Topologically Closed Packed Phase and Its Interaction with Dislocation Movement in Ni–Based Superalloy during High–Temperature Creep

Abstract

In superalloys, topologically close–packed (TCP) phases, which contain refractory elements, usually significantly influence the mechanical properties. The current work investigates the structure and composition of the TCP phase in an Al–Mo–rich Ni–based single crystal superalloy. It is shown that after 40 h of thermal exposure, a large number of strip–like TCP phases are formed, which are enriched in Mo and Re. The structure of the TCP phase is identified as the tetragonal σ phase with the lattice parameter a being 0.93 nm and c being 0.50 nm. During the creep process, the single crystal tilts obviously and leads to orientation variation from <1 1 0> direction. Two groups of dislocations are observed in the deformed sample. One group contains straight dislocation lines and another group contains dislocation networks. The interaction between TCP phase and dislocation in the single crystal superalloy is studied to reveal the effect of the TCP phase on the deformation behavior. During creep, the σ phase hinders the dislocation movement, which may contribute to the propagation of the cracks and the final fracture.

1. Introduction

Ni–based single crystal (SC) superalloys usually include many elements, such as Ni, Al, Mo, Cr, Co, W, Ti, Ta, Re and Ru [1]. These constituent elements can be classified into two categories, i.e., γ’ phase forming elements and γ phase forming elements [2,3,4,5,6,7]. The γ’ phase, Ni₃(Al, Ta), is a simple cubic L₁₂ crystal structure with definite occupation of Al(Ta) and Ni atoms. Al(Ta) atoms are at the corners of the cube and Ni atoms are at the center of the plane. Thus the γ’ stoichiometric composition is Ni₃(Al,Ta). The γ phase forming elements refer to alloying elements distributed mainly in the γ matrix. The representative elements include Mo, Co Cr, W, Re and Ru [8]. These γ phase forming elements could significantly increase the solidus temperature and decrease the stacking fault energy, enhancing thermal stability and creep properties [9,10]. In addition, γ phase forming elements will have a great influence on the microstructure, diffusion and dislocation motion of SC superalloys, which could bring tailored mechanical properties. For example, Re is often added into the γ phase to slow down the dislocation motion and affect the mechanical properties [11,12,13,14]. According to Wu, the Re enrichment in dislocations provides direct evidence of solute drag and pipe diffusion, resulting in a reduced dislocation velocity and transport of matrix elements into the γ’ precipitates [15].

However, the addition of refractory elements (e.g., Re, W and Mo) [15] often results in the precipitation of brittle topologically close–packed (TCP) phases (σ, μ, P, and R, etc.) [16,17] due to the congregation during high–temperature service, and these TCP phases are usually harmful to creep properties. Therefore, to guarantee the creep performance of SC superalloys, it is necessary to investigate the evolution of the TCP phase, as well as its influence on the deformation behavior. TCP phases are extensively studied in traditional Ni–based single crystal superalloys, such as the morphology of the TCP phase [18], the orientation relationships between TCP phases and matrix [19], and the effect of time and temperature on the TCP precipitation [20]. In recent years, single crystal superalloys with high Al and Mo contents draw more researchers’ attention. These single crystal superalloys usually have higher γ’ volume fraction and good creep performance at an ultrahigh temperature [21,22]. Meanwhile, there is a large tendency for the precipitation of the TCP phase in these alloys, since the low γ volume fraction easily leads to supersaturation of refractory elements and the precipitation of the TCP phase.

However, there is limited research on the structure and evolution of the TCP phase in single crystal superalloys with high Al and Mo contents. Ai Cheng et al. [23] studied the microstructure stability of TCP in Al–Mo–rich Ni–based single crystal superalloys under unstressed long–term (250 h) conditions, and the TCP phases in both long–term aged cellular and dendritic samples were identified as P phases. Whether the structure of the TCP phase will remain stable or change during thermal exposure is still unknown. Furthermore, the interaction of the TCP phase and dislocations during the creep process also remains unclear. In this work, we systematically investigate the structure of the TCP phase in the Al–Mo–rich Ni–based model SC superalloy and analyze the evolution of the TCP phase during the thermal exposure process at 1100 °C. The dislocation evolution and the interaction with the TCP phase of the Al–Mo–rich Ni–based SC superalloy are also investigated.

2. Experimental

The experimental alloy used in this work was an Al–Mo–rich Ni–based model SC superalloy with the nominal composition

Table 1. Composition of the used Ni₃Al–based single–crystal superalloy (mass %).

Element	Ni	Al	Cr	Mo	Re	Ta
Content	Excess	7.6–8.3	1.5–2.5	9–13	0.5–1.5	2.4–4.0

. Directional solidification experiments were carried out by a high rate solidification (HRS) apparatus. After directional solidification, the model SC superalloy underwent a solution treatment (1300 °C/2 h + 1305 °C/2 h + 1310 °C/2 h + 1315 °C/2 h + 1320 °C/4 h + 1325 °C/6 h, air cooling). Creep tests were carried out according to GB–2039–2012. The alloy composition is shown in Table 1. For the thermal exposure experiment, the sample after solution treatment was exposed to an air atmosphere at 1100 °C for 40 h. A normal creep (maximum strain is ~80%) experiment was conducted using the sample after solution treatment at 1100 °C and 140 MPa, and slices were cut from 1 mm and 10 mm away from the fracture surface for microstructure observation.

The cut surface was mechanically ground with silicon carbide abrasive paper (P200 to P3000), and then polished with diamond paste to a mirror level. The etching solution used to reveal the γ’ phase is 20 g CuSO₄ + 100 mL HCl + 100 mL C₂H₅OH. Microstructure characterizations were conducted by Apreo S LoVac scanning electron microscope (SEM) equipped with an energy dispersive spectroscopy (EDS) system operated at 20 kV and electron backscatter diffraction (EBSD) system (orientation analysis) operated at 25 kV. EDS experiments were conducted to analyze the composition of the TCP phase, and EBSD experiments were conducted to analyze the orientation variation after deformation. In order to further analyze the structure of the TCP phase, focused ion beam technology was used to make a transmission electron microscope sample using Helios G4 CX equipment. Then, the morphologies of TCP phase and dislocations of SC superalloy were identified by a JEM 2100 TEM operated at 200 kV.

3. Results and Discussion

Figure 1 shows the microstructure of the solution–treated Al–Mo–rich Ni–based SC superalloy. The microstructure consists of cubic γ′–precipitates embedded in the γ–matrix. The fractions of two phases are obtained by statistically counting white (γ′–phase) in the images. The volume fraction of γ′–precipitate of the solution–treated SC superalloy is calculated as 72 ± 3%, the γ′–cube sizes as 405 ± 20 nm, and the γ–channel width as 60 ± 6 nm. No TCP phase is observed after solution treatment. Figure 1c shows the creep time–strain curve under 1100 °C/140 MPa and the average rupture time is ~42 h and the average strain is ~86%. Figure 2a,b exhibit the microstructure of Al–Mo–rich Ni–based SC superalloy after 40 h thermal exposure. When observing from the (1 1 0) plane, a large amount of stripe–like TCP phases is regularly arranged, which can be classified into two types with different widths of 400 ± 25 nm and 125 ± 15 nm. Figure 2c shows the EDS results of spot 1 to spot 4 in Figure 2b, and both the two types of TCP phases are enriched in Re and Mo elements.

TEM is employed to study the structure of the TCP phase. Figure 3a shows the SEM image of Al–Mo–rich Ni–based SC superalloy after hot exposure. FIB is used to cut a TEM sample from the red rectangle region in Figure 3a, and the TEM lamellar sample is shown in Figure 3b with a TCP phase in the center. Figure 3c exhibits the SC superalloy’s bright field TEM image, which contains an interface of matrix and TCP phase. Figure 3d shows the diffraction patterns of the matrix and interface with the beam from [1 1 1] direction, and the TCP is confirmed as the tetragonal σ phase. According to the diffraction patterns, it has been determined that the crystallography relationship of (1 1 1)_γ//(0 0 1)_σ is maintained between the γ and σ phases.

In order to further characterize the crystal parameter of the σ phase, high–resolution TEM is employed. Figure 4a shows the high–resolution TEM image of the σ phase, and Figure 4b is the fast Fourier transform patterns of the red rectangle in Figure 4a, and the red circles in Figure 4b are the masks. Figure 4c exhibits the inverse fast Fourier transform image containing the crystal lattice. The crystal plane space in (0 0 1) plane of σ phase of Al–Mo–rich Ni–based SC superalloy after 40 h thermal exposure is calculated to be 0.93 nm, indicating that lattice parameter a of the σ phase is 0.93 nm.

In comparison, the TCP phase in Mo–rich Ni–based SC superalloy after 250 h thermal exposure treatment is identified as P phase [23]. Thus it is speculated that the temporality of the TCP phase evolution is from σ phase to P phase. Similarly, according to RC Reed’s work [24], in traditional Ni–based SC superalloy, the flaky σ phase is first precipitated and then transformed into the P phase. Increasing the molybdenum stabilizes P phase at the expense of µ phase and σ phase. In our Al–Mo–rich Ni–based SC superalloy, the initially precipitated σ phase is rich in refractory elements Mo and Re. As the precipitated phase grows up, the surrounding γ’ phase restricts the diffusion of refractory elements to the precipitated phase, which impedes the growth of the TCP phase. Meanwhile, the phase transformation (σ phase to P phase) is promoted by self–diffusion of the refractory element in the σ phase.

After creep tests, EBSD is used to study the orientation variation from <1 1 0> direction of the SC superalloy sample. Figure 5a–c shows the inverse pole figure of (1 −1 0) plane from X direction, Y direction and Z direction of the region which is 10 mm away from the fracture surface, respectively. The inserts are the misorientation angle of the line in each figure, and the maximum misorientation across the line (~400 µm) is ~8°. While Figure 5d–f shows the inverse pole figure of (1 −1 0) plane from X direction, Y direction and Z direction of the region which is 1 mm away from the fracture surface, respectively, and the maximum misorientation across the line (~400 µm) is ~13°. This misorientation indicates the single crystal is obviously twisted during creep, and the deformation mechanism and dislocation behavior are revealed in the following discussion part.

Figure 6 exhibits the bright field TEM image of the region which is 10 mm away from the fracture surface, and the insert is a diffraction pattern with the beam from the [1 1 0] direction. Two groups of dislocations are observed in the deformed sample. One group contains straight dislocation lines parallel to each other, which are shown in Area_1. Another group contains dislocation networks in Area_2, which is near the σ phase. We observed a series of images and did not find any twins, which indicates that the creep mechanism of the SC superalloy is dislocation motion. According to our previous work and other researchers’ work, the creep mechanism of Al–Mo–rich Ni–based SC superalloy is mainly dislocation slip. Figure 7 shows the enlarged TEM image of Area_2 in Figure 6. Figure 7a–c shows the bright field TEM images of the dislocation networks of the region which is 10 mm away from the fracture surface with two beam conditions. According to the extinction law in Figure 7d, the dislocations networks which are near the strip–like σ phase exhibit a b = a/2<1 −1 0> or b = a/2<0 1 −1> Burgers vector. Figure 8 shows the enlarged TEM image of area 1 in Figure 6. Figure 8a–c shows the bright field TEM images of the straight dislocation lines. According to the extinction law in Figure 7d, the straight dislocation lines exhibit a b = a/2<1 −1 0> or b = a/2< 1 0 1> Burgers vector. This suggests that during creep, the dislocations piled up near the strip–like σ phase may lead to the stress concentration, which will promote the initiation and propagation of microcracks and accelerate the fracture of the SC superalloy [25,26].

Figure 9 shows the interaction of the σ phase and dislocations in the deformed Al–Mo–rich Ni–based SC superalloy. Figure 9a exhibits the bright field image of the deformed sample 10 mm away from the fracture surface, and several bending dislocations are observed near the σ phase, indicating the stress is easily concentrated in this region, which will promote the initiation and propagation of the cracks to accelerate the fracture. The crystal structure of σ phase is confirmed by the diffraction patterns in Figure 9b,c. Figure 9e shows the diffraction pattern of the interface between σ phase and γ matrix in Figure 9d, which indicates the orientation relationship of [1 1 0]_γ//[1 1 0]_σ. The lattice parameter c is calculated to be 0.50 nm. According to RC Reed’s work [24], there is a consistent orientation relationship between σ and the γ matrix, which is also confirmed by our results, and dislocations in the interface which extend from the edges of the precipitates are indications of a high degree of coherency in the interface between σ and the γ matrix. In order to observe the dislocation movement, we tilt α in the range of −35° to 35° and record the σ phase every 2°. Figure 9f,g are bright field images of the σ phase with ~6° α tilt variation, and throughout emerging Figure 9f,g, a stereoscopic morphology is obtained in Figure 9h. Dislocations can be observed in the region which is near the σ phase. The σ phase which is enriched in Re and Mo experiences random nucleation and diffusion–controlled growth [20]. The growth of σ phase greatly consumes the surrounding elements to form a depleted zone that surrounds the σ phase. As the growth primarily consumes γ–rich elements, e.g., Mo and Re, the depleted zone is the intrinsic γ′ phase. Since the large depleted zone is prone to stress concentration, dislocations that emit from the interface between TCP and the depleted zone facilitate dislocation interaction and microcrack initiation in the large depleted zone, promoting the propagation of the cracks and accelerating the fracture. In the future, more quantitative research, e.g., dislocation–based phase field [27,28,29], should be made to further understand the interaction between dislocation and phase microstructures of Ni–based superalloys.

4. Conclusions

In the current work, the structure of the TCP phase in the Al–Mo–rich Ni–based model single crystal superalloy is systematically investigated, as well as analyzing the possible evolution of the TCP phase. Meanwhile, the dislocation behavior of the Al–Mo–rich Ni–based SC superalloy is studied to reveal the role the TCP phase played during the creep.

In Al–Mo–rich Ni–based model SC superalloy, no TCP phase is observed after solution treatment. After thermal exposure for 40 h, a large number of strip–like TCP phases enriched in Mo and Re elements are formed. The structure of the TCP phase is identified as the tetragonal σ phase with the lattice parameter a being 0.93 nm and c being 0.50 nm.

During the creep process, the single crystal is tilted to result in obviously orientation variation from <1 −1 0> direction. Two groups of dislocations are observed in the deformed sample. One group contains straight dislocation lines with Burgers vector of b = a/2< 1 −1 0> or b = a/2<0 1 −1>. Another group contains dislocation networks with Burgers vector of b = a/2<1 −1 0> or b = a/2< 1 0 1>. During creep, dislocations piled up near the strip–like σ phase lead to the stress concentration, which may promote the initiation and propagation of microcracks and accelerate the fracture of the SC superalloy.