Deformation behavior of a new Ni-Co base superalloy GH4251 during hot compression

The deformation behavior of a new Ni-Co base superalloy GH4251 under hot compression tests within the deformation strain window of 0.36 to 1.2 was investigated in the temperature range of 1050 ∼ 1170 °C and strain rate range of 0.001 ∼ 1 s−1. Based on the analysis of true stress-strain curves, constitutive equations were established to describe the rheological behavior during hot compression. Microstructure evolution was investigated by transmission electron microscope (TEM), scanning electron microscope (SEM) and optical metallography (OM). The results show that flow behavior of GH4251 alloy is combinedly determined by the effect of work hardening and dynamic recrystallization (DRX). The deformation activation energies at strain of 0.36 to 1.2 are calculated to be 311 ∼ 536 kJ mol−1 in the super-solvus temperature region, and 796 ∼ 1064 kJ mol−1 in the sub-solvus temperature region. The recrystallization nucleation mechanism of GH4251 alloy is strain induced grain boundary migration (SIBGM). The occurrence and expansion of recrystallization are strongly promoted by high deformation temperatures and high strain rates, while the DRX grain size increases with elevated deformation temperature. When the deformation temperature is below 1090°C, the recrystallized grain can be extremely small (<17μm), which is rather independent on strain and strain rate. However, above 1110 °C the grain size at strain rate of 0.001s−1 is significantly larger than that of higher strain rates. The difference can be ascribed to the presence of γ′ phase, with which the development of dynamic recrystallization is postponed, while the growth of recrystallized grains is inhibited as well.


Introduction
Aiming at the material demand for the development of aircraft engines and land-based gas turbines, superalloys with superior mechanical properties, microstructure stability, and especially excellent workability have been developed successively. At present, nickel-based superalloy remains the most widely used turbine disk materials. These alloys obtain high temperature capability and strength by continuously increasing the alloying level of aluminum, titanium and refractory metal elements [1][2][3]. However, the high alloying of γ′ forming elements makes the homogenization processes challenging and often results in the poor formability and cracking during subsequent fabrication [4].
Recently, a series of Ni-Co-base alloys were designed through combining the the characters of Ni-base and Co-base superalloy, which are both strengthened by γ/γ′ two-phase structure. This strengthening effect results from the large coherent strain between the disordered phase matrix γ (based on Ni or Co, face-centered cubic structure) and the coherent ordered phase γ′ (Ni 3 Al in Ni-base alloy and Co 3 (Al,W) in Co-base alloy, L1 2 crystal structure). These alloys are modified by increasing the contents of Co and Ti elements based on the chemical compositions of U720Li alloy (a hard-deformed nickel-based superalloy), where Co promotes the formation of a more stable microstructure and Ti enhances the effect of phase interface strengthening by substituting Al atom in γ′ phase. It has been reported in the literatures that the Ni-Co alloy has a high heat resistance temperature, exceeding 50°C of the most advanced commercial cast and wrought (C&W) Ni based alloy U720Li, as well as a wide processing window in traditional C&W route [5][6][7]. Therefore, it is considered to be a promising candidate material for the hot-end components in aircraft engines. However, problems are also reported for these materials, such as the structural instability due to η phase transformation at high temperature, and the cracking susceptibility under some deformation conditions [8].
Formability is considered the premise of manufacturing the core hot-end component with favorable microstructure and required properties. As for the prototype alloys of Ni-Co based alloys, plenty of studies have been carried on investigating the hot working performance of nickel-based alloy U720Li and cobalt-based alloys. For hard-deformed U720Li alloy, the hot working window is determined by considering the dynamic recrystallization (DRX) behavior [9,10], the dynamic precipitation of γ′ phase in hot deformation and its interaction with DRX [4,11], as well as the high temperature plasticity [12], and the corresponding microstructure control and superplasticity under dual-phase deformation [13][14][15]. On the other hand, the investigations on DRX nucleation mechanism (both discontinuous and continuous), microstructure evolution, constitutive equations and activation energy were widely conducted on the Co-based superalloys as well [16][17][18].
Over the past decade, many reports have been focused on the phase precipitation behavior and mechanical properties of the Ni-Co based alloy [7,19]. In contrast, few research can be found on the workability of these alloys, while the systematic researches on the hot deformation behavior of superalloys with medium cobalt content (20%-35%) are rather lacking [20,21]. Liu et al [21,22]investigated the hot deformation behavior of a Ni-Co alloy by isothermal compression tests and constructed the hot processing map to derive the optimal work conditions. The effect of strain rate on the microstructural evolution and deformation mechanism at γ′ subsolvus and γ′ super-solvus temperatures were studied, while the deformation feature of primary γ′ particle and its influence on DRX were also mentioned. However, dynamic recrystallization expansion and grain size evolution in single-phase and dual-phase deformation condition in a wider temperature range and larger strain rage is still unclear, which needs further investigations.
In this paper, the deformation behavior of a new type of Ni-Co alloy, GH4251, is studied by hot compression tests. This alloy is developed by optimizing Ti/Al ratio and Nb content of the existing Ni-Co alloy to solve the problem of η phase transformation at high temperature and the corresponding microstructure instability. In order to provide theoretical support for its turbine disc forging production, a wide range of true strain (maximum of 1.2) and temperature (from 50°C higher than γ′ solvus to 50°C lower than γ′ solvus) was chosen to evaluate the formability of this alloy. The selection of strain rate in testing is based on the equipment condition of forging hydraulic press, which is always lower than 1s −1 . The effects of dynamic precipitation and dissolution behaviors of γ′ particles on DRX process and grain size under different deformation conditions were investigated. Constitutive analysis of different strain was carried out under both single-phase and dual-phase deformation condition. The corresponding deformation activation energy was compared and interpreted.

Experimental
The nominal composition of GH4251 alloy in the present study is shown in table 1. The alloy was manufactured by vacuum induction melting (VIM) + protective atmosphere electroslag remelting (ESR) + vacuum arc remelting (VAR). After homogenization, the ingot was forged into a bar with a diameter of Φ120 mm (figure 1). Cylindrical compression samples with the size of Φ10 × 15 mm were cut from the mid-radius of the bar. The initial microstructure contains equiaxial grains of about 45 μm in size and quite a few primary γ′ precipitates (2-3μm), as well as secondary γ′ precipitates (500-600nm) and tertiary γ′ precipitates (less than 50nm).
Hot compression tests were carried out on a MTS810 thermo-mechanical simulator at the temperatures of 1050, 1070, 1090, 1110, 1130, 1150 and 1170°C, and strain rates of 0.001, 0.01, 0.1 and 1.0s −1 . Engineering strains of 30%, 50% and 70% (true strain of 0.36, 0.7, 1.2) were chosen to study the microstructural evolution during hot compression. The schematic diagram of the hot compression experiment is illustrated in figure 2. All specimens were heated to test temperature at a heating rate of 20°C s −1 and held for 15min to ensure temperature uniformity before deformation [12,16]. Graphite lubricant was placed between the specimen and compression die to reduce the effect of friction. The load-displacement data were automatically output by MTS testing machine and converted into true stress-true strain curve. Three samples were compressed in each deformation condition, with the average value being the final results. Deformed specimens were quenched immediately after compression and sectioned along the load direction. In addition, in order to study the initial microstructures just prior to hot compression, some samples were heated up to test temperature and held for 15 min followed by water quenching immediately. Compared to the microstructure of forged bar, the grain size of the samples heated at 1050°C-1090°C has no obvious change. The grain size of the samples heated at 1110 ∼ 1170°C increases to 67.5,151.6,198.4 and 221.1 μm, respectively (figure 3). Samples for optical metallographic examination were mechanically polished and boiled in the solution of 2.5gKMnO 4 + 10mlH 2 SO 4 + 90mlH 2 O for 30min [23,24]. Samples for morphology observation of precipitated phase were electro-etched (3V) in the solution of 170ml H 3 PO 4 + 10 ml H 2 SO 4 + 15 g CrO 3 [21,24,25]. Foils for transmission electron microscope (TEM) measurement were firstly machined from the deformed samples with a thickness of 0.3mm. After hand grinding to 50μm, samples were thinned using a twin-jet technique in the electrolyte of a solution of 10vol% HClO 4 in ethanol [22].
According to the national standard GB/T6394-2002, the intercept method was used to measure the grain size of three random fields of view for each sample. The morphology analysis of precipitates was performed on ZEISS sigma scanning electron microscopy (SEM) through secondary electron mode. TEM observations were carried out on a FEI Tecnai G2 F20 operated at 200kV. Bright field morphology and selected area electron diffraction were used to investigate the dislocation configuration and zone axis.   value of 0.2% residual deformation as its yield limit. The elastic modulus is obtained by calculating the ratio of stress and strain in the elastic proportional deformation stage. It can be seen that under all deformation conditions, the true stress increases rapidly at the beginning of compression. Then the curve exhibits a slow growth to the peak stress followed by a gradual decrease. Steady state was observed in the medium strain stage of hot compression. The stress increases again at large strain in some flow curves. On the contrary, steady state was kept until the end of compression in the flow curves of 1s −1 strain rate and 0.1s −1 /1170°C.  An abrupt decrease in the flow stress during initial straining (strain less than 0.03), known as the yield drop phenomenon, was observed in the flow curves of 1130 ∼ 1170°C at strain rate of 0.1 s −1 and 1110 ∼ 1170°C at strain rate of 1 s −1 (figures 4(c), (d)). The stress of these curves exhibits a further rise after yield drop until the peak stress is reached. In literature, such yield drop behavior has also been observed in other nickel-based superalloys [26,27]. Explanations for this phenomenon include the following theories: A general explanation is that after the pinning of dislocations by solute atoms a higher stress is needed to mobilize these locked dislocations [28]. According to the dynamic deformation theories, dislocation density increases rapidly at the initial stage of deformation, which necessitates a higher stress for further dislocation motion [27]. In the study of Guimaraes [29], short range ordering of γ′ forming elements (Ti and Al) as well as carbides are responsible for locking dislocations in Waspaloy in the yield stage. Sharma et al [26] attributes this effect to the decrease of stacking fault energy (SFE) due to Co and Nb elements which makes the slip of partial dislocations difficult. The yield drop phenomenon observed in the present work can be explained by almost all the above theories.
The true stress of GH4251 alloy increases with decreasing temperature and increasing strain rate during hot compression. At the same strain rate, stress difference between curves at 1050 ∼ 1110°C is significantly larger than that at 1130 ∼ 1170°C. The difference between peak stress and steady stress in a single curve is also larger in the former temperature range. According to the calculation result by JmatPro software, the solution temperature of the main strengthening phase γ′ in GH4251 is 1115°C ( figure 5). Therefore, γ′ phase is inferred to have a great influence on the hot deformation behavior of GH4251 alloy. The shape of flow curve reflects the alloy's work hardening and dynamic softening mechanism during hot compression. The flow curve of GH4251 alloy has typical dynamic recrystallization characteristics, while the subsequent microstructure analysis also reveals the present of DRX during hot compression. The interpretation of flow behavior through microstructure will be carried out later in this paper.

Microstructure analysis
Metallographic analysis indicates that DRX occurred in GH4251 alloy under all the deformation conditions in this paper, while the recrystallization fraction and grain size are affected by deformation conditions. No cracks are observed in all GH4251 samples during hot compression. Typical microstructural evolution of hot deformation and the corresponding flow curve is shown in figure 6 (take the sample of 1090°C/0.01s −1 as an example). A large number of fine equiaxed grains (16.0 μm) are observed at strain of 0.36 ( figure 6(b)). These uniform grains can be determined as dynamic recrystallization grains because of the smaller size than initial grains (45.6 μm). Large original elongated grains (marked by dashed lines figure 6(b)) also exist in the samples, indicating the partial DRX process. When the strain increases to 0.7, initial deformed grains disappear (figure 6(c)). The microstructure of the sample with 1.2 true strain does not change significantly, only with finer grains (figure 6(d)). It is note worthy that γ′ particles are found to dispersively distribute in the three samples of 1090°C/0.01 s −1 with different strains. Complete dynamic recrystallization process occurred in all the samples at strain of 1.2. The corresponding DRX grain size, which was measured by intercept method, is listed in table 5. Typical deformed microstructure at strain of 1.2 is shown in figure 7. Compared with the grain size of heated but uncompressed samples (figure 3), the microstructure after hot deformation has been refined to varying degrees. Very fine recrystallized grains (<7 μm) were obtained at the temperature of 1050 ∼ 1070°C (table 5). DRX grain size increases with increasing deformation temperature at the same strain rate, especially at 1150 ∼ 1170°C. The effect of strain rate is more complex. When compression temperature is lower than 1090°C, strain rate has little effect on the grain size. In contrast, when compression temperature is higher than 1110°C, DRX grains are significantly larger at strain rate      figure 8, where volume fraction of DRX greater than 95% is considered complete. It can be seen that all the samples of 1050 and 1070°C experiences partial recrystallization at strain of 0.36. When compression temperature is 1090°C, complete DRX process only occurs in the sample of 1s −1 strain rate. In contrast, as deformation temperature increases to 1110°C, only the samples of 0.001s -1 do not fully recrystallized ( figure 8(a)). Overall, low temperature and low strain rate are not conducive to the development of dynamic recrystallization. The influence of deformation parameters on recrystallization completion at strain of 0.7 is similar to that at 0.36, while parameter range of partial DRX trends towards lower temperature and lower strain rate ( figure 8(b)). Traditionally, high-speed deformation is considered to be harmful to the nucleation and expansion of recrystallization because of the insufficient time for atom diffusion under thermal activation [30]. However, some literatures reported the opposite results [22,31]. In this paper, large deformation resistance and strain rates may cause significant adiabatic heating during deformation, which will promote the nucleation and growth of DRX grains. Figure 9 shows the influence of deformation conditions on grain size evolution of GH4251 alloy during hot compression. At the temperature of 1110°C and strain of 0.36, the grain size of all deformed samples decreases sharply compared to the original size (heated but not compressed), and decreases with the increase of strain rate ( figure 9(a)). Furthermore, the grain size does not change significantly when strain increases to 0.7 and 1.2 under the strain rate of 0.01 ∼ 1 s -1 . On the contrary, when the strain rate is 0.001s -1 , the minimum grain size is obtained at strain of 0.7, while grains coarsen to nearly original size at 1.2 strain. This phenomenon is attributed to the slow development process of recrystallization at the strain rate of 0.001s −1 . According to figure 8(b) (strain of 0.7), DRX process is incomplete only at 0.001s −1 strain rate at 1110°C. At this time, more fine recrystallized grains appeared in the sample with 0.7 strain than that with 0.36 strain, leading to a smaller average grain size. When strain increases to 1.2, DRX has been completed, new grains grow apparently because of the very long time of deformation (duration of deformation at 0.001, 0.01, 0.1 and 1 s −1 strain rate is about 1200s, 120s, 12s, and 1.2s, respectively). At the same strain rate of 0.01s −1 , all samples had obvious grain refinement at the strain of 0.36, and the grain size remained basically unchanged or slightly decreased during the subsequent deformation process ( figure 9(b)). Under the same deformation, the grain size decreases with decreasing deformation temperature (average grain size at 1050°C was measured only for the sample with complete DRX at strain of 1.2).

Constitutive analysis for GH4251
In order to study the deformation behavior of GH4251 alloy in the whole process of hot compression, the constitutive equations were constructed at true strain of 0.36, 0.7 and 1.2, respectively. Three types of constitutive equations, all of which include an Arrhenius term, are commonly used to represent the correlation between work conditions and flow stress during hot deformation (equations (1)-(3) [30,32].
Z is the Zener-Hollomon parameter, representing the thermal activation corrected strain rate; Q is the activation energy of deformation; R is the universal gas constant; A 1 , A 2 , A, β and α are constants and n, n 1 are the stress exponents. The three constitutive equations are suitable for deformation conditions of different stress levels.
Hyperbolic-sine type equation (3), which was introduced by Sellars and Tegart, is considered to be a more general form suitable for a wide stress range [16,21]. Thus, equation (3) is chosen to describe the constitutive relationship of GH4251 in this study. Since the thermal deformation is different below and above the γ′ solvus temperature (1115°C in figure 5), the constants of constitutive equations were calculated for sub-solvus temperature region, super-solvus temperature region and full temperature range, respectively. Natural logarithm linear regression method was used to calculate the above parameters. In order to optimize the value of α (α = β/n 1 ), equation ( It can be seen from table 6 that the deformation activation energies of GH4251 in super-solvus temperature region and sub-solvus temperature region are in the range of 311 ∼ 536 kJ mol −1 and 796 ∼ 1064 kJ mol −1 , respectively. The former is comparable with the Q values of most Ni-based superalloys, such as precipitated strengthened superalloy 740H (25Cr-20Co-2Nb-1.5Al-1.5Ti, 357kJ mol −1 ) in single-phase deformation condition [31], solution strengthened superalloy G3(22Cr-20Fe-7Mo-1.5Co, 361kJmol −1 ), lower than alloy C276 (15.5Cr-6Fe-16Mo-4W, 515kJ mol −1 ) [33]. It should be noted that the activation energy of above alloys for comparison are all calculated at the peak strain, which is in the low or middle strain part of flow curves. It has been reported that activation energy of thermal deformation reflects the ability of work hardening and softening, which is related to the alloy composition and microstructure [33]. The addition of most alloying elements will increase the Q value in single-phase deformation condition. The first reason is lattice distortion caused by atomic radius difference between solute elements and matrix, leading to the impediment of dislocation movement. This solid solution strengthening effect increases the work hardening level in hot working. The second reason is the reduction of dynamic recovery softening. In the view of diffusion, the calculated activation energies of GH4251 are much larger than the self-diffusion activation energy of nickel (292 kJ mol −1 ), cobalt (265 kJ mol −1 ) or the diffusion of any other alloying element such as Al (269 kJ mol −1 ) or Ti (256 kJ mol −1 ) in nickel [20,33]. This means adding of solute atoms will reduce the effective diffusion rate in the lattice, which greatly increases the activation energy for elements diffusion. Since the diffusion-controlled vacancy movement and dislocation climbing are both the most important mechanisms for dynamic recovery, increased activation energy for elements diffusion will limit recovery-softening and increase the flow stress. The additional Co and Nb in GH4251 will further weaken the effect of dynamic recovery by reducing the stacking fault energy and inhibiting the cross-slip of dislocation [20,34,35]. It can be seen that the above two reasons are both related to the atomic radius of the solid solution element [34]. Therefore, the content of large atomic radius elements, e.g., W, Mo, Nb, Ti determines the activation energy of GH4251 alloy in single-phase deformation region, which is similar to 740H (3.5%Nb+Ti) and G3 (7%Mo), lower than C276 (16%Mo+4%W). The activation energies of GH4251 in dual-phase deformation range (1130°C-1170°C) is much larger, which is comparable with the Q values of some Co-based and Ni-based superalloys, e.g., 1182 kJ mol −1 for Co-20.4Ni-9.8Al-7.4W-2.7Ti, and 930 kJ mol −1 for IN939, lower than Waspaloy (1400 kJ mol −1 ) and U720Li (1552 kJ mol −1 ) [18,20,21,31]. It is believed that this aggravate hardening behavior is caused by γ′ phase precipitation. However, compared with Ni-based superalloys, the composition of γ′ precipitates in Ni-Co based GH4251 alloy is converted from Ni 3 Al to (Ni,Co) 3 (Al,Ti) after the addition of Co and Ti, and the strength of (Ni,Co) 3 (Al,Ti) is lower than that of Ni 3 Al above 1000°C [20].
It should be noted that the activation energy decreases with strain in sub-solvus temperature range, while exhibits an opposite trend in super-solvus temperature range. Mathematically in constitutive equation, activation energy reflects the sensitivity of true stress to strain rate and temperature, indicating the influence of deformation parameters on microstructure. In the dual-phase deformation region, DRX is not completed under most deformation conditions at strain of 0.36 (still in the expansion stage), and the recrystallized fraction is sensitive to the deformation parameters, leading to a higher activation energy. With the increase of strain, complete recrystallization microstructure can be obtained under more deformation conditions, which leads to less sensitivity of microstructure to deformation parameters, resulting in the decrease of Q value. At the strain of 1.2, DRX fraction does not change with deformation parameters (complete recrystallization), and grain size difference is not significant except for the condition of 1110°C/0.001 s −1 , which corresponds to the lowest value of deformation energy in sub-solvus temperature range. On the contrary, DRX is already complete at the strain of 0.36 and 0.7 in single-phase deformation region, resulting in the similar low activation energy values (table 6). However, when strain increases to 1.2, grain size varies greatly under different deformation parameters due to the absence of pinning effect of γ′ phase, which means the increasing sensitivity of microstructure to work conditions, leading to a larger Q value.

Nucleation mechanism of DRX
According to the metallographic analysis, DRX is the critical softening mechanism of GH4251 alloy during hot compression, which is consist with other superalloys. The addition of a large number of alloy elements reduces the stacking fault energy of GH4251 and inhibits cross slip and climbing of dislocations, weakening the effect of dynamic recovery [21,34,36]. The occurrence of complete DRX at lower deformation temperature also proves this view ( figure 7). Figure 11(a) shows the deformation microstructure of GH4251 compressed by 30% at 1070°C/0.01 s −1 . Fine recrystallized grains (<10 μm) are formed at the grain boundaries (GBs) and triple junctions (TJs) of prior grains. Bulges and serrations also appear at deformed grain boundaries (marked by arrows). The bright field transmission electron microscope (TEM) image under the same deformation condition is shown in figures 11(b) and (c). In figure 11(b), dislocation tangles and local dislocation density difference are observed, which indicates the heterogeneous deformation during hot compression. The weak dynamic recovery of this alloy suppresses the dislocation elimination, promoting the local accumulation and entanglement of dislocations. In addition, the barrier effect of grain boundary and twin boundary on dislocation movement accelerates the formation of dislocation density gradient. Local interaction between dislocations and grain boundary leads to the appearance of grain boundary bulges. It is commonly noted that DRX nuclei can be formed from these local bulges [16,37,38]. These nuclei expand into the inside of deformed grains with high dislocation density through grain boundary migration, which is driven by the reduction of strain energy. This discontinuous dynamic recrystallization mechanism is called strain induced grain boundary migration nucleation (SIGBM). For GH4251 alloy, dislocation tangles, local dislocation density gradient, DRX nucleus with very low dislocation density are observed simultaneously in sample of 1070°C/0.01 s −1 with 30% deformation ( figure 11(c)), which proves that the DRX nucleation mechanism of GH4251is SIGBM.
Both DRX nucleus with few dislocations and growing DRX grain containing poorly developed substructures were found in the sample in figure 11(c), which indicates the growth process of recrystallized grains during hot compression. The growth of DRX nucleus and development of dislocation substructure can be described by a schematic graph (figure 12) [39]. Here ρ 0 is the dislocation density of a fully annealed grain, ρ c is the critical value for DRX nucleation and D is the grain size. When DRX nucleus is formed, it contains few dislocations compared to adjacent deformed grains ( figure 11(c)). The driving force of grain growth for small DRX grains is the difference of dislocation density with parent grains. DRX grain nucleus grows with the hot deformation, while its interior dislocation substructure develops at the same time. When the dislocation density is equal to the adjacent grains, the growth of DRX grains will stop ( figure 12). With the increase of deformation, dislocation entanglement and local dislocation density gradient will also be formed in the DRX grains, which will bulge the DRX grain boundary and form new recrystallized grains, i.e., multi-cycle dynamic recrystallization. Figure 11(d) is the microstructure of GH4251 compressed by 70% at 1070°C/0.01 s −1 . When dynamic recrystallization has been completed, grain boundary bulges and serrations can also be found. Newly formed DRX grains are observed in the triple junction of previous recrystallized grains (marked in figure 11(d)).
Based on the DRX behavior of this alloy, the flow curves of GH4251 alloy can be interpreted as followed. At the beginning of hot compression, the true stress increases rapidly through dislocation accumulation, which is accelerated due to the retardation of cross slip and climbing. With the increase of deformation, dynamic recrystallization nucleated by strain-induced grain boundary migration. The soften effect of DRX increases with its expansion, leading to the slowdown of stress growth rate (figure 4). When softening effect increased to be equivalent to the hardening effect, the peak stress is obtained. Subsequently, the softening effect is further enhanced and the true stress continues to decrease. At the same time, work hardening occurs in the recrystallized grains. After reaching a critical deformation, multi-cycle DRX occurs. When work hardening and dynamic softening reach equilibrium again, the flow curve presents steady state. In figure 4, the stress of some curves increases again at large strain. Some literatures believe that this phenomenon is related to the weak recovery effect of low stacking fault energy materials and the multiple rounds of recrystallization nucleation [33]. The research on the deformation behavior of nickel -cobalt based alloys at very large strain is undergoing.

Effect of γ′ phase on deformation behaviors
The first influence of γ′ phase on hot deformation of GH4251 is the retardation of dynamic recrystallization. Incomplete DRX microstructure is observed at strain of 0.36 and 0.7 under 1050 ∼ 1110°C ( figure 8). The occurrence and expansion of DRX needs atomic diffusion under thermal activation to promote nucleation and grain boundary migration, so low deformation temperature is not conducive to recrystallization. In addition, many studies have shown that fine and dense second phases will hinder grain boundary migration, thereby inhibiting recrystallization nucleation [9,39,40]. The nucleation mechanism of GH4251 is SIGBM, which depends on the formation of local dislocation density gradient at grain boundary and triple junctions of grain boundary. A large number of dispersed γ′ phases will pin the moving dislocations, which impedes the increase of dislocation density gradient at the above nucleation sites and further retards DRX process. Figure 13(a) is the bright field TEM image of grains in the GH4251 alloy under 1050°C/0.1 s −1 /30% deformation conditions in [011] zone axis, which is proved by the selected area electron diffraction (SAED) pattern in figure 11(b). Besides, the SAED pattern also shows that this structure is γ′ phase, as marked in figure 13(a). Dislocation tangles can be found near the γ′ phase, however, the dislocation density gradient is not obvious in the adjacent grains. Under the combining effect of hindering dislocation movement and delaying recrystallization by γ′ phase, the work hardening of GH4251 alloy at 1050 ∼ 1100°C is remarkable, resulting in higher flow stress, especially peak stress ( figure 4).
Another effect of γ′ phase in the compression of GH4251 is to effectively hinder the growth of recrystallized grains, which is particularly obvious under low-speed deformation. The calculated dissolution temperature of γ′ is 1115°C, and the mass fraction of this phase decreases rapidly above 900°C (figure 5). Considering the kinetic factors of heating process, the actual solvus temperature of γ′ phase is slightly higher than the thermodynamic equilibrium calculation result. According to the microstructure analysis, DRX grains are extremely small when the compression temperature is lower than the solvus temperature of γ′ (1050 ∼ 1090°C), and do not grow significantly at the strain rate of 0.001 s −1 up to strain of 1.2 (compression time about 1200 s, table 5). In contrast, when compression is carried out in single-phase region, the grain size is obviously larger. It should be noted that when GH4251 alloy is deformed near the γ′ solvus temperature (1110 ∼ 1130°C), the grain size after long time compression at 0.001 s −1 is quite larger than that of other strain rate, which is attributed to the mass fraction decrease of γ′. In this case, long-time deformation with weak pinning effect leads to an increase in grain size.
It is note worthy that three types of γ′ particles with different sizes exist in the original microstructure of cogging bar. Figure 14 shows the comparison of microstructure between cogging bar and heated but not deformed samples of 1110°C and 1150°C. γ′ particles with uniform size of 20-30 nm are observed in both the samples of 1110°C and 1150°C, while another type of γ′ particles with a larger size of 1μm exist only in the sample with deformation temperature of 1110°C. Studies have shown that small γ′ particles are easier to dissolve and precipitate during heat treatment [15]. It can be inferred that water cooling cannot inhibit the precipitation of the smallest size γ′ phase, and the large size γ′ particle in the sample of 1110°C is the residual phase which is not fully dissolved during the heat process before compression. Furthermore, GH4251 alloy exhibits single-phase deformation during hot compression at 1150°C.
In summery, GH4251 alloy does not crack during single-phase and dual-phase deformation in the hot compression tests in this paper. Deformation defects, such as share bands and grain boundary separation are not observed. Although the existence of γ′ phase at low temperature delayed the development of DRX, fully recrystallized microstructure was still obtained at stain of 1.2. Therefore, the deformation ability of GH4251 alloy is better than that of difficult-to-deform superalloys such as U720Li and ЭК151.

Conclusions
The hot deformation behaviors of a new Ni-Co based alloy GH4251 were investigated over the deformation temperature range of 1050 ∼ 1170°C and strain rate range of 0.001 ∼ 1 s −1 with strain of 0.36, 0.7 and 1.2 by hot compression tests. The main conclusions can be drawn as follows: (1)The true stress of GH4251 during hot compression increases with increasing strain rate and decreasing temperature. Hyperbolic sine equation can properly model the hot deformation behavior. The activation energies at strain of 0.36, 0.7 and 1.2 were calculated to be 1064, 896, 796 kJ mol −1 in γ+γ′ dual-phase region and 311, 369, 536 kJ mol −1 in γ single-phase region, respectively.
(2)Complete dynamic recrystallization process occurred in all the samples at strain of 1.2. Strain induced grain boundary migration (SIGBM) is the critical nucleation mechanism of DRX. High deformation temperature and high strain rate are beneficial to the nucleation and expansion of DRX. Recrystallized grain size increases with deformation temperature. Very fine and uniform DRX grains (<17 μm) can be obtained at deformation temperature below 1090°C, regardless of strain and strain rate. When deformation temperature is above 1110°C, the grain size at strain rate of 0.001 s −1 is significantly larger than that at other rates.
(3)Water cooling cannot inhibit the precipitation of the tertiary γ′ phase in GH4251 alloy. γ′ phase delays the nucleation and development of dynamic recrystallization and hinders the growth of recrystallized grains, which leads to a larger peak stress in the flow curve of γ+γ′ dual-phase region.