Improving Intermediate Temperature Plasticity of Co–Al–W–Base Superalloy: Based on Optimizing MC Carbides

Abstract

The brittleness of the new Co–Al–W–base superalloys is said to be a fatal weakness for their application in the intermediate temperature range, especially for polycrystal Co–Al–W–base superalloys. MC carbide controlling is used to improve the intermediate temperature plasticity of the Co–Al–W–base superalloy in the present study. Both microstructure characterization and interrupted tensile tests were performed to investigate the mechanism of improving the plasticity of the superalloy by MC carbides at an intermediate temperature of 800 °C. The results show that the plasticity of the superalloy is mainly dependent on MC carbides. MC carbide breakers appear after yielding, which closely depends on the morphology and size of MC carbides. Based on the experimental and calculation results, it is found that the intermediate temperature plasticity of the superalloy can be guaranteed by controlling the critical mean size of MC carbides, and the desired mean size of MC carbides should be less than 47 μm.

1. Introduction

In 2006, Ishida’s group reported a type of Co–Al–W–base superalloy that is strengthened through L1₂-ordered γ’–Co₃(Al, W) [1]. This superalloy shows higher strength and better creep performance at elevated temperatures, compared to the traditional Co-based superalloys strengthened via carbides [2]. This superalloy also possesses higher solidus and liquidus temperatures (about 100–150 °C) than Nickel-based superalloys, and thus becomes a good candidate material for next-generation gas turbine blades [3]. Gas turbine blades serve in complicated stress and high-temperature conditions [4]. Therefore, the mechanical properties of Co–Al–W–base superalloys have been the main concern of researchers. Abundant research about the effect of alloy elements (Ti [5,6], Ta [7,8], Mo [9,10], Ce [11]) on the yield/tensile strength of Co–Al–W–base superalloys was performed. However, the plasticity of Co–Al–W–base superalloys is rarely mentioned. Moreover, the intermediate temperature embrittlement still remains today, which is the fatal weakness of Co–Al–W–base superalloys.

In the past several years, our group developed a new multicomponent Co–Al–W–base superalloy [12]. In order to improve the creep property of the superalloy, appropriate content of carbon has to be added, and thus MC carbides appear. MC carbide usually forms before solidification in the Co–Al–W–base superalloy [13], which is composed of metallic elements and carbon, such as TiC, TaC, WC, NbC, etc. MC carbides are mainly used to strengthen the high-temperature strength in traditional Cobalt-base superalloys due to their high resolution temperature (about 1280 °C) and microstructure stability [14]. In addition, MC carbide also can hinder dislocation slip, enhance high-temperature strength, and especially strengthen the grain boundary [15,16,17]. It is known that MC carbide is sensitive to the plasticity of the superalloy because it belongs to a kind of brittle hard phase [18]. The variations in the size, morphology, and distribution of MC carbide have significant effects on the high-temperature deformation process of conventional Co-based superalloys [19,20,21,22,23,24,25]. So far, there is no suitable method to eliminate the effect of MC carbide on fracture behavior except to control the C content [26]. Although in Co–base superalloys, MC carbides have been found to redissolve or convert into other types of carbides (M₆C, M₇C₃, or M₂₃C₆) during long-term aging, the effect still exists [21,25]. Therefore, finding how to improve the intermediate temperature plasticity of the superalloy containing MC is an urgent problem to be solved.

In the present study, a Co–Al–W–base superalloy with optimized components was prepared. The superalloy consisted of γ–Co, γ’–Co₃(Al, W), and MC phases. No other phases were observed in this superalloy, which reduces the effect of other phases on plasticity, and makes the role of MC carbides more prominent. It provides a favorable condition for investigating the effect of MC carbide on the plasticity of the alloy. All three phases have face-centered cubic structures. The γ’–Co₃(Al, W) phase has a coherent relationship with the matrix (γ–Co), while MC carbide has an incoherent relationship with the matrix. This will cause the interface between the matrix and MC carbides to generate microcracks. It is a common phenomenon that the γ/MC interface generates microcracks in Nickel-base superalloys, which can seriously affect the plasticity of the superalloys. However, the fracture mode of MC carbide and the effect of MC carbide on the plasticity in a polycrystalline Co–Al–W–base superalloy is not clear. Therefore, in the current paper, the effect of MC carbide on the plasticity of the superalloy was investigated by using a Co–Al–W–base superalloy that consisted of γ, γ’, and MC phases in order to improve the intermediate temperature plasticity of the superalloy. Both microstructure characterization and interrupted tensile tests were performed, in order to investigate the mechanical behavior of MC carbide in the deformation process of the superalloy at 800 °C.

2. Materials and Methods

The main chemical compositions (at%) of the Co–Al–W–base superalloy used for present study were 26.00Ni, 9.00Al, 1.50W, 2.30Ti, 2.30Ta, 2.00Mo, 8.00Cr, 0.12B, 0.10Zr, 0.50C, and Co balance. The superalloy was prepared by the double vacuum induction melting method, then cast into bars with a diameter of 13 mm. The microstructure of the superalloy was analyzed by X-ray diffraction (XRD, RIGAKU SmartLab, Cu Kα radiation) analyses in the θ–2θ mode, as shown in Figure 1. Tensile specimens were machined into round bars with a diameter of 5 mm and a gauge length of 25 mm. Tensile tests were carried out on a Shimadzu testing system (AG–X plus 100 kN) with a resistance-heated furnace, and at the strain rate of 1 × 10⁻³ s⁻¹. The real-time temperatures of the tensile specimen were measured by three thermocouples directly contacted with the specimen in the furnace. The tensile tests were interrupted at the residual strain of 0%, 2%, and 5.0%, respectively, for quantitative microstructure analysis. Secondary electron photographs of the MC carbides and tensile fracture morphology were then examined by a scanning electron microscope (SEM, TESCAN MIRA3), and the chemical compositions of the MC carbides were analyzed via electron probe microanalysis (EPMA, JXA 8350F). Evaluation of the mean size and area fraction of MC carbides was performed by the Image (J image processing and analysis software). Twelve SEM photographs with an area of 1.4 mm × 1.0 mm were selected from the tensile specimens (homogeneous plastic deformation zone on the longitudinal section) in order to compute the mean size and area fraction.

3. Results

3.1. MC Carbide Characteristics

Figure 2a shows the morphologies of MC carbides in Co–Al–W–base superalloy. The MC carbides both exist inside the grains and on the grain boundaries (Figure 2b), exhibiting block (Figure 2c), dot-linked (Figure 2d), or skeleton (Figure 2e) shape. The morphology of MC carbide is depended on the solidification conditions [27]. In a quasi-rapidly solidified state, the MC carbide core grows in dendrite-like patterns [28,29]. Due to the various initial nucleation times of MC carbides [30], the carbide cores transform from octahedrons to block, dot-linked, and skeleton types. The area fractions of MC carbides inside the grain and on the grain boundaries are summarized in Figure 2f, exhibiting a distribution area ratio of approximately 9:1. The skeleton-type carbide has the maximum area fraction in the present superalloy, while the dot-linked-type carbide has the minimum area fraction, and the block-type carbide has nearly half the area fraction of that of the skeleton-type carbide. The difference in the area fraction is due to the various initial nucleation times of MC carbides.

Each type of MC carbide has a remarkable variation in size (Figure 2g). The mean length and width of the block-type carbide are 1.98 ± 0.2 μm and 1.72 ± 0.2 μm, respectively, which has the minimum mean length. The mean length and width of the dot-linked-type carbide are 11.58 ± 0.2 μm and 1.27 ± 0.2 μm, respectively, which has the maximum aspect ratio. The mean length and width of the skeleton-type carbide are 38.84 ± 0.2 μm and 22.94 ± 0.2 μm, respectively, which is approximately 20 times the mean length and over 13 times the mean width of the block-type carbide. The MC carbide in the present superalloy is rich in Ta, Ti, W, and Zr (Figure 3), and the Zr promotes the formation of blocky MC carbides [31,32] and enhances its area fraction.

3.2. Effect of MC Carbides on Deformation Behavior of the Superalloy at 800 °C

The stress-strain curve and the fractography of the Co–Al–W–base superalloy performed at 800 °C are shown in Figure 4. It can be seen that the elongation at fracture is 10.2% (Figure 4a). The superalloy shows good plasticity with a ductile fracture, compared to the reported data (an elongation at fracture of 2–5%) [18,33]. On the fracture surface exists many dimples that possess transgranular fracture characteristics (Figure 4b). It has two abnormal-sized dimples (marked with 1 and 2) in the crack initiation area (marked with a red line), as shown in Figure 4c. Two huge skeleton-type MC carbides are observed embedded in these dimples (Figure 4d,e). This pattern of the fracture indicates that the crack initiation is associated with MC carbides.

In order to clarify the role of MC carbides in the deformation behavior of the superalloy, the interrupted tensile tests at 800 °C were performed at the residual strain of 0% (I), 0.2% (II), and 5.0% (III), respectively, as illustrated in Figure 4a. The morphologies of the MC carbides (uniform deformation zone) in these interrupted tensile specimens were then compared to that of the fractured specimen (10.2% (IV)). Figure 5 shows the micrographs of MC carbides in the specimens at various residual strain, and microcracks are pointed out with different colored arrows. The microcrack size is associated with MC carbide size, which means that the skeleton-type MC carbide can generate more and larger microcracks. The MC carbide breakers (Since the MC carbide is a kind of brittle phase, it is easily broken with certain deformation, so it was called the MC carbide breaker) can be observed in the specimens at the residual strain of 0.2%, 5.0%, and 10.2% (Figure 5b–d), respectively, but no MC/γ interface debonding can be found. This indicates that microcracks are preferentially initiated in MC carbides after yielding. The three different types of carbides generate microcracks with different sizes, resulting in different size dimples on the fracture surface. In addition, a phenomenon where several microcracks are generated in the same MC carbide (including block, dot-linked, and skeleton types of MC carbides) was found (Figure 5d). Due to an increase in strain, more severe stress concentrations occur around the carbides, resulting in this phenomenon being observed in more MC carbides.

The numbers of MC carbide breakers in various deformation stages were examined (Figure 6). It is found that the increasing tendency of cracks per unit area is similar to that of MC carbide breakers during plastic deformation, and the increasing rate of the number can be divided into three stages. At stage I (0–0.2%), the specimen undergoes elastic-plastic deformation, causing a small number of MC carbides to break; at stage II (0.2–5.0%), the specimen experiences uniform plastic deformation, and the number of MC carbide breakers significantly increases; and at stage III (5.0–10.2%), the uniform plastic deformation decreases and the non-uniform plastic deformation increases, resulting in a lower increasing rate of MC carbide breakers than that at stage II. The three stages indicate that MC carbide fracture runs through the entire deformation process. These microcracks are all generated from MC carbides. It suggests that MC carbide breakers decide the number of microcracks, which also means that MC carbide breakers can control the high-temperature plasticity of the superalloy at 800 °C. The increasing tendency of skeleton-type MC carbides is significantly higher than the other two types of MC carbides. The breakage of skeleton-type carbides reduces the stress concentration of dot-linked and block carbides, thereby reducing their numbers (Figure 6b).

3.3. High-Temperature Ductility Evaluation of the Superalloy by MC Carbides

It is found that the fracture paths of MC carbides are different comparing the MC carbide fracture characteristics of the longitudinal section (Figure 7a–c) and the fracture surface (Figure 7d). The difference is mainly dependent on the angle between the length direction of an MC carbide and the tensile loading direction. When the angle is greater than 45°, an MC carbide tends to fracture along its length direction. However, an MC carbide tends to fracture along its width direction when the angle is smaller than 45°. It indicates that the angle between the length direction and the loading direction affects the fracture path of MC carbide. Obviously, the microcrack sizes of the two fracture paths are different. In other words, the difference in fracture path will significantly affect the intermediate temperature plasticity of the superalloy at 800 °C.

It can be imaged that whole MC carbides fracture along their length direction or width direction in the plastic deformation process of the superalloy at 800 °C. When whole MC carbides fracture along their width direction, the sum size of microcracks can be at the minimum value. However, once the whole MC carbides fracture along their length direction, the sum size of microcracks would be at the maximum value. The sizes of microcracks formed in the two conditions are significantly different. The minimum value of microcrack size is an industrial goal. However, the maximum value of microcracks is the critical value for evaluating the reasonable C content, aiming to generate the appropriate number of MC carbides. Therefore, the condition of the maximum value is considered.

When the MC carbide fractures along its length direction, the microcrack size is equal to the length of the MC carbide (hereinafter referred to as MC carbide size). This indicates that the largest microcrack size is equal to the MC carbide size. Therefore, the sum of microcrack sizes is equal to the sum of MC carbide sizes, which means the microcrack size can be replaced by the MC carbide size. Thus, MC carbide size can be used to describe the relationship between breaking stress and internal crack. In tensile testing, the crack size has a certain relationship with the fracture strength of the sample, as shown below [34,35]:

$${\sigma }_c=\frac{K}{\sqrt{a_i}}$$

(1)

where σ_c is the breaking stress and a_i is the internal crack size.

The sum of MC carbide sizes is used to replace the internal crack size. Then, Equation (1) can be expressed as follows:

$${\sigma }_c=\frac{K}{\sqrt{\sum l_{MC}}}$$

(2)

where ∑l_MC is the sum of MC carbide breaker sizes. However, during the plastic deformation of the specimen, the number of MC carbide breakers increases with the strain. According to the stress concentration factor equation [36]:

$${\sigma }_{max}=2\sigma {\left(\frac{a}{\rho }\right)}^{\frac{1}{2}}$$

(3)

where σ_max is the maximum stress of the crack tip, σ is the nominal tensile stress, a is half of the crack size, and ρ is the crack curvature radius. The values (under different residual strains) of σ and a are extracted from the stress-strain curve (Figure 4a) and the statistical results (Figure 6b). Thus, the stress concentration coefficient under different residual strains is obtained as follows: ${\frac{a}{\rho }}_{0.2\%}<{\frac{a}{\rho }}_{5.0\%}<{\frac{a}{\rho }}_{10.2\%}$. With the increasing strain, the ratio of MC carbide breakers that corresponds to skeleton-type carbides also increases (Figure 6b). This suggests that the breakage of skeleton carbides reduces the stress concentration of dot-linked and block carbides, thereby reducing their numbers. Hence, the crack curvature radius can be added to Equation (2) to give the following:

$${\sigma }_c=\frac{K\sqrt{\rho }}{\sqrt{\sum l_{MC}}}$$

(4)

The value of ∑l_MC can be obtained by measuring the sum of MC carbide breaker sizes at the residual strain of 10.2%, and the values of K and ρ can be obtained by substituting the breaking stress and the value of ∑l_MC into Equation (4). Thus, the relationship between breaking stress and MC carbide size can be obtained. The total residual strain is equal to the sum of residual strains at the various plastic deformation stages. Importantly, the residual strain at the breaking stress contains the elastic-plastic deformation stage and the work hardening stage. Therefore, the equation for the total residual strain can be shown below:

$${\epsilon }_r={\epsilon }_c+{\epsilon }_n$$

(5)

where ε_r is the total residual strain, ε_c is the residual strain of the breaking stress, and ε_n is the residual strain at the necking stage. In this study, it is found that the value of ε_n is approximately 1.3%, taking C as a constant. Then, Equation (5) can be expressed as the following:

$${\epsilon }_r={\epsilon }_c+C$$

(6)

Here, only ε_c is needed to be determined. The key is to establish the relationship between stress and strain. Additionally, the stress-strain curve shows this relationship. Interestingly, the stress-strain curve in the work-hardening stage is very flat. It can be approximated as a straight line. Then drawing two intersecting lines, A and B, at the breaking stress point shows an approximately linear relationship that exists between σ_c and ε_c, as shown in Figure 8.

In Figure 8, the work hardening stage can be fitted as a straight line (replaced by line A). Line B is the mathematical relationship between σ_c and ε_c. Line A and line B intersect at the point (ε_c, σ_c), where ε_c is the engineering strain. Therefore, according to the linear strain-hardening model [37], the equations of lines A and B can be derived as follows:

$$\sigma =K_w\epsilon +{\sigma }_0$$

(7)

$$\sigma =K_e\left(\epsilon -{\epsilon }_c\right)$$

(8)

where K_w is the slope of line A, σ_o is the y-intercept of line A, K_e is the slope of line B, and ε_c is the x-intercept of line B. Equation (8) shows a specific relationship between σ and ε_c. Substituting Equation (7) into (8) can eliminate the variable ε. Here, the relationship between σ and ε_c is well established. σ_c is just substituted into the variable σ. Thus, when the variable σ = σ_c, the equation for the relationship between σ_c and ε_c is as follows:

$${\epsilon }_c=\frac{\left(K_e-K_w\right){\sigma }_c-K_e{\sigma }_0}{K_w{K}_e}$$

(9)

Equation (9) shows the relationship between σ_c and ε_c. The relationship between σ_c and l_MC was obtained in Equation (4). By combining these two relations and eliminating σ_c, the relationship between ε_c and l_MC can be obtained. However, ε_c is not the sum strain ε_r, which needs plus C. Then, by substituting Equations (4) and (6) into (9), we can derive the equation that relates ε_r and l_MC:

$${\epsilon }_r=\frac{\left(K_e-K_w\right)K\sqrt{\rho }}{K_e{K}_w\sqrt{\sum l_{MC}}}-\frac{{\sigma }_0}{K_w}+C$$

(10)

Equation (10) shows the relationship between the total residual strain and MC carbide size, which indicates that the residual strain is inversely proportional to the MC carbide size. Meanwhile, Equation (10) clearly indicates the effect of MC carbide on the plasticity of Co–Al–W–base superalloy at 800 °C. The plasticity of the superalloy is mainly dependent on MC carbides. The values of K_w, K_e, σ_c, and ε_c can be obtained from the stress-strain curve. Additionally, the constant C is 1.3%. The value of ∑l_MC needs to count the number of all fractured MC carbides in the cross-section perpendicular to the loading direction. Meanwhile, three types of MC carbides need to be calculated separately and summed. Then, the ∑l_MC can be obtained. Substituting the value of ∑l_MC into Equation (4), the values of K and ρ can be obtained. Hence, Equation (10) can be used to evaluate the effect of MC carbide size on the residual strain. The elongation at fracture of a ductile material is larger than 5%. It means that the sum microcrack size obtained at 5% strain condition is the sum critical size of MC carbides. Substituting ${\epsilon }_r=5\%$ into Equation (10), the sum size of MC carbide breakers can be obtained. Then, dividing this value by the amount of MC carbide breakers can obtain the mean critical size of an MC carbide, $l_{MC}=47 \mathsf{\mu }\mathrm{m}$. The mean critical size of an MC carbide ignores the morphology of carbides. It is worth noting that the condition considered in Equation (10) is that all MC carbides fracture along the length direction, and the microcrack size is maximized, which is the maximum length of MC carbides. In polycrystalline Co–Al–W–base superalloys, the orientations of the grains are varied, which means that the angle between the MC carbide in each grain and the loading direction is different. This causes all MC carbides to fracture not along the length direction, but along the width direction, as shown in Figure 5b–d. In this case, the length of microcracks generated by MC carbide breakers during plastic deformation is not equal to the length of the MC carbide. However, the actual fracture of MC carbides is complicated, and the microcrack size is difficult to measure. Therefore, Equation (10) assumes that all MC carbides fracture along the length direction, and the microcrack size is replaced with the MC carbide size. This method is not only simple to calculate, but the calculated critical value can also make the C content conform to a specific safe range, which is more convenient for controlling the content C. It can not only give full play to the high-temperature strengthening effect of MC carbides, but also reduce the damage caused by excessive MC carbides to the plasticity of the superalloy at 800 °C.

MC carbide has a high melting point and high-temperature microstructure stability, which is very suitable for application in Co–Al–W–base superalloys. Due to the brittleness of MC carbide, it has been very sensitive to its application. After generations of scholars’ research, the brittleness of MC carbide has still not been solved, resulting in the application of MC carbide being severely limited. In the industry, the morphology of MC carbides is controlled by content C and solidification rate. The ideal morphology of MC carbide is spherical in small size. In this study, two fracture paths of MC carbides were found. It opens up the possibility of large-scale MC carbide applications. Under uniaxial loading, making sure the angle between all MC carbides and loading direction is less than 45°, the microcracks generated by MC carbide breakers can be significantly reduced. This has an important effect on the intermediate temperature plasticity of Co–Al–W–base alloys at 800 °C.

4. Conclusions

The relationship between MC carbide and intermediate temperature plasticity of a γ/γ’/MC Co–Al–W–base superalloy was investigated at 800 °C, the mechanical behavior of MC carbide in the whole deformation process was explored, and the relationship between MC carbide size and the residual strain was established. The main conclusions are as follows:

(1)

A Co–Al–W–base superalloy that consisted of γ–Co, γ’–Co₃(Al, W), and MC carbide was developed;

(2)

MC carbide breakers are the source of microcracks, and directly affect the plasticity of Co–Al–W–base superalloy at 800 °C;

(3)

MC carbide breakers appear after yielding. It is mainly dependent on the morphology and the size of the MC carbides. Fractures along the length direction and the width direction are two primary fracture paths of MC carbides;

(4)

The intermediate temperature plasticity of the Co–Al–W–base superalloy can be guaranteed by controlling the critical size of MC carbides.