Microstructural Changes of a Creep-Damaged Nickel-Based K002 Superalloy Containing Hf Element under Different HIP Temperatures

3.1.1 The original and degraded γ′ morphology

Figure 1:

The γ′ morphology for K002 superalloy. (a) Original microstructure; (b) damaged microstructure.

Cuboidal-shaped γ′ precipitates in the virgin superalloy are shown in Figure 1(a), which can excellently enhance the mechanical properties of superalloy through hindering the dislocation movement [5]. However, these γ′ precipitates were converted from cuboidal into irregular shape after long-term service, as shown in Figure 1(b). The γ′ precipitates were coalesced with the adjacent ones and their corners became blunt, and some of them were elongated resulting in the fact that the γ matrix channels between the γ′ precipitates gradually diminished. It is indicated that there was remarkable microstructural degradation in the service-exposed sample. The size of degraded γ′ precipitates was somewhat larger than that of initial alloy due to the γ′ coarsening, as seen in Table 2, but the volume fraction of γ′ precipitates was reduced as compared to that of initial alloy. The γ′ morphological evolution is driven by the reduction of the total internal energy in this system, and the microstructural changes are attributed to the competition between interfacial energy and elastic strain energy between γ and γ′ phases [15]. Before long-term service, the γ′ precipitates present cubic morphology since the elastic strain energy plays an important role in the competition between the interfacial and elastic strain energy. However, with the increase of service time, the contribution of elastic strain energy was gradually reduced, and the coherency between γ and γ′ was also gradually lost, resulting in the formation of irregular γ′ precipitates.

3.1.2 The γ′ morphology in HIP specimens under various HIP schedules

Figure 2:

The γ′ morphology under various HIP conditions. (a) 1,190°C/200 MPa/4 h; (b) 1,220°C/200 MPa/4 h; (c) 1,250°C/200 MPa/4 h.

Figure 2(a) illustrates the γ′ morphology under HIP condition of 1,190°C/200 MPa/4 h, consisting of coarsened primary γ′ and fine secondary γ′ precipitates. These primary γ′ precipitates still represented coarse and irregular morphology. Figure 3 shows the result of differential scanning calorimeter (DSC) measurement. The γ′ solvus temperature for K002 superalloy was about 1,222°C obtained from the DSC curve, indicating that HIP treatment under condition of 1,190°C/200 MPa/4 h cannot completely dissolve the coarse and irregular γ′ precipitates into the γ matrix owing to the lower HIP temperature. Both the size and volume fraction of γ′ precipitates were decreased as compared to those of the damaged alloy, as given in Table 2, implying that the γ′ dissolution behavior has played an important role in HIP processing and the residual coarse and irregular γ′ precipitates might be experiencing a dissolution process. When increasing HIP temperature to 1,220°C, the γ′ precipitates were transformed into relatively cubic shape since the temperature of HIP was comparatively close to the γ′ solvus temperature for K002 superalloy, as given in Figure 2(b). It is demonstrated that the primary γ′ precipitates can be completely dissolved into the γ matrix at HIP temperature of 1,220°C, and subsequently the γ′ phases can be re-precipitated and continuously grow as cubic morphology during the slower cooling stage of HIP, indicating that the elastic strain energy re-dominated the γ′ morphology during the γ′ growth. As HIP temperature was increased to 1,250°C, the γ′ precipitates were evolved to more regular shape, as shown in Figure 2(c). However, the size of γ′ precipitates under HIP conditions of 1,220°C/200 MPa/4 h and 1,250°C/200 MPa/4 h was remarkably smaller than that of initial alloy, as shown in Table 2, probably resulting from the insufficient γ′ growth. Therefore, it is necessary for the damaged superalloy to perform subsequent RHTs to obtain more regular and cubic γ′ morphology analogous to the original γ′ microstructure after HIP. However, HIP temperature cannot be excessively high since there is a risk of incipient melting of the superalloy [16, 17].

Figure 3:

DSC result for K002 superalloy.

3.2.1 Morphology of creep cavity after service

Figure 4 illustrates the large-sized creep cavities of turbine blade after long-term service, indicating that the blade has suffered serious degeneration. These cavities usually located at the triple junction, the interface between carbides and γ matrix, and grain boundary, can be coalesced with each other, finally leading to the formation of cracks which are quite dangerous for operation.

Figure 4:

Creep cavities after a long-term service.

3.2.2 Morphology of creep cavity after HIP

For the sample subjected to 1,190°C/200 MPa/4 h, the so-called concentrically oriented γ′ rafts elongated perpendicularly to the radial direction of creep cavity can be clearly seen in the microstructure around the HIP creep cavities, as shown in Figure 5(a). However, the γ′ precipitates far away from the creep cavity still present irregular orientations without significant directional coarsening. In our previous work [18–20], similar findings have been obtained for K465 superalloy. The concentrically oriented γ′ rafts ought to form when HIP temperature is not high enough to dissolve γ′ precipitates into the matrix in a large amount. It is sure that the evolution of the concentrically oriented γ′ rafts was complicated. The most probable reason could be explained as follows. The γ′ precipitates around the creep cavity were elongated under a complex stress field when the applied hydrostatic pressure was kept at P_h. The complex stress field can be simplified by a thick spherical shell model as follows. According to the Lame solution of elasticity theory [21], the triaxial stress state can be decomposed into three principal stress components, including a radial component σ_r and two tangential components σ_θ and σ_φ at a spherical polar coordinate. They can be expressed by the following equations:

where r is the radial distance from the cavity center, R_i and R_e are the internal radius and external radius of the shell, respectively. An effective stress σ_eff can be then obtained by the radial component σ_r and the tangential components σ_θ and σ_φ in accordance with the definition of Von Mises stress as follows:

Therefore, the complex stress state can be equivalently expressed by the effective stress σ_eff which actually represents a uniaxial tensile stress state in the radial direction, and it is found that there was a negative stress gradient around the cavity for σ_eff. In addition, the nickel-based superalloy before HIP usually had a negative value of the lattice misfit between γ matrix and γ′ precipitates owing to long-term service [22]. The preferential coarsening of γ′ precipitates was directionally arranged with a parallel direction (P type) or normal (N type) to the applied stress. According to Peng et al. [23], the γ′ precipitates would be elongated perpendicularly to the tensile stress direction when the lattice misfit was negative, which exhibited N-type γ′ rafting. Therefore, the concentrically oriented γ′ was N-type rafting. Meanwhile, the negative stress gradient near the cavity resulted in a chemical potential gradient causing the solute diffusion. Hence, the solute atoms were diffused inward the cavity replacing the vacancies. The concentrically oriented γ′ rafting only formed in the healing regions of creep-induced cavities, and it did not occur in the regions far away from the creep cavities. It was indicated that HIP at 1,190°C was insufficient to completely heal creep-induced creep cavities, implying that these cavities were experiencing a healing process.

As observed in Figure 5(b), the similar concentrically oriented γ′ rafting structure was also found, whereas creep cavity can be hardly observed. It is probable that some relatively small creep cavities in the sample subjected to 1,190°C/200 MPa/4 h had been completely healed, while the concentrically oriented γ′ rafting had not sufficiently vanished through solute diffusion.

Figure 5:

Concentrically oriented γ′ raft structure near the creep-induced cavity under HIP condition of 1,190°C/200 MPa/4 h. (a) Creep cavity can be observed; (b) creep cavity cannot be observed.

When increasing HIP temperature to 1,220°C which was extremely close to the γ′ solvus temperature for K002 superalloy, large amounts of creep-induced cavities had been healed and only a few small-sized cavities on grain boundary can be occasionally observed, as shown in Figure 6(a). The remnant cavity was generally located on the surfaces of γ/γ′ eutectics or carbides which were also the favorable nucleation positions for creep cavity owing to their different coefficient of thermal expansion with γ matrix during service [24]. Meanwhile, the concentrically oriented γ′ rafting around the cavity vanished since HIP at the sufficiently high temperature leads to the complete dissolution of γ′ precipitates into the γ matrix. Under HIP condition of 1,250°C/200 MPa/4 h, there were hardly large-sized creep cavities in the HIP sample. There still existed some small-sized creep cavities incompletely healed near γ/γ′ eutectics, and the cubic γ′ precipitates around the cavities were re-precipitated during the cooling stage of HIP, as shown in Figure 6(b).

Besides, some unstable morphology at the neck was observed indicating a healing feature of cavities during HIP. The morphological instability was analogous to the classical Rayleigh instability of a fluid cylinder under surface tension [25]. It is probably that some healed interfaces between the opposite surface of the cavity might also form when HIP temperature exceeds the γ′ solvus temperature due to the similar results found in our previous work for K465 superalloy [18, 19]. Therefore, it can be inferred that the solute atoms diffused inward the cavity to nucleate and grow the γ matrix in order to heal the creep cavities.

Figure 6:

Microstructure characteristics of different HIP samples. (a) 1,220°C/200 MPa/4 h; (b) 1,250°C/200 MPa/4 h.

3.2.3 Carbide changes before and after HIP

HIP treatment for K002 superalloy containing Hf seems to contribute to the formation of large amounts of blocky carbides, as shown in Figure 7. However, the phenomenon was not found in our previous work for K465 superalloy. The EDX results show that, as given in Table 3, main elements of carbides before HIP are Ta, W, indicating that the carbides are MC₍₁₎ type [17]. However, after HIP main elements of carbides are changed to Ta and Hf, Hf atoms have diffused into the carbides to substitute for W atoms to form MC₍₂₎-type carbides. The longer holding time of HIP might lead to the formation of lots of blocky carbides since Hf can be amply combined with C to form the MC₍₂₎-type carbides. Figure 8 shows the changes of volume fraction of carbides for K002 superalloy under different HIP conditions. The volume fraction of carbides in the damaged sample was slightly increased as compared to that of the initial specimen. For the sample subjected to 1,190°C/200 MPa/4 h, the volume fraction of carbides has not been obviously changed as compared to that of the initial and damaged specimen, but with the increase of HIP temperature to 1,220°C and 1,250°C, the volume fraction of carbides was markedly increased, demonstrating that HIP treatment at higher temperatures (e.g., 1,220°C and 1,250°C) can promote the formation of a large number of blocky carbides.

Figure 7:

Carbides for K002 superalloy after HIP.

Figure 8:

Changes of volume fraction of carbides for K002 superalloy under different HIP conditions.