Temperature Effects on the Deformation Mechanisms in a Ni-Co-Based Superalloys

Abstract

The tensile properties of a Ni-Co-based superalloy were investigated from room temperature to 900 °C. From 25 to 650 °C, the yield strength and tensile strength of the alloy decreased slightly, while the elongation decreased sharply. From 760 °C to 900 °C, the yield strength and tensile strength were greatly reduced, while the elongation also had a low value. With the increase in temperature, the deformation mechanism transformed from anti-phase boundary shearing to stacking fault shearing, and then from deformation twinning to Orowan bypassing, respectively. Deformation twins were generated in the deformed alloy with high-density stacking faults and they can contribute to the high strength. The alloy in this study has good mechanical properties and hot working characteristics below 760 °C and can be used as a turbine disk, turbine blade, combustion chamber, and other aircraft structural parts.

1. Introduction

Nickel-based superalloys are widely used in aircraft engines because of their excellent strength and creep resistance at high temperatures [1,2]. The outstanding strength of these alloys mainly stems from solid solution strengthening elements and high volume fractions of γ′ precipitates [3,4]. The superalloy turbine discs are used at temperatures of 650~700 °C and sometimes reach 815 °C, according to the actual service conditions [5]. The temperature-bearing capacity and mechanical properties of turbine disks determine the thermal efficiency and service life of the engine system. Conventional strengthening mechanisms, such as solid-solution strengthening, work hardening, grain boundaries strengthening, and precipitation hardening, have previously been used to improve the properties of superalloys [6]. However, the above-mentioned defects are not thermally stable at high temperatures. Lu et al. reported that nanoscale coherent twin boundaries (TB) can effectively enhance nano and ultra-fine grained materials without adversely affecting the ductility [7,8]. The addition of Co will reduce the stacking fault energy (SFE) of superalloys [9]. The decrease of SFE can change the deformation mechanism of the alloys, which may lead to twin boundary strengthening [6]. Therefore, it is necessary to study the deformation mechanism and mechanical properties of Ni-Co-based superalloys from room temperature to high temperatures.

The intermediate and high-temperature deformation mechanisms have been extensively investigated in the past decades [10,11,12,13,14,15,16,17]. The research results of Yuan et al. [18] show that the addition of the Co element reduces the stacking fault (SF) energy of the Ni-based superalloy and influences its deformation mechanism. The dominant deformation mechanism of superalloys is that a pair of dislocations shear γ′ precipitates at low temperatures, while SF is formed by the movement of partial dislocations at intermediate temperatures. Xu et al. [19] studied the variation trend of SF energy and SF density with temperatures, showing that the SF density increased with the increase in temperatures. Yuan et al. [20] reported that the mechanisms of antiphase boundary (APB) cutting and SF cutting both appeared in the alloy. With the increase in temperature, deformation twins are easier to form than APB.

In this study, the temperature dependence of the tensile behavior of a Ni-Co-based superalloy was investigated. On the basis of the microstructure analysis, the transformation of the deformation mechanism (with temperature) and the intermediate temperature brittleness were discussed.

2. Experimental Procedures

The studied Ni-Co-based superalloy has the following chemical compositions: 25.0 Co, 14.0 Cr, 2.8 Mo, 1.2 W, 8.0 (Al, Ti, Nb), 0.02 C, 0.02 B, 0.03 Zr, and balanced Ni (wt%). The alloy was prepared by vacuum induction melting and forging. The samples were cut by electric spark cutting and homogenized at 1100 °C to eliminate the segregation. Aging treatments were carried out as follows: 650 °C/24 h/air cooling + 760 °C/16 h/air cooling. The equipment used for the heat treatment was a KSL-1400X-A1 high-temperature box furnace. Specimens with a gauge length of 25 mm and a diameter of 12 mm were machined out for tensile tests. The tensile testing equipment was an electronic universal material testing machine produced by INSTRON Company. The tensile tests were conducted with a strain rate of 3 × 10⁻⁴ s⁻¹ at various temperatures, i.e., 25 °C, 650 °C, 760 °C, 850 °C, and 900 °C, respectively. A temperature fluctuation not exceeding ±2 °C was maintained over the gauge length. At least three identical specimens were tested at each temperature. After the tensile tests, the fracture surfaces of the specimens were investigated using a Zeiss Gemini 500 field emission scanning electron microscope (FESEM) operated at 15 kV. Slices were cut perpendicular to the tensile direction from the deformed samples and mechanically thinned down to 50 μm, and then thinned by ion-beam milling. The microstructures of the deformed samples were investigated using a Talos F200X transmission electron microscope (TEM) operated at 200 kV. It was also equipped with an X-ray energy dispersion analyzer for chemical composition analysis.

3. Results and Discussion

3.1. Mechanical Properties Analysis

The tensile properties of the superalloys at different temperatures are shown in Figure 1. The yield strength (YS) and ultimate tensile strength (UTS) of the superalloy had similar variation trends with the temperature. The YS and UTS decreased slightly when the temperature rose to 650 °C and then decreased rapidly above this temperature. The elongation decreased rapidly with temperature, from room temperature to 650 °C, and then maintained a low value above 650 °C, demonstrating an intermediate temperature brittleness (ITB) behavior.

The γ′ precipitation has a unique feature in that its strength increases with the increase in temperature [21,22,23]. The superalloy studied in this work contained a high-volume fraction of γ′ precipitates. The decrease in strength of the γ matrix was compensated by the increased strength of γ′ precipitates with the increase in temperature [24], due to the thermally-activated cross-slip of dislocations from {111} planes to {110} planes [21,25].

The YS of the alloy depended greatly on the testing temperature. According to References [25,26], the yielding behavior at a high temperature is a thermally activated process. Therefore, the Arrhenius type of relationship may be suitable for characterizing the plastic flow behavior of this alloy, i.e.,

σ_y/E = A{exp[Q/(RT)]}

(1)

where σ_y is the yield strength, E is the elastic modulus, Q is the activation energy of the deformation process, R is the gas constant (8.314 J mol⁻¹ K⁻¹), T is the temperature, and A is the material constant. Equation (1) was applied to several theoretical and experimental studies of superalloys [27,28]. A plot of the ratio of normalized values of YS by the elastic modulus vs. the reciprocal of temperature is shown in Figure 2. It is obvious that there are three different stages as functions of temperatures. At high temperatures (above 760 °C), the curve had a positive slope, at intermediate temperatures, the curve exhibited a negative slope, and at low temperatures (below 650 °C), the curve dropped slowly. This indicates that the rate-controlling processes for the deformation mechanisms had great changes. The main strengthening mechanism of the alloy was the γ′ precipitation, which could effectively block dislocation movement. So, at low and intermediate temperatures, the high applied stress could force dislocations to cut into γ′ precipitates. With the increase in temperature, the applied stress value is less than the critical stress required for dislocation cutting γ′ precipitates, and then the enhanced thermal activation causes the dislocation climbing mechanism [29].

3.2. Mechanical Property Analysis

3.2.1. Initial Microstructure

A typical microstructure of the alloy after the heat treatment is shown in Figure 3. One can see that the γ′ precipitates are uniformly distributed in the alloy. According to the forming sequence, the γ′ precipitates in the alloy can be divided into primary, secondary, and tertiary γ′ precipitates. As shown in Figure 3a,c,e, the primary γ′ precipitates tend to appear at grain boundaries with larger sizes, while the secondary γ′ precipitates are located within moderately sized grains. In addition, a large number of nanoscale tertiary γ′ precipitations can be seen inside the grains. The γ′ precipitates are L1₂ structures confirmed by the SAED patterns (Figure 3b,d,f) and are coherent with the γ matrix.

By calibrating the super diffraction spots and the matrix diffraction spots, the orientation relationships between the precipitates and the matrix can be obtained as follows (m represents matrix, p represents γ′ precipitates):

$$[110]\mathrm{m}//[110]\mathrm{p}\hspace{1em}\hspace{1em}\hspace{1em}(002)\mathrm{m}//(002)\mathrm{p}\hspace{1em}\hspace{1em}\hspace{1em}(111{)}_{\mathrm{m}}//(112{)}_{\mathrm{p}}$$

The element content in the alloy was analyzed by TEM equipped with an X-ray energy dispersive analyzer; the results are shown in

Table 1. Alloy element distribution (wt %).

	Ni	Co	Cr	Ti	Al	Mo	W
Primary γ′	66.1	17.2	3.0	8.7	3.6	0.7	0.7
Secondary γ′	63.0	18.7	7.0	6.4	2.4	1.5	1.0
γ matrix	50.8	24.8	13.6	4.8	1.9	2.8	1.3
Cprimary γ′/Cγ	1.301	0.694	0.221	1.813	1.895	0.250	0.538
Csecondary γ′/Cγ	1.240	0.754	0.515	1.333	1.263	0.536	0.769

. The contents of C, B, and other light elements are difficult to be measured by the energy spectrometer, so only some elements are listed in Table 1. It can be seen that Co, Cr, Mo, and W are enriched in the γ matrix and play the role of solid solution strengthening, while Ti, Al, and Ni are enriched in the γ′ precipitates. These three elements mainly form (Ni,Co)₃(Al,Ti) during the heat treatment process, and the γ′ precipitates can play a role in precipitation strengthening.

3.2.2. Microstructure Analysis of the Deformed Samples

Figure 4 shows the deformed microstructure of the room temperature-tested sample observed by TEM. Apparently, there are some dislocations in both the γ matrix and the γ′ precipitation at room temperature. Most of the dislocations move alone, some in pairs. Due to the smaller size of the tertiary γ′ precipitate, its density is higher [14]. When a partial dislocation is cut into the γ′ precipitation, high-energy APBs are generated due to the a/2<110> displacement of the superlattice. So dislocations occur through the γ′ precipitation in closely spaced pairs to minimize the APB [14].

The TEM observations of this alloy tested at different temperatures are shown in Figure 5, Figure 6 and Figure 7. Figure 5 presents the deformed microstructures after a tensile test at 650 °C. When the testing temperature is 650 °C, the dislocation density becomes high and a number of SFs are produced both inside the γ′ precipitates and the γ matrix. Moreover, the formation of dislocation entanglement in the γ matrix indicates that the deformation occurs mainly in the γ matrix. The SFs were observed in two different directions, which formed structures similar to Lomer–Cottrell locks. As indicated by the arrows in Figure 5b, these SFs on different planes intersect and interact with each other. The observed SFs have been investigated in some alloys [29]. Chu et al. [30] believed that the SFE of DZ951 alloy decreased with temperature from 20 °C to 760 °C. The transition temperature from APB shearing to the SF shearing mechanism decreases due to the combination of higher APB energy and lower SFE [20]. Therefore, SFs will appear as the temperature increases.

When the test temperature is 760 °C, it can be seen that the dislocation density in the γ matrix further increases, and they become entangled with each other. In addition, when the test temperature is increased, a large number of deformation twins are generated in the alloy. A study of FCC crystals showed that SFs and deformation twins were sequentially generated in the alloy with increasing plastic strains and deformation temperatures [31]. Figure 6d shows a high-resolution transmission electron microscope (HRTEM) image of the superalloy. The twin boundaries are indicated by green lines. The thickness of the twinned lamella is 16 layers of atoms.

When the test temperature was 850 °C, it can be seen that more dislocations were formed in the γ matrix and tangled with each other (Figure 7a,b), but it is difficult to find any deformation twins, and the stacking faults parallel to each other are clearly visible after the tensile test at 850 °C. When the tensile temperature was increased from 760 °C to 850 °C, there were many dislocation loops, as seen in Figure 7b. Figure 7c,d show that the deformation mechanism mainly involves the dislocation bypassing the γ′ precipitation through the Orowan process, and the dislocations pile up around the γ′ precipitation at high temperatures. A large number of dislocations surrounded the γ′ precipitates, while only a few dislocations were located in the γ′ precipitates. In contrast to the deformation at room temperature, the deformation was dominated by Orowan bypassing at 900 °C. When the test temperature increased, the applied stress was less than the critical stress required for dislocation cutting the γ′ precipitate, then the enhanced thermal activation caused the dislocation climbing mechanism. It can be seen that as the temperature increases, the deformation activation energy decreases [29].

3.2.3. Fracture Surface Analysis

The morphologies of the fracture surfaces of the tensile specimens tested at different temperatures are shown in Figure 8. Clearly, the deformation is homogeneous, and the fracture has a large number of dimples. It can be seen that the tear-off dimples are elongated, so the plasticity is good, as shown in Figure 8a,b. This is consistent with the results obtained for higher elongation after room temperature tensile testing. When the testing temperature increased to 650 °C, the fracture morphology changed, as shown in Figure 8c,d. Clearly, the main features of the specimen tested at 650 °C were the decohesion combined with some dimples, as shown in Figure 8c. Although there were some dimples, the tear-off dimples were not elongated, so the deformation was limited. Further observations showed that there were many secondary cracks and debonding phenomena along the γ′ precipitate/γ matrix interfaces on the fracture surface. Secondary cracks readily propagated along rigid precipitates and the matrix, which is detrimental to ductility [32,33,34,35]. Therefore, it is easy to understand why the elongation of the alloy is lower at 650 °C. The fracture of the alloy tested at 760 °C is shown in Figure 8e,f. It can be seen that the sample exhibited a heavier brittle feature, with cleavage fracture being the main feature, as shown in Figure 8e. The secondary cracks propagated along the grain boundary. Moreover, small dimples and precipitates were distributed on the cleavage facets. The deformation was not homogeneous, so the elongation was low. The morphologies of the fracture surfaces of the tensile specimens tested at 850 °C and 900 °C are shown in Figure 8g,h, respectively. The fracture pattern was mainly in the intergranular mode. Decohesion along the γ′ precipitates/γ matrix interfaces existed on the fracture surface.

4. Discussion about the Mechanism of Intermediate-Temperature Brittleness

Although the elongation is high at room temperature, the elongation is greatly decreased when the test temperature is higher than 650 °C. The intermediate temperature brittleness was observed in several superalloys with high volume fractions of γ′ precipitates. There are a number of reasons for this behavior, and the most important factors for superalloys are [36]: (1) carbide particles are rich in trace elements or oxygen, resulting in the embrittlement of grain boundaries; (2) changes in the deformation mechanisms lead to strain localization; (3) the instability of γ′ precipitates that are exposed to high temperatures. Clearly, no carbide particles are found in the present superalloy, so intermediate-temperature brittleness is not caused by carbide particles. Chu et al. [30] showed that deformation was mainly dominated by pairs of a/2<110> dislocation shearing γ′ precipitation at low temperatures, and it was mainly dominated by Orowan bypassing at high temperatures. They found the intermediate temperature brittleness at 760 °C was due to strain localization mainly caused by SFs. Kim et al. [13] suggested that the generation of deformation twins has an important effect on the intermediate temperature brittleness. At the same time, they showed that deformation twins are easy to form in superalloys with low SFE.

In general, when a new deformation mechanism appears, the ductility decreases simultaneously [37]. The study of the deformed microstructure provides information for analyzing possible fracture mechanisms. The dominant deformation mechanisms is that the γ′ precipitates is cut by dislocations at low temperatures (below 650 °C). Orowan bypass is the main deformation mechanism at high temperatures (above 900 °C). At 760 °C, a high density of dislocations was formed in the γ matrix and many of them were tangled. The operation of a multi-slip system will generate the delivery and reaction of dislocations, thereby creating more barriers for cross-slips. Most importantly, there are a large number of SFs and deformation twins, which can also create barriers for cross-slips. The deformation is very inhomogeneous, and the local strain reaches the maximum simultaneously, resulting in intermediate temperature brittleness at 760 °C. Similarly, there are a number of SFs both at 650 °C and 850 °C. So the intermediate temperature brittleness is mainly caused by the strain localization caused by SFs or deformation twinning.

5. Conclusions

Through the above comprehensive analysis of the microstructural changes of the Ni-Co-based superalloy (after tensile testing at different temperatures), the following conclusions can be drawn:

(1)

From room temperature to 650 °C, the YS and TS of the alloy decrease slightly, while the elongation decreases sharply. From 760 °C to 900 °C, YS and TS are greatly reduced; the elongation also has a low value. There is an intermediate temperature brittleness.

(2)

After heat treatment, the alloy has three different γ′ precipitate sizes, which play important roles in precipitation strengthening. With the increase in temperature, the deformation mechanism changes from anti-phase boundary shearing to stacking fault shearing, and then deformation twins are generated in the deformed alloy with high-density stacking faults. The alloy has good mechanical properties and hot working characteristics below 760 °C and can be used as a turbine disk, turbine blade, combustion chamber, and other aircraft structural parts.

(3)

From the analysis of the results, the mechanical properties of the alloy decrease sharply from 850 °C to 900 °C, and the deformation mechanism also changes to Orowan bypass.