Hot Deformation Behavior and Dynamic Recrystallization Mechanism of GH3230 Superalloy

Abstract

An isothermal hot compression test of GH3230 was carried out under deformation conditions with deformation temperatures ranging from 1020 to 1110 °C and strain rates ranging from 1 to 0.001 s⁻¹. On this basis, the corresponding constitutive equation of the alloy was established. $\dot{\epsilon }=exp\left(36.123\right){\left[sinh\left(0.00587\sigma \right)\right]}^{4.7946}exp\left[-451.507/\left(RT\right)\right]$. At the same time, a power dissipation diagram and thermal processing diagram were created. The peak value η can reach 0.36, and the optimum hot working parameter window of the GH3230 superalloy is 1020~1110 °C/0.1~0.001 s⁻¹. The microstructure evolution of the alloy under different conditions was studied by EBSD. With an increase in deformation temperature and a decrease in strain rate, the grain size significantly improved; the average grain size of the GH3230 alloy increased from 16.86 to 35.06 μm, and the degree of recrystallization of the alloy also improved. The maximum recrystallization volume fraction is 75.2%. At low temperature and high strain rate, the recrystallization mechanism of the microstructure is mainly CDRX, and DDRX is the auxiliary mechanism. At high temperature and low strain rate, the main corresponding recrystallization mechanism gradually transforms into DDRX.

1. Introduction

There are two kinds of softening phenomena in nickel-based superalloys during high-temperature processing, namely dynamic recovery and dynamic recrystallization [1,2,3]. The whole hot working process of the alloy involves work hardening, during which these two softening phenomena occur together, leading to changes in the microstructure and properties of the alloy [4,5]. Nickel-based superalloys are very typical low-stacking-fault-energy alloys. In the process of thermal processing, a large number of dislocations in the microstructure are decomposed, and these generated edge dislocations and screw dislocations are difficult to slip by climb or cross-slip [6,7]. With the continuous accumulation of dislocations, dislocation tangles and dislocation cells form. As deformation continues, dislocations accumulate to a certain threshold, and new grains are produced. This process is dynamic recrystallization, and the small new grains are called recrystallized grains [8]. The dynamic recrystallization process can not only change the microstructure of the alloy but also greatly improve its processing and performance [9].

Azarbarmas et al. [10] analyzed the deformation behavior of the Inconel 718 alloy by hot compression. The results show that when deformation is small, DDRX with bow-shaped nuclei at grain boundaries plays a leading role. With an increase in deformation, the effect of CDRX is enhanced. At the same time, when the deformation rate is relatively large and the temperature is low, CDRX is also promoted. Song et al. [11] studied and analyzed the hot deformation and recrystallization behavior of new nickel-based superalloys for ultra-supercritical applications at different processing temperatures (950~1150 °C) and strain rates (0.01~10 s⁻¹). It was revealed that the recrystallization mechanism of the alloy is dominated by discontinuous dynamic recrystallization. Zhang et al. [12] studied Inconel 718 by conducting an isothermal compression test. The smaller size and rod-like δ phase make the DRX fraction increase faster. Compared with the alloy with a rod-like δ phase, the alloy with a needle-like δ phase requires a larger strain to reach the whole DRX fraction. For studied alloys with a smaller δ phase size, finer grain size can be obtained after the deformation stage.

GH3230 is a nickel-based superalloy designed for corrosion resistance and high-temperature applications. The main chemical composition is Ni-Cr-W, which makes it a solid-solution-strengthened superalloy. GH3230 has high stability in the temperature range from 700 °C to 1100 °C. The GH3230 alloy is effective in the temperature range from 700 °C to 1100 °C and can still maintain high tensile strength (about 900 MPa at room temperature) and yield strength (about 500 MPa) at high temperatures (above 800 °C) [13,14]. Compared with other nickel-based alloys, the GH3230 alloy has higher strength, thermal fatigue and cold fatigue properties, and heat corrosion resistance. Because of its excellent performance in terms of thermal stability, fatigue resistance, and oxidation resistance, GH3230 is widely used in key components under long-term high-temperature conditions, especially in the fields of aero-engines and gas turbine stator blades [15]. The microstructure of the alloy determines its mechanical properties, and the strain rate of the material in each part of the engine is also different. Different strain rates and different temperatures also have different effects on the mechanical behavior of the superalloy. The hot deformation property of GH3230 makes it extremely sensitive to temperature. If the processing temperature is lower than 1000 °C, the plasticity of the material decreases significantly, and cracks are likely to occur [16]. However, temperatures that are too high (such as >1150 °C) may lead to excessive grain coarsening or local melting, affecting the final mechanical properties. Therefore, it is necessary to study the microstructure evolution of the GH3230 alloy [17].

Zhang et al. [18] established a corresponding constitutive model by stretching the CH3230 alloy at different strain rates. The results show that the activation energy of the alloy is more susceptible to strain rate. Hao et al. [19] studied the effects of different heat treatments on the mechanical properties of GH3230 prepared by L-PBF and found that hot isostatic pressing (HIP) and solid solution treatment (ST) can effectively reduce the elongation and anisotropy of mechanical properties. Through the post-treatment of the GH3230 alloy prepared by SLM, Yuan et al. [20] found that the ductility and plasticity of the alloy were improved after heat treatment at 1250 °C/30 min in the HIP state.

At present, there are few reports on the microstructure evolution of the GH3230 alloy during hot deformation. It is necessary to study the hot deformation behavior of the GH3230 alloy. In this paper, the hot deformation behavior of the alloy at a deformation temperature of 1020~1110 °C and strain rate of 1~0.001 s⁻¹ was studied by designing a compression test. Based on the experimental results, a constitutive model and hot processing map were established, and the optimum process parameters were determined. The effects of different deformation parameters on the microstructure evolution of the alloy were analyzed in detail by Electron Backscatter Diffraction (EBSD), and the DRX mechanisms of GH3230 under different deformation conditions were clarified. This study provides a practical reference for the rational formulation of the process and the use of the GH3230 superalloy.

2. Materials and Experimental Procedures

The chemical composition of the GH3230 superalloy is shown in

Table 1. Chemical composition of GH3230.

Ni	Cr	W	Mo	Al	Ti	Fe
Bal	24	15	3	0.5	0.1	3

(wt %). A hot-rolled bar was cut into several small cylindrical specimens with a size of 8 mm × 12 mm (sampling along the rolling direction of the bar to ensure that the height direction is consistent with the rolling direction), and a hot compression experiment was carried out on the testing machine.

The sample was heated to the deformation temperature at a speed of 10 °C/s, and the temperature was kept evenly distributed. Then the experiment was carried out according to the different strain rates, and the deformation degree of each sample was 50%. The deformation temperatures used in this experiment are 1020, 1050, 1080, and 1110 °C; the holding time is 10 min; and the strain rates are 1, 0.1, 0.01, and 0.001 s⁻¹. The load–stroke curve obtained from the compression experiment was recorded and converted into a true stress–true strain curve by using a formula. After hot compression was completed, the sample was quenched in water immediately so that the microstructure morphology could be preserved after high-temperature deformation. After that, the deformed samples were cut along the deformation direction, and the samples were ground with 240-, 400-, 800-, 1200-, 1500-, 2000-, 2500-, and 3000-mesh sandpaper. Each piece of sandpaper was rotated 90° until the surface was smooth and scratch-free. Finally, the surface was polished with a polishing machine. The grinding paste is a SiO₂ suspension and was polished until the mirror surface was clean. For electrochemical corrosion at room temperature, the corrosion solution is C₂H₂O₄: HClO₄: C₂H₅OH = 1:1:9. Corrosion lasted for 30 s, with a voltage of 20 V and a time of 30 s. The microstructure of the center of the deformed part was observed by EBSD characterization. EBSD data were analyzed using Channel 5.

3. Results and Discussion

3.1. Stress–Strain Curve

After thermal simulation compression deformation at different temperatures and strain rates, the true force–strain curve of nickel-based superalloy GH3230 was established and is shown in Figure 1. It can be seen from Figure 1 that at different temperatures and strain rates, the change in flow stress with strain is divided into three stages: When the material yields, the flow stress rapidly reaches the peak with an increase in strain, that is, work hardening. Then, under the combined effect of deformation temperature and strain rate, the flow stress decreases with an increase in strain; that is, softening occurs. Finally, as the strain continues to increase, the magnitude of the flow stress tends to be stable, the degree of stress softening gradually decreases, and the alloy enters a stable deformation stage [21]. The relationship between peak stress and deformation parameters is shown in Figure 2. It can be seen that the peak stress gradually decreases with an increase in temperature and a decrease in strain rate.

3.2. Constitutive Model

The flow stress of the GH3230 nickel-based superalloy is obviously affected by the deformation rate and deformation temperature, which is the same as in the plastic deformation process of other metals. Flow stress is mainly controlled by the thermal activation process. The Arrhenius hyperbolic sine constitutive equation is usually used to describe the relationship between stress and strain rate and temperature. In this paper, the process of determining the material constant of the alloy at a true strain of 0.6 is taken as an example.

$$\dot{\epsilon }=A{\left[\mathit{sinh}\left(\alpha \sigma \right)\right]}^{″}\mathit{exp}\left[-Q/\left(RT\right)\right]$$

(1)

Formula (1) can also be expressed as follows:

$$Z=A{\left[\mathit{sinh}\left(\alpha \sigma \right)\right]}^{″}=\epsilon \mathit{exp}\left[Q/\left(RT\right)\right]$$

(2)

Formula (1) can be obtained by Taylor series expansion:

When the flow stress is low (ασ < 0.8), Equation (1) can be simplified to a power function relation, which is obtained by taking the logarithms of both sides

$$\mathit{ln}\dot{\epsilon }=\mathit{ln}{A}_1+{ n}_1\mathit{ln}\sigma -Q/\mathit{RT}$$

(3)

When the flow stress is high (ασ > 1.2), Equation (1) can simplify the exponential function relationship, and the logarithm of both sides is obtained

$$ln\dot{\epsilon }=ln{A}_2+\beta \alpha -Q/RT$$

(4)

where σ is the flow stress (MPa); Q is the deformation activation energy (kJ/mol), which is related to the material. $\dot{\mathsf{\epsilon }}$ is the strain rate (s⁻¹); T is the absolute temperature (K); R is the molar gas constant; A, A₁, A₂, n, n₁, α, and β are temperature-independent constants. n is the stress index; α and β are stress adjustment factors; and the relationship between α, β, and n₁ satisfies the following:

$$\alpha =\beta /n_1$$

(5)

Each parameter in Formula (1) can be derived from Formula (1) to Formula (4)

$$ln\sigma =ln\dot{\epsilon }/n_1-ln{A}_1/n_1$$

(6)

$$\sigma =ln\dot{\epsilon }/\beta -ln{A}_2/\beta$$

(7)

According to the stress–strain curve, the coordinates of lnσ and lnε and σ and lnε were plotted, and the lines were drawn by the least squares method. The linear relationship as shown in Figure 3 can be obtained, and the value of constant n₁ is shown in Figure 3a (a is the average value of the slope of the line at different temperatures, n₁ = 6.4733), and the value of β is shown in Figure 3b (b is the average value of the slope of the line at different temperatures, β = 0.0379; thus α = β/n₁= 0.0058).

Assuming that the activation energy of deformation is independent of temperature, it can be obtained by Equation (1)

$$ln\dot{\epsilon }=lnA+nln\left[\mathit{sinh}\left(\alpha \sigma \right)\right]-Q/\left(RT\right)$$

(8)

The partial integral of temperature T is obtained by deforming

$$Q=R\left\{{\displaystyle \frac{\partial ln\epsilon }{\partial ln\left[sinh\left(\alpha \sigma \right)\right]}}\right\}T\left\{{\displaystyle \frac{\partial ln\left[\mathit{sinh}\left(\alpha \sigma \right)\right]}{\partial \left(1/T\right)}}\right\}\dot{\mathsf{\epsilon }}$$

(9)

The relation curve between α and ln[sinh(ασ)] − lnε and the relation curve for ln[sinh(ασ)] − 1/T obtained above are drawn, as shown in Figure 4a, where the average reciprocal slope of each line in Figure 4b is the value of the stress index n, n = 4.7946. Combined with the obtained n value, it is brought into Equation (9), and the average deformation activation energy Q = 451.507 kJ/mol under different deformation temperatures and strain rates can be obtained.

The deformation activation energy obtained is put into Equation (2)

$$Z=\epsilon \mathrm{e}\mathrm{x}\mathrm{p}\left[600656.3/\left(RT\right)\right]={A\left[\mathit{sinh}\left(\alpha \sigma \right)\right]}^n$$

(10)

Taking the logarithm of both sides, we obtain

$$lnZ=lnA+nln\left[\mathit{sinh}\left(\alpha \sigma \right)\right]$$

(11)

By introducing the 4 relative strain rates at 4 temperatures into the calculation of the above equation, 16 relative Z values can be obtained, and a scatter plot of lnZ and ln[sinh(ασ)] can be drawn, and the linear relationship diagram shown in Figure 5 can be obtained by fitting. The intercept of the line in Figure 5 is the value of lnA, which is 36.123. Figure 5 shows that within the range of experimental parameters, the linear relationship between lnZ and ln[sinh(ασ)] is in good agreement, indicating that the relationship among flow stress, strain rates, and deformation temperature can be expressed by Equation (1). Substituting the material constant and the calculated deformation activation energy Q = 451.507 kJ/mol, the stress index n = 4.7946, the stress level parameter α = 0.00587 mm²/N, and other parameters into Formulas (1)–(4), the constitutive equation of nickel-based superalloy GH3230 in the thermal simulation compression test at 1020~1110 °C can be obtained.

3.3. Processing Map

The deformation process of the alloy is approximately regarded as a nonlinear and irreversible thermodynamic process. Through the evaluation of the hot working performance of the alloy, it can not only reflect the characteristics of the microstructure’s deformation but also clearly divide the safe area and instability area of the deformation parameters during the processing of metal materials so as to clarify the processing window of the material. According to the dynamic material model (DMM) theory proposed by Prased [22], the alloy component can be regarded as an energy dissipation body during processing and forming. The total energy P input to the workpiece by the outside world is partially converted into heat energy during the deformation process. Dissipation is denoted by G; the other part is consumed by the organizational transformation of the workpiece itself, which is represented by the dissipative covariate J.

$$P=\sigma \dot{\epsilon }=G+J={\int }_0^{\dot{\epsilon }}\sigma d\dot{\epsilon }+{\int }_0^{\sigma }\dot{\epsilon }d\sigma$$

(12)

The ratio of G to J depends on the strain rate sensitivity factor m. Generally, for ordinary plastic deformation materials, the value of m is between 0 and 1. If m = 1, ideal linear dissipation occurs; that is, J_max = P/2. The ratio of J to J_max is closely related to the microstructure evolution of the material during the deformation process, so it is defined as the power dissipation factor, expressed by η. The calculation of the η value is as follows:

$${\displaystyle \frac{\partial J}{\partial G}}={\displaystyle \frac{\partial \left(ln\sigma \right)}{\partial \left(ln\dot{\epsilon }\right)}}=m$$

(13)

According to Formulas (15) and (16), the expression of η is deduced. After calculating the value of η, a two-dimensional equivalent curve of the power dissipation value η of the material at five temperatures and four strain rates is drawn, which is called the power dissipation diagram of the material during the deformation process, as shown in Figure 6a.

$$\eta ={\displaystyle \frac{J}{J_{max}}}=\left\{\begin{array}{c}1, J=J_{max}\\ {\displaystyle \frac{2m}{m+1}}, J<J_{max}\end{array}\right.$$

(14)

In DMM theory, Sivakesavam et al. [23] proposed a material flow instability criterion based on the principle of an irreversible thermodynamic extremum under severe plastic deformation:

$$\xi \left(\dot{\epsilon }\right)={\displaystyle \frac{\partial \mathit{lg}\left(\frac{m}{m+1}\right)}{\partial \mathit{lg}\dot{\epsilon }}}+m<0$$

(15)

The region of $\xi \left(\dot{\epsilon }\right)<0$ represents the rheological instability region of the material. When plastic deformation is carried out at the corresponding temperature and strain rate in this region, the instability phenomenon of the alloy will occur, as shown in Figure 6b.

The thermal processing map of the material is drawn by superimposing the power dissipation map with the rheological instability region of the material, as shown in Figure 7. The gray part is the material instability area. By observing the hot processing map, it can be found that the area with η > 0.3 is mainly concentrated at a deformation temperature of 1020 °C~1040 °C and a strain rate of 0.001 s⁻¹. The microstructure evolution mechanism of GH3230 can be determined by the power dissipation coefficient. In the process of high-temperature deformation, the microstructure evolution mechanism of the superalloy from the initial deformation to the end mainly includes dynamic recrystallization, dynamic recovery, and super-plasticity. The instability behaviors in the deformation process mainly include local flow, the formation of adiabatic shear bands, and so on [24]. Generally speaking, in the case when the alloy shows no instability behavior, the higher power dissipation value reflects the fact that the alloy is easy to process, and an excellent microstructure and properties can be obtained under this parameter condition [25]. The η values of other regions are all less than 0.3. At the same time, it can be found that the instability region of the material is mainly concentrated at a deformation temperature of 1020 °C~1040 °C and a strain rate of 0.01 s⁻¹~1 s⁻¹, and the distribution is the greatest when the strain rate is large. This region should be avoided as much as possible during processing.

3.4. Structural Evolution

Figure 8 shows an inverse pole figure (IPF) of the typical microstructure of GH3230 after isothermal compression under different conditions. The red region represents the orientation tending to the <001> direction, the blue part represents the orientation tending to the <111> direction, and the green part represents the orientation tending to the <101> direction. It can be seen from the IPF of the typical structure that the hot deformation process under most conditions is accompanied by dynamic recrystallization. At the same time, it can also be observed that there are obvious differences in the deformed microstructures obtained under different deformation conditions, and the average grain size changes significantly under different conditions. When the deformation temperature is consistent, the average grain size of GH3230 increases gradually with the slow strain rate. The most significant change is at 1080 °C. With the change in strain rate, the average grain size increases from 16.86 μm to 30.12 μm. When the strain rate remains the same, the average grain size of the alloy increases with an increase in deformation temperature. The most significant change is at the strain rate of 0.001 s⁻¹, the average grain size increases from 15.15 μm to 35.06 μm with an increase in deformation temperature.

Figure 9(a₁–e₁) show the recrystallization volume fraction of the typical microstructure of the GH4093 alloy after isothermal compression. There are obvious differences in the deformation microstructure obtained under different conditions, which indicates that the process parameters have a great influence on the nucleation and growth of grains during dynamic recrystallization. In the figure, the red part represents the deformed grains in the alloy, the yellow part represents the sub-grains, and the blue part represents the recrystallized grains in the alloy. The recrystallization degree of the alloy is closely related to the change in deformation temperature. Figure 9(a₂–e₂) show an average nuclear orientation difference (KAM) map of the deformation center region of the sample after deformation under different conditions.

When the deformation condition was 1020 °C/1 s⁻¹, the proportion of recrystallized grains was 36%, the proportion of sub-grains was 25.2%, and the proportion of deformed grains was 38.8%. When the deformation temperature was 1050 °C/1 s⁻¹, the proportion of recrystallized grains increased slightly to 38.6%, the proportion of sub-grains increased rapidly to 39%, and the proportion of deformed grains was 22.4%. At the same temperature, when the strain rate gradually slowed down, the volume fraction of recrystallized grains remained basically unchanged at 38.8%, the proportion of sub-grains increased sharply to 50.9%, and the proportion of deformed grains continued to decrease to 10.3%. When the deformation condition was 1050 °C/0.1 s⁻¹. When the strain rate was consistent, with an increase in deformation temperature, when the deformation condition was 1080 °C/ 0.1 s⁻¹, the degree of recrystallization increased rapidly, the proportion of recrystallized grains was 72.3%, the proportion of sub-grains decreased rapidly to 22.4%, and the proportion of deformed grains was 5.25%. At this deformation temperature, with a decrease in strain rate, under the deformation condition of 1080 °C/ 0.01 s⁻¹, the proportion of recrystallized grains remained unchanged at 75.2%, the proportion of sub-grains decreased slightly to 20.1%, and the deformed grains continued to decrease to 4.75%. In summary, it can be found that the deformation conditions have a great influence on the recrystallization behavior of GH3230. With an increase in temperature and a decrease in strain rate, the recrystallization degree of the alloy increases, the percentage of sub-grains increases first and then decreases, and the proportion of deformed grains decreases all the time.

When the deformation temperature is relatively low and the deformation rate is relatively large, the activation energy for the initiation of recrystallization is insufficient, and the recrystallization nucleation time is too short. Under this condition, only a very small number of recrystallized grains are formed, and the original structure basically maintains the deformation state. With a continuous increase in deformation temperature, dynamic recrystallization fully occurs and grains grow, and there are basically no original deformed grains in the microstructure. The driving force of grain boundary migration increases, which accelerates the occurrence of recrystallization. Due to the relatively large deformation rate, the recrystallization nucleation time is limited, so the recrystallization content is relatively low. When the deformation rate continues to decrease, the recrystallization nucleation rate increases slightly. This is because the longer the forming time, the longer the time the dislocations in the alloy grains have to move and aggregate to the grain boundaries, prompting more new fine grains to nucleate and grow at the grain boundaries, and the degree of recrystallization further increases [26].

KAM diagrams can reflect the dislocation distribution inside the grains and grain boundaries. When the deformation temperature is low, the stored energy cannot be consumed quickly, and the stored energy becomes particularly significant in the non-recrystallization deformed grains, as shown in Figure 9(a₂,b₂). As the deformation temperature gradually increases and the strain rate gradually decreases, the stored energy is gradually consumed as recrystallization proceeds, and the energy inside the alloy gradually tends to transfer from the inside of the deformed grain to the grain boundary. As shown in Figure 9(d₂,e₂), the higher the degree of recrystallization, the more the KAM value is biased towards blue [27].

As shown in Figure 10, for the grain boundary distribution of the alloy microstructure under different deformation conditions, it is generally considered that 2~10° represents LAGBs (Low-Angle Grain Boundaries), 10~15° represents MAGBs (Medium-Angle Grain Boundaries), and more than 15° is considered to represent HAGBs (high-angle grain boundaries).

Through observation, it can be found that when the strain rate is kept consistent, with an increase in deformation temperature, LAGBs gradually decrease, and the content of HAGBs gradually increases. This shows that with an increase in deformation temperature, more LAGBs are converted into HAGBs by stacking and aggregating to absorb dislocations, and the maximum content of HAGBs can reach 71.8%. In the microstructure of the sample with a high degree of recrystallization, the content of LAGBs is relatively low, there are mainly HAGBs, and LAGBs are mainly distributed at the edge of the grain. At this time, the KAM value is very small and basically accumulates at the edge of the grain. In the microstructure of the sample with a low degree of recrystallization, the deformed grains are filled with a large amount of LAGBs, while the content of HAGBs is very small. At this time, the area with a large KAM value is also concentrated near the LAGBs inside the deformed grains.

3.5. Twinning Evolution

Studies have shown that nickel-based superalloys are prone to forming twins during their high-temperature forming process. The formation of twins will cause lattice distortion in the crystal, resulting in a larger grain boundary orientation angle, which will be beneficial to the recrystallization process [28,29]. Figure 11 shows a twin boundary microstructure diagram of the GH3230 alloy deformed under different thermal process conditions. The red line segment represents the twin boundary, and the black line segment represents the ordinary grain boundary.

It can be found from the diagram that there is a great percentage of twins in the alloy structure throughout the whole dynamic recrystallization process. Based on this, we study the evolution of twins in the alloy under different processing conditions. The twins formed by deformation have a certain orientation relationship with the original grains in the alloy, which can be expressed as <100>/60°. There is not only one orientation relationship, so it is necessary to introduce a related Coincidence Site Lattice (CSL) model. This model refers to the overlapping part of the intersection of the two lattices. The reference symbol ∑ represents the proportion of the overlapping part. Based on this model, the relationship between the twin boundaries of ∑3n (n = 1, 2, 3) is as follows:

$$\sum 3+\sum 9\to \sum 9$$

$$\sum 3+\sum 9\to \sum 27$$

The overall distribution of twin boundary ∑3ⁿ (n = 1, 2, 3) in the typical microstructure of the alloy was obtained using Channel 5. From Figure 12a, it can be seen that at a deformation temperature of 1020 °C, with a gradual decrease in the deformation rate, the proportion of ∑3 twin boundaries decreases slightly from 18.1% to 13.8%. This may be due to the relatively low deformation rate. During the dynamic recrystallization process, dislocations are piled up at the ∑3 twin boundaries in the original structure, resulting in a small amount of twin boundaries changing the original orientation relationship and gradually becoming random high-angle grain boundaries [30,31]. The trend in the proportion of recrystallization in this process is also the same, and the proportion of the twin boundaries of Σ9 and Σ27 is less than 1%.

Based on Figure 12b, when the strain rate is consistent, as the temperature increases from 1020 °C to 1080 °C, the proportion of the ∑3 twin boundary decreases slowly, which is similar to the change in recrystallization. On the one hand, twins begin to form in the early stage of recrystallization nucleation. With a temperature increase, recrystallization heat is fully activated. The stacking defects and grain growth mechanism in the process of recrystallization nucleation promote the change in the ∑3 twin boundary ratio [32]. On the other hand, the dynamic recovery process is more active at high temperature. Dislocations are rearranged by slip and climb to form a stable sub-grain structure, which reduces the driving force of recrystallization. The combined effect of the two causes the twin boundary change to decrease from 13.0% to 12.8%, and the change is not significant. Meanwhile, the twins will aggravate the bending of grain boundaries and further promote the nucleation rate of recrystallized grains, thus affecting the dynamic recrystallization kinetics [33].

According to Figure 12c, dynamic recrystallization fully occurs and grains grow at a deformation temperature of 1080 °C. With a decrease in deformation rate, the proportion of the ∑3 twin boundary decreases from 34.2% to 12.6%. This is due to the low deformation rate, which makes ‘growth accidents’ more likely to occur, and some twin boundaries extend throughout the whole grain as the recrystallized grains grow. In addition, under different deformation parameters, the percentage of the ∑9 and ∑27 twin boundaries is below 2.0%, which reflects the fact that there is less transformation between ∑3ⁿ (n = 1, 2, 3) twin boundaries, which shows that the ‘growth accident’ mechanism is the main source of ∑3 twin boundaries [34,35]. At the same time, it also comes from the grain mechanism in the process of nucleation and growth.

3.6. Dynamic Recrystallization Behavior

CDRX and DDRX are the two main recrystallization mechanisms of materials during thermal deformation. CDRX occurs through dislocation recombination and sub-grain gradual evolution to form new grains, and the process is continuous and has no obvious boundary. There is no obvious original grain boundary trace, and the sub-grain boundary gradually transforms into a large-angle grain boundary. DDRX refers to the strain-induced local bowing (protrusions) of the original grain boundaries to form new dislocation-free nuclei, which then rapidly grow and replace the deformed matrix [36]. There is a typical structure: the necklace structure (new grains are distributed along the original grain boundaries) [37].

3.7. Continuous Dynamic Recrystallization

The distribution of grains in the deformed microstructure of the GH3230 alloy obtained at 1080 °C/0.01 s⁻¹ is shown in Figure 13. In order to explain the dynamic recrystallization process at this time, two kinds of grains A and B were selected. Grain A refers to incompletely continuous dynamic recrystallization grains, using A1, A2, A3, A4, and A5 to represent it. Grain B refers to completely continuous dynamic recrystallization grains, using B1 and B2 to represent it.

The orientation angle distribution of point to point and point to origin along the black arrow in grain A and grain B is shown in Figure 13. The maximum cumulative orientation difference is 58°, and the minimum is 12°, From A1 to A5, the Euler angles are (61.4°, 34.1°, 87.8°), (41.3°, 35,9°, 5.5°), (42.1°, 44°, 6.1°), (159°, 42.6°, 85.8°), and (42.3°, 44.7°, 6.3°). It can be seen that there is a significant change in the Euler angle along the straight-line direction in the unrecrystallized grains; that is, the sub-grain rotation occurs inside the deformed grains. From Figure 13d,e, it can be seen that the cumulative orientation difference in B1 and B2 reaches 58°. This indicates that there are dislocation accumulation and sub-grains in the grains of the unrecrystallized region [38]. The results show that when the orientation angle from the inner point to the origin of the grain is more than 15°, the combination and rotation of the sub-grains will occur in the structure, and then CDRX will occur.

3.8. Discontinuous Dynamic Recrystallization

The distribution of grains and nuclear mean orientation difference (KAM) in the deformed microstructure of the GH3230 alloy obtained at 1110 °C/0.1 s⁻¹ is shown in Figure 14. In this figure, the fine white line represents LAGBs, and the coarse black line represents HAGBs. It can be clearly seen that a huge number of dislocations accumulate near the grain boundaries, while the dislocation density inside the deformed grains is very small. In addition, due to the huge number of dislocations, the intragranular boundaries and grain boundaries and both sides of the grain boundaries show obvious inconsistency, and the phenomenon of grain boundary bow-shaped nucleation appears at the grain boundary. Moreover, there are a number of sub-grains inside the deformed grains and near the GBs. The above situation indicates that DDRX occurs during the deformation of the alloy.

In the figure, the HAGB (48°) side in the blue region bends from the low-density dislocation region toward the high-density dislocation region. Meanwhile, the sub-grain boundaries formed by dislocation accumulation and entanglement cut off the local bending part of the deformed grains. With the further accumulation of dislocations around the sub-grain boundaries, the sub-grain boundaries are gradually transformed into high-angle grain boundaries, and finally the new grains are separated from the original grains. This process is repeated with continuous high-temperature deformation, resulting in a gradual reduction in the deformation area [39]. As the deformation temperature gradually increases, the DRX mode gradually changes from CDRX to DDRX.

4. Conclusions

The constitutive equation of GH3230 was established according to the Arrhenius hyperbolic sine form. Under the conditions of a deformation temperature of 1020~1110 °C and strain rate of 1~0.001 s⁻¹, the constitutive relation can be expressed as follows: $\dot{\epsilon }=exp\left(36.123\right){\left[sinh\left(0.00587\sigma \right)\right]}^{4.7946}exp\left[-451.507/\left(RT\right)\right]$. The average activation energy of deformation at different deformation temperatures and different strain rates is 451.507 kJ/mol, and dynamic recrystallization occurs during deformation.

A power dissipation diagram and hot processing diagram of the GH3230 alloy under different deformation parameters were established. The peak value of η is 0.34. Under 1020~1050 °C and 1 s⁻¹, the deformed alloy is prone to flow instability defects. The suitable hot processing range of GH3230 is 1020~1110 °C, and the strain rate is 0.1~0.001 s⁻¹.

With an increase in deformation conditions, the average grain size of the GH3230 alloy grows from 16.86 to 35.06 μm. The deformation temperature and strain rate have an enormous impact on the dynamic recrystallization of the GH3230 alloy. With an increase in deformation temperature, the degree of recrystallization of the alloy increases. With a decrease in strain rate, the degree of recrystallization of the alloy increases, and the maximum recrystallization volume fraction can reach 75.2%. During the dynamic recrystallization process, a large number of twins are produced, and the change trend in the ∑3 grain boundary is similar to that in the recrystallization volume fraction.

The recrystallization mechanism of the GH3230 alloy under different process parameters includes CDRX and DDRX. The continuous dynamic recrystallization process is dominated by the rotation of sub-grains accompanied by the transformation of LAGBs to HAGBs, and DDRX is dominated by the nucleation of grain boundary bowing. With an increase in deformation temperature and a decrease in strain rate, the recrystallization mechanism changes from continuous dynamic recrystallization to discontinuous dynamic recrystallization.