γ matrix channel facilitate the movement of mobile dislocations, which mainly cause a significant decreasing of stress rupture life of the thermally exposed alloy within 5000 hours. Although γ′ coarsening continues between 5000 and 15,000 hours thermal exposure, the γ′ particles become a more spherical and irregular morphology and semi-coherent or non-coherent, and a higher magnitude of mismatch may result in a finer spacing of misfit dislocations, which in turn would act as a stronger barrier to mobile dislocations. Moreover, the discrete M23C6 precipitates along GB would improve the resistance to GB slip and have a beneficial effect on stress rupture properties. Due to these effects, the stress-rupture lifetime of the alloy between 5000 and 15,000 hours presents no obvious change. The severe MC decomposition and continuous coarsened M23C6 and γ′ precipitates at the curve grain boundaries facilitate the initiation and propagation of microcracks after the thermal exposure from 20,000 up to 24,000 hours and result in the drastic drop of the stress rupture life.