- Email: [email protected]
Effects of annealing parameters on microstructural evolution of a typical nickel-based superalloy during annealing treatment
Materials Characterization 141 (2018) 212–222
Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Effects of annealing parameters on microstructural evolution of a typical nickel-based superalloy during annealing treatment ⁎
T
⁎
Ming-Song Chena,b, , Zong-Huai Zoua,b, Y.C. Lina,b,c, , Hong-Bin Lid, Wu-Quan Yuana,b a
School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China State Key Laboratory of High Performance Complex Manufacturing, Changsha 410083, China c Light Alloy Research Institute of Central South University, Changsha 410083, China d College of Metallurgy and Energy, Hebei United University, Tangshan 063009, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Superalloy Annealing treatment Microstructure Homogeneity
It is usually one of important goals to obtain the homogeneous microstructure for forgings by hot working, such as hot die forging. However, there is always a large inhomogeneous deformation in die forgings during hot deformation. It often results in an inhomogeneous microstructure of forging due to the different dynamic recrystallization (DRX) fractions and grain sizes in different strain regions. In this study, a method was proposed to improve the homogeneity of deformed microstructure by annealing treatment. The microstructural evolution of a typical nickel-based superalloy during annealing treatment has been investigated. The Optical Microscope (OM), Electron Backscatter Diffraction (EBSD) and Scanning Electron Microscope (SEM) technologies were applied to observe microstructures. It is found that adopting an annealing treatment after deformation can effectively improve the homogeneity of microstructure due to the static recrystallization. There is an incubation period which is not less than 5 min for the occurrence of static recrystallization. The initial heterogeneous grains become homogeneous firstly, and then heterogeneous again when the annealing time is increased from 5 to 30 min. The deformation accelerates the dissolution of delta phase in annealing treatment. The fraction of delta phase decreases at the early stage of annealing whether the temperature exceeds the solution temperature. Based on a comprehensive consideration of all factors, the optimum annealing parameters, which are the temperature of 980 °C and the time of 10 min, have been obtained.
1. Introduction
established a DRX kinetics model by a new method. CA models were established by Liu et al. [25] and Reyes et al. [26]. Prithiv et al. [27] evaluated the efficacy of recrystallization vs. strain induced boundary migration in achieving grain boundary engineered microstructure in a Ni-based superalloy. In addition, the hot deformation behavior of superalloy FGH96 was investigated by Zhang et al. [28] and Liu et al. [29]. Yeom et al. [30] and Brand et al. [31] simulated the microstructural evolution by finite-element (FE) method. The properties of parts mainly depend on their microstructures [32,33]. Generally, fine and homogeneous microstructures can effectively improve the fatigue resistance and mechanical strength [34–36]. However, the microstructure after deformation is usually heterogeneous due to the inhomogeneous deformation [37]. It has been found that there is a severe inhomogeneous deformation on the sheet surface during uniaxial tensile testing [38]. Inhomogeneous deformation is determined by thermal deformation parameters such as local loading parameters, friction and die-forging mode [39]. It is difficult to find the perfect thermal deformation parameters that can avoid the
Nickel-based superalloys are applied to make the key parts in energy and aerospace industries due to their fine corrosion and fatigue resistance, as well as prominent high-temperature characteristics [1–3]. These key parts, such as turbine disk, often need to be formed by hot die forging. To make a proper hot working processing, some studies on the flow behaviors as well as microstructural evolution of Ni-based superalloys have been done [4–7]. For example, Lin et al. [8,9], Ning et al. [10], Zhang et al. [11], Wu et al. [12], Etaati et al. [13], Yu et al. [14] and Momeni et al. [15] established the constitutive models to accurately predict flow behaviors. Wen et al. [2,16], Zhang et al. [17], Liu et al. [18], He et al. [19] and Wang et al. [20] established the relevant processing maps to gain optimal hot deformation parameter domains. Mignanelli et al. [21] and Retima et al. [22] analyzed the effects of heat treatment on microstructural evolution. To predict the DRX behavior, Chen et al. [23] researched the DRX behavior through thermal compressive tests, and established a segmented model. Chen et al. [24]
⁎
Corresponding authors at: School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China. E-mail addresses: [email protected] (M.-S. Chen), [email protected] (Y.C. Lin).
https://doi.org/10.1016/j.matchar.2018.04.056 Received 10 February 2018; Received in revised form 27 April 2018; Accepted 30 April 2018 Available online 01 May 2018 1044-5803/ © 2018 Elsevier Inc. All rights reserved.
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
inhomogeneous deformation. In other words, the inhomogeneous deformation is unavoidable during die forging. It is well known that the fraction and grain size of DRX depend on the temperature, strain and strain rate [40–43]. Therefore, the DRX fraction and grain size in different regions of forging with different strains are also different, which often results in the inhomogeneous microstructure in forgings. So, it is necessary to adopt a method to obtain fine and homogeneous microstructures. However, few attentions focus on the way to improve the homogeneity of deformed microstructure, although there are a lot of researches focus on the flow behaviors and microstructural evolution [44–47]. In this study, a method to improve the homogeneity of deformed microstructure by annealing treatment was proposed because of the great influence of annealing parameters on phase evolution and recrystallization behavior [48]. Firstly, the feasibility to obtain homogeneous microstructure by annealing treatment has been verified. Secondly, the effects of annealing parameters on microstructural evolution have been studied to obtain the optimum annealing parameters.
Table 1 The chemical compositions (wt%) of nickel-based superalloy GH4169. Ni
Cr
Nb
Mo
Ti
Al
Co
C
Fe
52.82
18.96
5.23
3.01
1.00
0.59
0.01
0.03
Bal
Holding
Temperature
1 min Compression
T: 950, 965,980, 1010 C T: 950 C : 0.1s
t: 5,10,30,60,180min
-1
: 0.69 Heating rate 10 C/s
Water quenching
Heating rate 10 C/s
Water quenching
2. Materials and Experiments A commercial nickel-based superalloy of which chemical compositions are shown in Table 1 was used. The experimental cylindrical specimens with a size of ϕ10 mm × 15 mm were cut from the initial forged bar. Before testing, there is a solution treatment (T = 1040 °C, t = 0.75 h) for all specimens to obtain uniform initial microstructures, and then cooled by water. Then, the treated specimens were aged at
Time Fig. 1. Experimental procedure for tests.
Fig. 2. Microstructure of the studied superalloy; (a) and (b) show the microstructure at the region of P1 after deformation and annealing treatment, respectively; (c) and (d) show the microstructure at the region of P2 after deformation and annealing treatment, respectively. 213
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
°
°
°
°
Fig. 3. Orientation image microscopy maps and misorientation angle distributions of the deformed specimens at different annealing temperature and annealing time of 10 min (a, b) the deformed one; (c, d) 950 °C; (e, f) 980 °C; (g, h) 1010 °C (the LAGBs and HAGBs are indicated by thin-gray lines and thick-black lines, respectively).
214
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
(°C) Fig. 4. Grain size distribution maps of the deformed specimens at the annealing time of 10 min and annealing temperature of: (a) 950 °C; (b)980 °C; (c)1010 °C; (d) Histogram of average grain size.
900 °C for 24 h and then cooled by water again. The detailed experimental scheme is shown in Fig. 1. Hot compression tests were performed on a Gleeble-3500 at the deformation temperature of 950 °C and the strain rate of 0.1 s−1. Before loading, all specimens were heated to the deformation temperature at 10 °C/s, and then held for 1 min to eliminate the temperature gradient in specimen. The height reduction for all specimens is 50%. After hot compression test, all the deformed specimens were immediately quenched by water. Then the annealing tests were carried out in a heat treatment furnace. The annealing times are 5, 10, 30, 60 and 180 min. The annealing temperatures are 950, 965, 980 and 1010 °C. All the specimens were quenched by water after annealing heat. The microstructures were observed through OM, SEM and EBSD. For OM observations, the samples were firstly mirror polished, and then chemically etched with a corrosion solution of which compositions are HCl (50 ml) + CH3CH2OH (50 ml) + CuCl2 (2.5 g). For SEM observations, the samples were mirror polished. For EBSD observations, the slices were cut from the samples, and then the disks of 70–80 μm thick and 3 mm diameter were acquired. Finally, the disks were electro-polished, and the compositions of corrosive solution are HClO4 (30 ml) + CH3CH2OH (270 ml).
is only ~0.3. Fig. 2a and c show the microstructures at the region of P1 and P2 after deformation, respectively. It is easy to find that the microstructure is very different due to the uneven deformation. The grains at the region of P1 have been elongated, and a mass of fine DRX grains (marked by ellipse) distributed around the deformed grain boundaries, while there are only few DRX grains (marked by ellipse) at the region of P2, as shown the partial enlarged view in Fig. 2c. The DRX is not fully complete both on the regions of P1 and P2. It has been proved that the fraction of DRX increases with the increased temperature and strain or decreased strain rate [49–51]. This indicates that it needs a higher temperature and strain or lower strain rate to make sure a full completion of DRX, which is often difficult to achieve in die forging. However, the microstructure is similar due to the occurrence of recrystallization after annealing treatment, and both of them are fine and homogeneous, as shown in Fig. 2b and d. It indicates that adopting an annealing treatment after deformation can effectively improve the homogeneity of the microstructure. But the optimum annealing parameters that make the microstructure most uniform are still unknown. In order to obtain the optimum annealing parameters, it is necessary to study the influences of annealing parameters on microstructural evolution.
3. Results and Discussion
3.2. Microstructure Evolution
3.1. Feasibility to Obtain Homogeneous Microstructure by Annealing Treatment
3.2.1. Influences of Annealing Temperature Fig. 3 represents the microstructures of the deformed specimens (at the region of P1) annealed at various temperatures with the annealing time of 10 min. From Fig. 3, it can be found that the elongated grains have turned into small equiaxial grains due to the occurrence of static recrystallization. Additionally, there are some annealing twins in the microstructure. It can also be found that the annealing temperature has an obvious effect on the microstructural evolution. As the annealing
A verified experiment has been carried out to verify the feasibility to obtain homogeneous microstructure by annealing treatment. To obtain the distribution of true strain in the deformed specimen, the finite element simulation was done using Deform-3D, as shown in Fig. 2. The value of strain at the region of P1 is ~0.9, while that at the region of P2 215
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
°
°
Fig. 5. Orientation image microscopy maps and misorientation angle distributions of the deformed specimens at annealing temperature of 980 °C and annealing time of (a, b) 5 min (c, d)30 min (the LAGBs and HAGBs are indicated by thin-gray lines and thick-black lines, respectively).
be found that the necklace-like microstructures have turned into uniform and equiaxial grains after annealing treatment. There are two reasons. On the one hand, the DRX recrystallization nuclei formed in the deformation have grown up during the annealing process. On the other hand, a static recrystallization occurred in the original elongated grains where still contain a few strain and dislocation energies. The large elongated grains are gradually replaced by recrystallization one. Therefore, the final grains are uniform and equiaxial. When the annealing time is 5 min, there are still some deformed grains surrounded by a lot of fine grains, as shown in Fig. 5a. This is because the characteristics of nucleation mechanism are the bulging at grain boundary and the combination of subgrain [55]. There is an incubation period for nucleation before the occurrence of recrystallization during annealing treatment. When the time is short, some deformed grains are still in the incubation period. Therefore, there are still some original deformed grains. When the annealing time is 10 min, all the deformed grains have been replaced by fine grains. The grains are fine and homogeneous. There are even a few large grains and the average grain size increases, as shown in Fig. 6, when the annealing time is up to 30 min. The grains are not uniform any more, as shown in Fig. 5c. This is because there are some small unstable recrystallization nuclei during annealing treatment. In order to keep a stable structure, grain growth takes place by minishing the grain boundary area per unit volume [56]. Due to the different energy in each grain, there are big differences on the final grain size. Hence, the grains become heterogeneous again with the increase of time. Thus, it can be concluded that the recrystallized grains grow up gradually when the annealing time is increased from 5 to 30 min. Meanwhile, the initial heterogeneous grains become homogeneous firstly, and heterogeneous again.
temperature is increased from 950 °C to 1010 °C, the average grain size decreases firstly and then increases, as shown in the Fig. 4. The reasons are as follows. In general, the promoting effect on recrystallization becomes more and more obvious as the annealing temperature is increased. Meanwhile, it also improves the migration rate of grain boundaries by providing more energy [52]. However, for the studied alloy, the solution temperature of delta phase is 980 °C. In the process of annealing, the large long needle-like delta phases gradually become small spherical. Previous studies [53,54] have shown that the second phases with different size have different effects on recrystallization. There are high dislocation density areas around the second phase with large size, which accelerates the occurrences of recrystallization by increasing nucleation sites. The fine second phases dispersed and distributed in matrix have pinning effects on the migration of grain boundaries. So, it leads to the restriction on the growth of grains. The fraction of fine second phase at 980 °C is higher than that at 950 °C.·This is the reason why the average grain size decreases firstly. When the annealing temperature is increased to 1010 °C, the average grain size increases obviously, as shown in Fig. 4c. This is because the delta phases dissolve quickly at the temperature of 1010 °C. Therefore, it cannot inhibit the grain growth any more. From Fig. 4, it can also be found that the microstructures treated at the annealing temperature of 950 °C and 980 °C for 10 min are similar, but there are obvious differences in the fraction of delta phase, which will be discussed in Section 3.4.
3.2.2. Influences of Annealing Time Fig. 5 indicates the microstructure of the deformed specimens annealed at the temperature of 980 °C with different annealing time. It can 216
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
0.5
0.4
d=9.2µm
0.3 0.2
0.2 0.1
0.1 0.0
d=11.94µm
0.3 Fraction
Fraction
0.4
0.0 0
5 10 15 20 25 30 35 40 45 50
0
5
10
(a) 0.5
Average grain size (µm)
Fraction
0.3 0.2 0.1 10
20
30
40
50
25
25
d=22.51µm
0
20
30
35
(b)
0.4
0.0
15
Grain size(µm)
Grain size(µm)
60
70
80
22.51
20 15
11.94 9.2
10 5 0
0
5
Grain size(µm)
10
30
Annealing time (min)
(c)
(d)
Fig. 6. Grain size distribution maps of the deformed specimens at the annealing temperature of 980 °C, and the time of: (a) 5 min; (b) 10 min; (c) 30 min; (d) Histogram of average grain size.
direction of grain boundaries is hard to over 5°, while the cumulative misorientation is easy to over 15°, as shown in Fig. 7a and b. This illustrates there is a relatively steady misorientation gradient both in the parallel and vertical direction of grain boundaries. In addition, the large cumulative misorientation indicates that there is a large orientation gradient both in the parallel and vertical direction of grain boundaries, where the recrystallization grain boundaries can easily transform into HAGBs because it is easy to accumulate enough misorientations [59]. For the annealed specimens, it can be found that both the local and cumulative misorientation in the parallel and vertical direction of grain boundaries are hard to over 1°, as shown in Fig. 7c and d. It shows that the orientation gradient and little deformation energy in the both directions are little. Similar phenomenon can be found from Fig. 7e–h. It indicates that the static recrystallization is fully completed, and all of the large elongated deformed grains are replaced by recrystallization grains. The appearance of large grains shown in Fig. 3g is due to the growth of recrystallization grains rather than the residue of the deformed grains.
3.3. Misorientation Angle Evolution 3.3.1. Influences of Annealing Temperature Fig. 3 shows the orientation image microscopy (OIM) maps of the deformed specimens and the specimens annealed at various annealing temperatures, respectively. It is easy to find that the microstructure after deformation is heterogeneous. Most of grain boundaries still belong to LAGBs, as shown in Fig. 3a marked by thin-gray lines. The fraction of LAGBs is relatively high and the average misorientation angle is 11.3°, as shown in Fig. 3b. It indicates that the fraction of recrystallization is very low after deformation. Compared to deformed specimens, the fraction of LAGBs after annealing treatment is significantly low, and the average misorientation angle increases obviously, as shown in Fig. 3. Additionally, the fraction of 60° boundaries increases largely after the annealing treatment. It shows that there are some annealing twins formed during annealing treatment. Generally, the boundary energy of growing grain is decreased by annealing twins, while the mobility of grain boundary is enhanced [57].Therefore, the annealing twins can accelerate the process of static recrystallization. When the annealing temperature is increased from 950 to 1010 °C, the fraction of LAGBs decreases and the average misorientation angle increases slightly. It indicates that the fraction of recrystallization for the latter is higher. This is because the migration capacity of the grain boundaries can be improved by elevating annealing temperature, which promotes the recrystallization. That is to say the recrystallization is easy to occur at higher temperatures [58]. The fraction of LAGBs decreases, and the fraction of recrystallization and the average misorientation angle increases with the increased annealing temperature. The point to point (local) misorientation and point to origin (cumulative) misorientations along the lines marked in Fig. 3 are calculated, as shown in Fig. 7. For the deformed specimens, it is easy to discover that the local misorientation both in the parallel and vertical
3.3.2. Influences of Annealing Time Fig. 5 shows the OIM maps and misorientation angle distributions of the deformed specimens at the annealing temperature of 980 °C with different annealing times. Compared to the initial deformed specimens, the specimens after annealing treatment at 980 °C for 5 min is dominated by HAGBs, and the fraction of LAGBs is extremely low. Meanwhile, the average misorientation angle is 48.5°. As the annealing time is further increased to 30 min, it is almost impossible to observe LAGBs again, and the average misorientation angle increases to 50.8°. It can be concluded that the fraction of LAGBs decreases and the average misorientation angle increases with the increasing annealing time. It is because there are many small recrystallization nuclei in the deformed specimens, and there is a certain strain energy accumulated in the 217
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
Point to origin Point to point
16 12 8 4 0
0
10
20
30
40
Point to origin Point to point
30 Misorientation angle (°)
Misorientation angle (°)
20
25 20 15 10 5 0
50
0
5
10
Distance (µm)
(b) Point to origin Point to point
0.3 0.2 0.1
5
10
15
20
25
30
Point to origin Point to point
0.4 Misorientation angle (°)
Misorientation angle (°)
0.4
0
0.3 0.2 0.1 0.0
35
0
2
4
Distance (µm)
0.2
0.1
10
15
20
0.2 0.1 0.0
25
0
5
10
Point to origin Point to point
Misorientation angle (°)
Misorientation angle (°)
0.30
0.20 0.15 0.10 0.05 0
5
20
(f)
0.25
0.00
15
Distance (µm)
(e) 0.30
12
0.3
Distance (µm)
0.35
10
Point to origin Point to point
0.4 Misorientation angle (°)
Misorientation angle (°)
Point to origin Point to point
5
8
(d)
0.3
0
6
Distance (µm)
(c)
0.0
20
Distance (µm)
(a)
0.0
15
0.20 0.15 0.10 0.05 0.00
10 15 20 25 30 35 40 45
Point to origin Point to point
0.25
0
5
Distance (µm)
10
15
20
Distance (µm)
(g)
(h)
Fig. 7. Changes of misorientation angle along the lines marked in Fig. 3: (a) A1; (b) A2; (c) B1; (d) B2; (e) C1; (f) C2; (g) D1; (h) D2.
annealing treatment at 980 °C for 5 min are higher than that for 10 min and 30 min, as shown in Fig. 8. It indicates there is still a little orientation gradient in both directions. The recrystallization is still not fully complete. Therefore, the large grains shown in Fig. 5a are the residue of the deformed grains. It illustrates that more time is needed for the occurrence of recrystallization. With the increase of annealing time, the local and cumulative misorientations in parallel and vertical directions of grain boundaries are little. It indicates there are only small orientation gradients in both directions, and the deformation energy
deformation zone without complete recrystallization. Therefore, the new recrystallization nuclei grew up in the process of annealing, as well as the DRX nuclei formed in the deformation, which cause the LAGBs change to HAGBs. There is more time for recrystallization with the increase of annealing time. So, the fraction of recrystallization increases. This causes the fraction of LAGBs to decrease and the average misorientation angle to increase. Meanwhile, the cumulative misorientations in the parallel and vertical directions of grain boundaries in the specimen after the 218
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
3.5
Point to origin Point to point
3.0
Misorientation angle (°)
Misorientation angle (°)
3.5 2.5 2.0 1.5 1.0 0.5 0.0
0
10
20
30
40
50
60
2.5 2.0 1.5 1.0 0.5 0.0
70
Point to origin Point to point
3.0
0
5
Distance (µm)
10
Misorientation angle (°)
Misorientation angle (°)
(b) 0.6
Point to origin Point to point
1.0 0.8 0.6 0.4 0.2 0.0
0
10
20
30
20
Distance (µm)
(a) 1.2
15
40
50
0.4 0.3 0.2 0.1 0.0
60
Distance (µm)
Point to origin Point to point
0.5
0
5
10
15
20
25
30
35
Distance (µm)
(c)
(d)
Fig. 8. Changes of misorientation angle along the lines marked in Fig. 5: (a) E1; (b) E2; (c) F1; (d) F2.
Fig. 9. SEM micrographs of the deformed specimens at the annealing time of 60 min, and the temperature of: (a)950 °C; (b) 965 °C; (c) 980 °C.
3.4. Delta Phase Evolution
has been depleted. Therefore, it is difficult to migrate for grain boundaries by dislocation slip or climb [52]. The grain growth behavior, as shown in Fig. 5c, is mainly caused by the HAGBs migration triggered by atom motion [60]. In addition, it should be pointed out that there is seemingly a slight increase in the cumulative misorientations inside the recrystallization grains when the annealing time is increased from 10 to 30 min, as shown in Figs. 7e–f and 8c–d. However, the increase is not caused by the increase of annealing time. It only attributes to the wave of misorientation inside each recrystallization grain. The difference between the values shown in Figs. 7e–f and 8c–d results from the uncertainty of the position of underlines.
3.4.1. Influences of Annealing Temperature Fig. 9 shows the SEM micrographs of the deformed samples annealed at different temperatures with the annealing time of 60 min. The needle-like delta phases become rod-like and spherical during annealing treatment, as shown in Fig. 9. The fraction of needle-like delta phase decreases with the increase of temperature. When the temperature reaches 980 °C, all the needle-like delta phases are dissolved. The change of morphology in delta phase is a phase boundary migration process, which is controlled by the atomic diffusion. There is more energy for the atomic diffusion in high temperature. Therefore, the proportion of needle-like delta phase decreases with the increase of temperature. The effect of annealing temperature on the fraction of delta phase is shown in Fig. 10. It can be found that the fraction of delta 219
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
The delta phase which is perpendicular to the direction of compression is subjected to tensile stress, which leads to the phenomenon of tension fracture during deformation. While the delta phases distributed along the direction of compression are bent or broken to adapt to the deformation. Thus, there are a lot of delta phases that have been broken after deformation. The morphology of these delta phases has changed from needle-like to rod-like and spherical. It means that the decomposition of these phase have finished in deformation. During the annealing, these delta phases will dissolve with a high rate, which lead to the dissolution rate being higher than the precipitation one. Therefore, the fraction of delta phase decreases at the early stage of annealing. It indicates that the deformation accelerates the dissolution of delta phase. However, all of these delta phases have dissolved as the annealing time further increases. Thus, the precipitation becomes stronger than dissolution. So, the fraction of delta phase increases. Although the fraction of delta phase at 30 min is lower than the one at 10 min, the grains are heterogeneous after annealing at 980 °C for 30 min, as shown in Fig. 5c. All things considered, the annealing time should be chosen as 10 min.
Fig. 10. Effects of annealing temperature on the fraction of delta phase (The annealing time is 60 min).
phase decreases from 3.9% to 0.52% when the annealing temperature is increased from 950 to 980 °C. This is because the annealing temperature has a decisive effect on the diffusion process of delta phase. With the increased temperature, the dissolution rate increases owing to the high diffusion coefficient at high temperature [61]. The excessive content of delta phase can promote crack propagation rate, and reduce the fatigue life of workpiece [35].Therefore, both considering the fatigue life and the uniformity of microstructure, the annealing temperature should be chosen as 980 °C.
4. Conclusions (1) It has been found that adopting an annealing treatment after deformation can effectively improve the homogeneity of the microstructure. The optimum parameters are the annealing temperature of 980 °C and time of 10 min. (2) The grains become homogeneous due to the occurrence of static recrystallization. There is an incubation period which is not less than 5 min during the static recrystallization. The initial heterogeneous grains become homogeneous firstly, and then heterogeneous again when the annealing time is increased from 5 to 30 min. (3) The deformation accelerates the dissolution of delta phase in annealing treatment. The fraction of delta phase decreases at the early stage of annealing whether the annealing temperature exceeds the solution temperature.
Acknowledgements This work was supported by Hunan Provincial Natural Science Foundation of China (2017JJ3380), National Natural Science Foundation of China (No. 51775564), the Science and Technology Leading Talent in Hunan Province (No. 2016RS2006), Program of Chang Jiang Scholars of Ministry of Education (No. Q2015140), the Open-End Fund for the Valuable and Precision Instruments of Central South University (No. CSUZC201821)，and Hebei Iron and Steel Joint Funds (No. E2015209243).
7
7
6
6 Fraction (%)
Fraction (%)
3.4.2. Influences of Annealing Time Figs. 11 and 12 show the effects of annealing time on the fraction and morphology of delta phase, respectively. It can be seen that the annealing time has a great effect on the fraction of delta phase. As the annealing time is increased, the fraction of delta phase decreases gradually at the annealing temperature of 980 °C. The fraction of delta phase decreases from 6.8% to 0.52% when the annealing time is increased to 60 min, as shown in Fig. 11(a). Meanwhile, the needle-like delta phase becomes spherical, as shown in Fig. 12. However, at the annealing temperature of 950 °C which is lower than the solution temperature of delta phase, the fraction of delta phase firstly decreases and then increases with the increasing annealing time. The reasons are as follows. According to the existing literature [61,62], the static dissolution of delta phase mainly consists of three parts: (1) the decomposition of delta phase; (2) the short-range diffusional process of atoms; (3) a longer-range diffusional process of atoms. But, the time for the decomposition of delta phase is much more than the short-range diffusional process of atoms and longer-range diffusional process of atoms.
5 4 3 2 1 0
5 4 3 2 1
0
10
20
30
40
50
0
60
Annealing time (min)
(a)
0
10 60 Annealing time (min)
180
(b)
Fig. 11. Effects of annealing time on the fraction of delta phase at the annealing temperature of: (a) 980 °C; (b) 950 °C. 220
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
Fig. 12. SEM micrographs of the deformed specimens at the annealing temperature of 980 °C, and the time of: (a) 10 min; (b) 30 min.
References
[20] Y. Wang, Q. Pan, Y. Song, C. Li, Z. Li, Hot deformation and processing maps of X750 nickel-based superalloy, Mater. Des. 1 (2013) 154–160. [21] P.M. Mignanelli, N.G. Jones, K.M. Perkins, M.C. Hardy, H.J. Stone, Microstructural evolution of a delta containing nickel-base superalloy during heat treatment and isothermal forging, Mater. Sci. Eng. A 621 (2015) 265–271. [22] M. Retima, S. Bouyegh, H. Chadli, Effect of the heat treatment on the microstructural evolution of the nickel based superalloy, Meta 17 (2011) 71–77. [23] X.M. Chen, Y.C. Lin, D.X. Wen, J.L. Zhang, M. He, Dynamic recrystallization behavior of a typical nickel-based superalloy during hot deformation, Mater. Des. 57 (2014) 568–577. [24] M.S. Chen, Y.C. Lin, K.K. Li, Y. Zhou, A new method to establish dynamic recrystallization kinetics model of a typical solution-treated Ni-based superalloy, Comput. Mater. Sci. 122 (2016) 150–158. [25] Y.X. Liu, Y.C. Lin, Y. Zhou, 2D cellular automaton simulation of hot deformation behavior in a Ni-based superalloy under varying thermal-mechanical conditions, Mater. Sci. Eng. A 691 (2017) 88–99. [26] L.A. Reyes, P. Páramo, A. Salas Zamarripa, M. de la Garza, M.P. Guerrero Mata, Grain size modeling of a Ni-base superalloy using cellular automata algorithm, Mater. Des. 83 (2015) 301–307. [27] T.S. Prithiv, P. Bhuyan, S.K. Pradhan, V.S. Sarma, S. Mandal, A critical evaluation on efficacy of recrystallization vs. strain induced boundary migration in achieving grain boundary engineered microstructure in a Ni-base superalloy, Acta Mater. 146 (2018) 187–201. [28] C. Zhang, L. Zhang, M. Li, W. Shen, S. Gu, Effects of microstructure and γ′ distribution on the hot deformation behavior for a powder metallurgy superalloy FGH96, J. Mater. Res. 29 (2014) 2799–2808. [29] Y.H. Liu, Y.Q. Ning, X.M. Yang, Z.K. Yao, H.Z. Guo, Effect of temperature and strain rate on the workability of FGH4096 superalloy in hot deformation, Mater. Des. 95 (2016) 669–676. [30] J.T. Yeom, C.S. Lee, J.H. Kim, N.K. Park, Finite-element analysis of microstructure evolution in the cogging of an alloy 718 ingot, Mater. Sci. Eng. A 449 (2007) 722–726. [31] A.J. Brand, K. Karhausen, R. Kopp, Microstructural simulation of nickel base alloy inconel 718 in production of turbine discs, Mater. Sci. Technol. 12 (2014) 963–969. [32] Y.C. Lin, X.M. Chen, A critical review of experimental results and constitutive descriptions for metals and alloys in hot working, Mater. Des. 32 (2011) 1733–1759. [33] X. You, Y. Tan, J. Li, P. Li, C. Dong, S. Shi, S. Qin, Effects of solution heat treatment on the microstructure and hardness of inconel 740 superalloy prepared by electron beam smelting, J. Alloys Compd. 638 (2015) 239–248. [34] A. Ghosh, N.P. Gurao, Effect of crystallographic texture on ratcheting response of commercially pure titanium, Mater. Des. 115 (2017) 121–132. [35] J. An, L. Wang, Y. Liu, W. Cai, X. Song, The role of δ phase for fatigue crack propagation behavior in a Ni base superalloy at room temperature, Mater. Sci. Eng. A 684 (2017) 312–317. [36] D. Hu, Q. Yang, H. Liu, J. Mao, F. Meng, Y. Wang, R. Wang, Crack closure effect and crack growth behavior in GH2036 superalloy plates under combined high and low cycle fatigue, Int. J. Fatigue 95 (2017) 90–103. [37] G.Z. Quan, Y.L. Li, L. Zhang, X. Wang, Evolution of grain refinement degree induced by dynamic recrystallization for Nimonic 80A during hot compression process and its FEM analysis, Vacuum 139 (2017) 51–63. [38] J. Kang, D.S. Wilkinson, R.K. Mishra, W. Yuan, R.S. Mishra, Effect of inhomogeneous deformation on anisotropy of AZ31 magnesium sheet, Mater. Sci. Eng. A 567 (2013) 101–109. [39] Z. Sun, H. Yang, N. Sun, Effects of parameters on inhomogeneous deformation and damage in isothermal local loading forming of Ti-alloy component, J. Mater. Eng. Perform. 21 (2012) 313–323.
[1] H.Y. Li, J.D. Hu, D.D. Wei, X.F. Wang, Y.H. Li, Artificial neural network and constitutive equations to predict the hot deformation behavior of modified 2.25 Cr–1Mo steel, Mater. Des. 42 (2012) 192–197. [2] D.X. Wen, Y.C. Lin, H.B. Li, X.M. Chen, J. Deng, L.T. Li, Hot deformation behavior and processing map of a typical Ni-based superalloy, Mater. Sci. Eng. A 591 (2014) 183–192. [3] M.R. Jahangiri, H. Arabi, S.M.A. Boutorabi, Comparison of microstructural stability of IN939 superalloy with two different manufacturing routes during long-time aging, Trans. Nonferrous Metals Soc. China 24 (2014) 1717–1729. [4] Z.H. Yao, Q.Y. Wang, M.C. Zhang, J.X. Dong, Microstructure control and prediction of GH738 superalloy during hot deformation, Acta Metall. Sin. 47 (2011) 1591–1599. [5] Y.C. Lin, D.X. Wen, J. Deng, G. Liu, J. Chen, Constitutive models for high-temperature flow behaviors of a Ni-based superalloy, Mater. Des. 59( (2014) 115–123. [6] Y. Liu, Z. Yao, Y. Ning, Y. Nan, H. Guo, C. Qin, Z. Shi, The flow behavior and constitutive equation in isothermal compression of FGH4096-GH4133B dual alloy, Mater. Des. 63( (2014) 829–837. [7] F.M. Shore, M. Morakabati, S.M. Abbasi, A. Momeni, Hot deformation behavior of incoloy 901 through hot tensile testing, J. Mater. Eng. Perform. 23 (2014) 1424–1433. [8] Y.C. Lin, K.K. Li, H.B. Li, J. Chen, X.M. Chen, D.X. Wen, New constitutive model for high-temperature deformation behavior of inconel 718 superalloy, Mater. Des. 74( (2015) 108–118. [9] Y.C. Lin, X.M. Chen, D.X. Wen, M.S. Chen, A physically-based constitutive model for a typical nickel-based superalloy, Comput. Mater. Sci. 83 (2014) 282–289. [10] Y. Ning, Z. Yao, X. Liang, Y. Liu, Flow behavior and constitutive model for Ni–20.0 Cr–2.5 Ti–1.5 Nb–1.0 Al superalloy compressed below γ′-transus temperature, Mater. Sci. Eng. A 551 (2012) 7–12. [11] P. Zhang, C. Hu, Q. Zhu, C.G. Ding, H.Y. Qin, Hot compression deformation and constitutive modeling of GH4698 alloy, Mater. Des. 65( (2015) 1153–1160. [12] H.Y. Wu, F.J. Zhu, S.C. Wang, W.R. Wang, C.C. Wang, C.H. Chiu, Hot deformation characteristics and strain-dependent constitutive analysis of inconel 600 superalloy, J. Mater. Sci. 47 (2012) 3971–3981. [13] A. Etaati, K. Dehghani, G.R. Ebrahimi, H. Wang, Predicting the flow stress behavior of Ni-42.5 Ti-3Cu during hot deformation using constitutive equations, Met. Mater. Int. 19 (2013) 5–9. [14] Q. Yu, J. Liu, W. Zhang, L. Yu, F. Liu, W. Sun, Modification of Johnson-Cook model for hot deformation behavior of inconel 625 alloy, Rare Metal Mater. Eng. 41 (2013) 1679–1684. [15] A. Momeni, S.M. Abbasi, M. Morakabati, H. Badri, X. Wang, Dynamic recrystallization behavior and constitutive analysis of Incoloy 901 under hot working condition, Mater. Sci. Eng. A 615( (2014) 51–60. [16] D.X. Wen, Y.C. Lin, J. Chen, J. Deng, X.M. Chen, J.L. Zhang, M. He, Effects of initial aging time on processing map and microstructures of a nickel-based superalloy, Mater. Sci. Eng. A 620 (2015) 319–332. [17] P. Zhang, C. Hu, C.G. Ding, Q. Zhu, H.Y. Qin, Plastic deformation behavior and processing maps of a Ni-based superalloy, Mater. Des. 65 (2015) 575–584. [18] Y. Liu, Y. Ning, Z. Yao, H. Guo, Y. Nan, Effect of true strains on processing map for isothermal compression of Ni–20.0 Cr–2.5 Ti–1.5 Nb–1.0 Al Ni-base superalloy, J. Alloys Compd. 612 (2014) 56–63. [19] D.G. He, Y.C. Lin, M.S. Chen, J. Chen, D.X. Wen, X.M. Chen, Effect of pre-treatment on hot deformation behavior and processing map of an aged nickel-based superalloy, J. Alloys Compd. 649 (2015) 1075–1084.
221
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
[51] Y.C. Lin, D.G. He, M.S. Chen, X.M. Chen, C.Y. Zhao, X. Ma, Z.L. Long, EBSD analysis of evolution of dynamic recrystallization grains and δ phase in a nickel-based superalloy during hot compressive deformation, Mater. Des. 97 (2016) 13–24. [52] X.M. Chen, Y.C. Lin, F. Wu, EBSD study of grain growth behavior and annealing twin evolution after full recrystallization in a nickel-based superalloy, J. Alloys Compd. 724 (2017) 198–207. [53] R.L. Goetz, Particle stimulated nucleation during dynamic recrystallization using a cellular automata model, Scr. Mater. 52 (2005) 851–856. [54] F.J. Humphreys, Particle stimulated nucleation of recrystallization at silica particles in nickel, Scr. Mater. 43 (2000) 591–596. [55] X. Wei, W. Zheng, Z. Song, T. Lei, Q. Yong, Q. Xie, Static recrystallization behavior of inconel 718 alloy during thermal deformation, J. Wuhan Univ. Technol. 29 (2014) 379–383. [56] Y.S. Jang, D.C. Ko, B.M. Kim, Application of the finite element method to predict microstructure evolution in the hot forging of steel, J. Mater. Process. Technol. 101 (2000) 85–94. [57] S.P. Coryell, K.O. Findley, M.C. Mataya, E. Brown, Evolution of microstructure and texture during hot compression of a Ni-Fe-Cr superalloy, Metall. Mater. Trans. A 43 (2012) 633–649. [58] X.M. Chen, Y.C. Lin, M.S. Chen, H.B. Li, D.X. Wen, J.L. Zhang, Microstructural evolution of a nickel-based superalloy during hot deformation, Mater. Des. 77 (2015) 41–49. [59] H. Zhang, K. Zhang, S. Jiang, H. Zhou, C. Zhao, X. Yang, Dynamic recrystallization behavior of a γ′-hardened nickel-based superalloy during hot deformation, J. Alloys Compd. 623 (2015) 374–385. [60] V. Randle, P.R. Rios, Y. Hu, Grain growth and twinning in nickel, Scr. Mater. 58 (2008) 130–133. [61] D. Cai, W. Zhang, P. Nie, W. Liu, M. Yao, Dissolution kinetics of δ phase and its influence on the notch sensitivity of inconel 718, Mater. Charact. 58 (2007) 220–225. [62] V. Beaubois, J. Huez, S. Coste, O. Brucelle, J. Lacaze, Short term precipitation kinetics of delta phase in strain free Inconel 718 alloy, Mater. Sci. Technol. 20 (2004) 1019–1026.
[40] H. Mirzadeh, M.H. Parsa, Hot deformation and dynamic recrystallization of NiTi intermetallic compound, J. Alloys Compd. 614 (2014) 56–59. [41] S. Saadatkia, H. Mirzadeh, J. Cabrera, Hot deformation behavior, dynamic recrystallization, and physically-based constitutive modeling of plain carbon steels, Mater. Sci. Eng. A 636 (2015) 196–202. [42] G.R. Ebrahimi, A. Momeni, H.R. Ezatpour, Modeling the viscoplastic behavior and grain size in a hot worked Nb-bearing high-Mn steel, Mater. Sci. Eng. A 714 (2018) 25–35. [43] C. Zhang, L.W. Zhang, W.F. Shen, Q.H. Xu, Y. Cui, The processing map and microstructure evolution of Ni-Cr-Mo-based C276 superalloy during hot compression, J. Alloys Compd. 728 (2017) 1269–1278. [44] Y.Q. Ning, C. Zhou, H.Q. Liang, M.W. Fu, Abnormal flow behavior and necklace microstructure of powder metallurgy superalloys with previous particle boundaries (PPBs), Mater. Sci. Eng. A 652 (2016) 84–91. [45] S.K. Pradhan, S. Mandal, C.N. Athreya, K.A. Babu, B. de Boer, V.S. Sarma, Influence of processing parameters on dynamic recrystallization and the associated annealing twin boundary evolution in a nickel base superalloy, Mater. Sci. Eng. A 700 (2017) 49–58. [46] X.F. Tang, B.Y. Wang, H.C. Ji, X.B. Fu, W.C. Xiao, Behavior and modeling of microstructure evolution during metadynamic recrystallization of a Ni-based superalloy, Mater. Sci. Eng. A 675 (2016) 192–203. [47] G.Z. Quan, J. Pan, X. Wang, Z.H. Zhang, L. Zhang, T. Wang, Correspondence between grain refinements and flow softening behaviors at Nimonic 80A superalloy under different strain rates, temperatures and strains, Mater. Sci. Eng. A 679 (2017) 358–371. [48] A. Sarkar, S. Sanyal, T.K. Bandyopadhyay, S. Mandal, Influence of annealing parameters on phase evolution and recrystallization kinetics of a Mn-Al-Si alloyed duplex steel, Mater. Charact. 134 (2017) 213–224. [49] Y.C. Lin, X.Y. Wu, X.M. Chen, J. Chen, D.X. Wen, J.L. Zhang, L.T. Li, EBSD study of a hot deformed nickel-based superalloy, J. Alloys Compd. 640 (2015) 101–113. [50] H. Jiang, J. Dong, M. Zhang, Z. Yao, Evolution of twins and substructures during low strain rate hot deformation and contribution to dynamic recrystallization in alloy 617B, Mater. Sci. Eng. A 649 (2016) 369–381.
222
Contents lists available at ScienceDirect
Materials Characterization journal homepage: www.elsevier.com/locate/matchar
Effects of annealing parameters on microstructural evolution of a typical nickel-based superalloy during annealing treatment ⁎
T
⁎
Ming-Song Chena,b, , Zong-Huai Zoua,b, Y.C. Lina,b,c, , Hong-Bin Lid, Wu-Quan Yuana,b a
School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China State Key Laboratory of High Performance Complex Manufacturing, Changsha 410083, China c Light Alloy Research Institute of Central South University, Changsha 410083, China d College of Metallurgy and Energy, Hebei United University, Tangshan 063009, China b
A R T I C LE I N FO
A B S T R A C T
Keywords: Superalloy Annealing treatment Microstructure Homogeneity
It is usually one of important goals to obtain the homogeneous microstructure for forgings by hot working, such as hot die forging. However, there is always a large inhomogeneous deformation in die forgings during hot deformation. It often results in an inhomogeneous microstructure of forging due to the different dynamic recrystallization (DRX) fractions and grain sizes in different strain regions. In this study, a method was proposed to improve the homogeneity of deformed microstructure by annealing treatment. The microstructural evolution of a typical nickel-based superalloy during annealing treatment has been investigated. The Optical Microscope (OM), Electron Backscatter Diffraction (EBSD) and Scanning Electron Microscope (SEM) technologies were applied to observe microstructures. It is found that adopting an annealing treatment after deformation can effectively improve the homogeneity of microstructure due to the static recrystallization. There is an incubation period which is not less than 5 min for the occurrence of static recrystallization. The initial heterogeneous grains become homogeneous firstly, and then heterogeneous again when the annealing time is increased from 5 to 30 min. The deformation accelerates the dissolution of delta phase in annealing treatment. The fraction of delta phase decreases at the early stage of annealing whether the temperature exceeds the solution temperature. Based on a comprehensive consideration of all factors, the optimum annealing parameters, which are the temperature of 980 °C and the time of 10 min, have been obtained.
1. Introduction
established a DRX kinetics model by a new method. CA models were established by Liu et al. [25] and Reyes et al. [26]. Prithiv et al. [27] evaluated the efficacy of recrystallization vs. strain induced boundary migration in achieving grain boundary engineered microstructure in a Ni-based superalloy. In addition, the hot deformation behavior of superalloy FGH96 was investigated by Zhang et al. [28] and Liu et al. [29]. Yeom et al. [30] and Brand et al. [31] simulated the microstructural evolution by finite-element (FE) method. The properties of parts mainly depend on their microstructures [32,33]. Generally, fine and homogeneous microstructures can effectively improve the fatigue resistance and mechanical strength [34–36]. However, the microstructure after deformation is usually heterogeneous due to the inhomogeneous deformation [37]. It has been found that there is a severe inhomogeneous deformation on the sheet surface during uniaxial tensile testing [38]. Inhomogeneous deformation is determined by thermal deformation parameters such as local loading parameters, friction and die-forging mode [39]. It is difficult to find the perfect thermal deformation parameters that can avoid the
Nickel-based superalloys are applied to make the key parts in energy and aerospace industries due to their fine corrosion and fatigue resistance, as well as prominent high-temperature characteristics [1–3]. These key parts, such as turbine disk, often need to be formed by hot die forging. To make a proper hot working processing, some studies on the flow behaviors as well as microstructural evolution of Ni-based superalloys have been done [4–7]. For example, Lin et al. [8,9], Ning et al. [10], Zhang et al. [11], Wu et al. [12], Etaati et al. [13], Yu et al. [14] and Momeni et al. [15] established the constitutive models to accurately predict flow behaviors. Wen et al. [2,16], Zhang et al. [17], Liu et al. [18], He et al. [19] and Wang et al. [20] established the relevant processing maps to gain optimal hot deformation parameter domains. Mignanelli et al. [21] and Retima et al. [22] analyzed the effects of heat treatment on microstructural evolution. To predict the DRX behavior, Chen et al. [23] researched the DRX behavior through thermal compressive tests, and established a segmented model. Chen et al. [24]
⁎
Corresponding authors at: School of Mechanical and Electrical Engineering, Central South University, Changsha 410083, China. E-mail addresses: [email protected] (M.-S. Chen), [email protected] (Y.C. Lin).
https://doi.org/10.1016/j.matchar.2018.04.056 Received 10 February 2018; Received in revised form 27 April 2018; Accepted 30 April 2018 Available online 01 May 2018 1044-5803/ © 2018 Elsevier Inc. All rights reserved.
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
inhomogeneous deformation. In other words, the inhomogeneous deformation is unavoidable during die forging. It is well known that the fraction and grain size of DRX depend on the temperature, strain and strain rate [40–43]. Therefore, the DRX fraction and grain size in different regions of forging with different strains are also different, which often results in the inhomogeneous microstructure in forgings. So, it is necessary to adopt a method to obtain fine and homogeneous microstructures. However, few attentions focus on the way to improve the homogeneity of deformed microstructure, although there are a lot of researches focus on the flow behaviors and microstructural evolution [44–47]. In this study, a method to improve the homogeneity of deformed microstructure by annealing treatment was proposed because of the great influence of annealing parameters on phase evolution and recrystallization behavior [48]. Firstly, the feasibility to obtain homogeneous microstructure by annealing treatment has been verified. Secondly, the effects of annealing parameters on microstructural evolution have been studied to obtain the optimum annealing parameters.
Table 1 The chemical compositions (wt%) of nickel-based superalloy GH4169. Ni
Cr
Nb
Mo
Ti
Al
Co
C
Fe
52.82
18.96
5.23
3.01
1.00
0.59
0.01
0.03
Bal
Holding
Temperature
1 min Compression
T: 950, 965,980, 1010 C T: 950 C : 0.1s
t: 5,10,30,60,180min
-1
: 0.69 Heating rate 10 C/s
Water quenching
Heating rate 10 C/s
Water quenching
2. Materials and Experiments A commercial nickel-based superalloy of which chemical compositions are shown in Table 1 was used. The experimental cylindrical specimens with a size of ϕ10 mm × 15 mm were cut from the initial forged bar. Before testing, there is a solution treatment (T = 1040 °C, t = 0.75 h) for all specimens to obtain uniform initial microstructures, and then cooled by water. Then, the treated specimens were aged at
Time Fig. 1. Experimental procedure for tests.
Fig. 2. Microstructure of the studied superalloy; (a) and (b) show the microstructure at the region of P1 after deformation and annealing treatment, respectively; (c) and (d) show the microstructure at the region of P2 after deformation and annealing treatment, respectively. 213
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
°
°
°
°
Fig. 3. Orientation image microscopy maps and misorientation angle distributions of the deformed specimens at different annealing temperature and annealing time of 10 min (a, b) the deformed one; (c, d) 950 °C; (e, f) 980 °C; (g, h) 1010 °C (the LAGBs and HAGBs are indicated by thin-gray lines and thick-black lines, respectively).
214
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
(°C) Fig. 4. Grain size distribution maps of the deformed specimens at the annealing time of 10 min and annealing temperature of: (a) 950 °C; (b)980 °C; (c)1010 °C; (d) Histogram of average grain size.
900 °C for 24 h and then cooled by water again. The detailed experimental scheme is shown in Fig. 1. Hot compression tests were performed on a Gleeble-3500 at the deformation temperature of 950 °C and the strain rate of 0.1 s−1. Before loading, all specimens were heated to the deformation temperature at 10 °C/s, and then held for 1 min to eliminate the temperature gradient in specimen. The height reduction for all specimens is 50%. After hot compression test, all the deformed specimens were immediately quenched by water. Then the annealing tests were carried out in a heat treatment furnace. The annealing times are 5, 10, 30, 60 and 180 min. The annealing temperatures are 950, 965, 980 and 1010 °C. All the specimens were quenched by water after annealing heat. The microstructures were observed through OM, SEM and EBSD. For OM observations, the samples were firstly mirror polished, and then chemically etched with a corrosion solution of which compositions are HCl (50 ml) + CH3CH2OH (50 ml) + CuCl2 (2.5 g). For SEM observations, the samples were mirror polished. For EBSD observations, the slices were cut from the samples, and then the disks of 70–80 μm thick and 3 mm diameter were acquired. Finally, the disks were electro-polished, and the compositions of corrosive solution are HClO4 (30 ml) + CH3CH2OH (270 ml).
is only ~0.3. Fig. 2a and c show the microstructures at the region of P1 and P2 after deformation, respectively. It is easy to find that the microstructure is very different due to the uneven deformation. The grains at the region of P1 have been elongated, and a mass of fine DRX grains (marked by ellipse) distributed around the deformed grain boundaries, while there are only few DRX grains (marked by ellipse) at the region of P2, as shown the partial enlarged view in Fig. 2c. The DRX is not fully complete both on the regions of P1 and P2. It has been proved that the fraction of DRX increases with the increased temperature and strain or decreased strain rate [49–51]. This indicates that it needs a higher temperature and strain or lower strain rate to make sure a full completion of DRX, which is often difficult to achieve in die forging. However, the microstructure is similar due to the occurrence of recrystallization after annealing treatment, and both of them are fine and homogeneous, as shown in Fig. 2b and d. It indicates that adopting an annealing treatment after deformation can effectively improve the homogeneity of the microstructure. But the optimum annealing parameters that make the microstructure most uniform are still unknown. In order to obtain the optimum annealing parameters, it is necessary to study the influences of annealing parameters on microstructural evolution.
3. Results and Discussion
3.2. Microstructure Evolution
3.1. Feasibility to Obtain Homogeneous Microstructure by Annealing Treatment
3.2.1. Influences of Annealing Temperature Fig. 3 represents the microstructures of the deformed specimens (at the region of P1) annealed at various temperatures with the annealing time of 10 min. From Fig. 3, it can be found that the elongated grains have turned into small equiaxial grains due to the occurrence of static recrystallization. Additionally, there are some annealing twins in the microstructure. It can also be found that the annealing temperature has an obvious effect on the microstructural evolution. As the annealing
A verified experiment has been carried out to verify the feasibility to obtain homogeneous microstructure by annealing treatment. To obtain the distribution of true strain in the deformed specimen, the finite element simulation was done using Deform-3D, as shown in Fig. 2. The value of strain at the region of P1 is ~0.9, while that at the region of P2 215
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
°
°
Fig. 5. Orientation image microscopy maps and misorientation angle distributions of the deformed specimens at annealing temperature of 980 °C and annealing time of (a, b) 5 min (c, d)30 min (the LAGBs and HAGBs are indicated by thin-gray lines and thick-black lines, respectively).
be found that the necklace-like microstructures have turned into uniform and equiaxial grains after annealing treatment. There are two reasons. On the one hand, the DRX recrystallization nuclei formed in the deformation have grown up during the annealing process. On the other hand, a static recrystallization occurred in the original elongated grains where still contain a few strain and dislocation energies. The large elongated grains are gradually replaced by recrystallization one. Therefore, the final grains are uniform and equiaxial. When the annealing time is 5 min, there are still some deformed grains surrounded by a lot of fine grains, as shown in Fig. 5a. This is because the characteristics of nucleation mechanism are the bulging at grain boundary and the combination of subgrain [55]. There is an incubation period for nucleation before the occurrence of recrystallization during annealing treatment. When the time is short, some deformed grains are still in the incubation period. Therefore, there are still some original deformed grains. When the annealing time is 10 min, all the deformed grains have been replaced by fine grains. The grains are fine and homogeneous. There are even a few large grains and the average grain size increases, as shown in Fig. 6, when the annealing time is up to 30 min. The grains are not uniform any more, as shown in Fig. 5c. This is because there are some small unstable recrystallization nuclei during annealing treatment. In order to keep a stable structure, grain growth takes place by minishing the grain boundary area per unit volume [56]. Due to the different energy in each grain, there are big differences on the final grain size. Hence, the grains become heterogeneous again with the increase of time. Thus, it can be concluded that the recrystallized grains grow up gradually when the annealing time is increased from 5 to 30 min. Meanwhile, the initial heterogeneous grains become homogeneous firstly, and heterogeneous again.
temperature is increased from 950 °C to 1010 °C, the average grain size decreases firstly and then increases, as shown in the Fig. 4. The reasons are as follows. In general, the promoting effect on recrystallization becomes more and more obvious as the annealing temperature is increased. Meanwhile, it also improves the migration rate of grain boundaries by providing more energy [52]. However, for the studied alloy, the solution temperature of delta phase is 980 °C. In the process of annealing, the large long needle-like delta phases gradually become small spherical. Previous studies [53,54] have shown that the second phases with different size have different effects on recrystallization. There are high dislocation density areas around the second phase with large size, which accelerates the occurrences of recrystallization by increasing nucleation sites. The fine second phases dispersed and distributed in matrix have pinning effects on the migration of grain boundaries. So, it leads to the restriction on the growth of grains. The fraction of fine second phase at 980 °C is higher than that at 950 °C.·This is the reason why the average grain size decreases firstly. When the annealing temperature is increased to 1010 °C, the average grain size increases obviously, as shown in Fig. 4c. This is because the delta phases dissolve quickly at the temperature of 1010 °C. Therefore, it cannot inhibit the grain growth any more. From Fig. 4, it can also be found that the microstructures treated at the annealing temperature of 950 °C and 980 °C for 10 min are similar, but there are obvious differences in the fraction of delta phase, which will be discussed in Section 3.4.
3.2.2. Influences of Annealing Time Fig. 5 indicates the microstructure of the deformed specimens annealed at the temperature of 980 °C with different annealing time. It can 216
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
0.5
0.4
d=9.2µm
0.3 0.2
0.2 0.1
0.1 0.0
d=11.94µm
0.3 Fraction
Fraction
0.4
0.0 0
5 10 15 20 25 30 35 40 45 50
0
5
10
(a) 0.5
Average grain size (µm)
Fraction
0.3 0.2 0.1 10
20
30
40
50
25
25
d=22.51µm
0
20
30
35
(b)
0.4
0.0
15
Grain size(µm)
Grain size(µm)
60
70
80
22.51
20 15
11.94 9.2
10 5 0
0
5
Grain size(µm)
10
30
Annealing time (min)
(c)
(d)
Fig. 6. Grain size distribution maps of the deformed specimens at the annealing temperature of 980 °C, and the time of: (a) 5 min; (b) 10 min; (c) 30 min; (d) Histogram of average grain size.
direction of grain boundaries is hard to over 5°, while the cumulative misorientation is easy to over 15°, as shown in Fig. 7a and b. This illustrates there is a relatively steady misorientation gradient both in the parallel and vertical direction of grain boundaries. In addition, the large cumulative misorientation indicates that there is a large orientation gradient both in the parallel and vertical direction of grain boundaries, where the recrystallization grain boundaries can easily transform into HAGBs because it is easy to accumulate enough misorientations [59]. For the annealed specimens, it can be found that both the local and cumulative misorientation in the parallel and vertical direction of grain boundaries are hard to over 1°, as shown in Fig. 7c and d. It shows that the orientation gradient and little deformation energy in the both directions are little. Similar phenomenon can be found from Fig. 7e–h. It indicates that the static recrystallization is fully completed, and all of the large elongated deformed grains are replaced by recrystallization grains. The appearance of large grains shown in Fig. 3g is due to the growth of recrystallization grains rather than the residue of the deformed grains.
3.3. Misorientation Angle Evolution 3.3.1. Influences of Annealing Temperature Fig. 3 shows the orientation image microscopy (OIM) maps of the deformed specimens and the specimens annealed at various annealing temperatures, respectively. It is easy to find that the microstructure after deformation is heterogeneous. Most of grain boundaries still belong to LAGBs, as shown in Fig. 3a marked by thin-gray lines. The fraction of LAGBs is relatively high and the average misorientation angle is 11.3°, as shown in Fig. 3b. It indicates that the fraction of recrystallization is very low after deformation. Compared to deformed specimens, the fraction of LAGBs after annealing treatment is significantly low, and the average misorientation angle increases obviously, as shown in Fig. 3. Additionally, the fraction of 60° boundaries increases largely after the annealing treatment. It shows that there are some annealing twins formed during annealing treatment. Generally, the boundary energy of growing grain is decreased by annealing twins, while the mobility of grain boundary is enhanced [57].Therefore, the annealing twins can accelerate the process of static recrystallization. When the annealing temperature is increased from 950 to 1010 °C, the fraction of LAGBs decreases and the average misorientation angle increases slightly. It indicates that the fraction of recrystallization for the latter is higher. This is because the migration capacity of the grain boundaries can be improved by elevating annealing temperature, which promotes the recrystallization. That is to say the recrystallization is easy to occur at higher temperatures [58]. The fraction of LAGBs decreases, and the fraction of recrystallization and the average misorientation angle increases with the increased annealing temperature. The point to point (local) misorientation and point to origin (cumulative) misorientations along the lines marked in Fig. 3 are calculated, as shown in Fig. 7. For the deformed specimens, it is easy to discover that the local misorientation both in the parallel and vertical
3.3.2. Influences of Annealing Time Fig. 5 shows the OIM maps and misorientation angle distributions of the deformed specimens at the annealing temperature of 980 °C with different annealing times. Compared to the initial deformed specimens, the specimens after annealing treatment at 980 °C for 5 min is dominated by HAGBs, and the fraction of LAGBs is extremely low. Meanwhile, the average misorientation angle is 48.5°. As the annealing time is further increased to 30 min, it is almost impossible to observe LAGBs again, and the average misorientation angle increases to 50.8°. It can be concluded that the fraction of LAGBs decreases and the average misorientation angle increases with the increasing annealing time. It is because there are many small recrystallization nuclei in the deformed specimens, and there is a certain strain energy accumulated in the 217
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
Point to origin Point to point
16 12 8 4 0
0
10
20
30
40
Point to origin Point to point
30 Misorientation angle (°)
Misorientation angle (°)
20
25 20 15 10 5 0
50
0
5
10
Distance (µm)
(b) Point to origin Point to point
0.3 0.2 0.1
5
10
15
20
25
30
Point to origin Point to point
0.4 Misorientation angle (°)
Misorientation angle (°)
0.4
0
0.3 0.2 0.1 0.0
35
0
2
4
Distance (µm)
0.2
0.1
10
15
20
0.2 0.1 0.0
25
0
5
10
Point to origin Point to point
Misorientation angle (°)
Misorientation angle (°)
0.30
0.20 0.15 0.10 0.05 0
5
20
(f)
0.25
0.00
15
Distance (µm)
(e) 0.30
12
0.3
Distance (µm)
0.35
10
Point to origin Point to point
0.4 Misorientation angle (°)
Misorientation angle (°)
Point to origin Point to point
5
8
(d)
0.3
0
6
Distance (µm)
(c)
0.0
20
Distance (µm)
(a)
0.0
15
0.20 0.15 0.10 0.05 0.00
10 15 20 25 30 35 40 45
Point to origin Point to point
0.25
0
5
Distance (µm)
10
15
20
Distance (µm)
(g)
(h)
Fig. 7. Changes of misorientation angle along the lines marked in Fig. 3: (a) A1; (b) A2; (c) B1; (d) B2; (e) C1; (f) C2; (g) D1; (h) D2.
annealing treatment at 980 °C for 5 min are higher than that for 10 min and 30 min, as shown in Fig. 8. It indicates there is still a little orientation gradient in both directions. The recrystallization is still not fully complete. Therefore, the large grains shown in Fig. 5a are the residue of the deformed grains. It illustrates that more time is needed for the occurrence of recrystallization. With the increase of annealing time, the local and cumulative misorientations in parallel and vertical directions of grain boundaries are little. It indicates there are only small orientation gradients in both directions, and the deformation energy
deformation zone without complete recrystallization. Therefore, the new recrystallization nuclei grew up in the process of annealing, as well as the DRX nuclei formed in the deformation, which cause the LAGBs change to HAGBs. There is more time for recrystallization with the increase of annealing time. So, the fraction of recrystallization increases. This causes the fraction of LAGBs to decrease and the average misorientation angle to increase. Meanwhile, the cumulative misorientations in the parallel and vertical directions of grain boundaries in the specimen after the 218
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
3.5
Point to origin Point to point
3.0
Misorientation angle (°)
Misorientation angle (°)
3.5 2.5 2.0 1.5 1.0 0.5 0.0
0
10
20
30
40
50
60
2.5 2.0 1.5 1.0 0.5 0.0
70
Point to origin Point to point
3.0
0
5
Distance (µm)
10
Misorientation angle (°)
Misorientation angle (°)
(b) 0.6
Point to origin Point to point
1.0 0.8 0.6 0.4 0.2 0.0
0
10
20
30
20
Distance (µm)
(a) 1.2
15
40
50
0.4 0.3 0.2 0.1 0.0
60
Distance (µm)
Point to origin Point to point
0.5
0
5
10
15
20
25
30
35
Distance (µm)
(c)
(d)
Fig. 8. Changes of misorientation angle along the lines marked in Fig. 5: (a) E1; (b) E2; (c) F1; (d) F2.
Fig. 9. SEM micrographs of the deformed specimens at the annealing time of 60 min, and the temperature of: (a)950 °C; (b) 965 °C; (c) 980 °C.
3.4. Delta Phase Evolution
has been depleted. Therefore, it is difficult to migrate for grain boundaries by dislocation slip or climb [52]. The grain growth behavior, as shown in Fig. 5c, is mainly caused by the HAGBs migration triggered by atom motion [60]. In addition, it should be pointed out that there is seemingly a slight increase in the cumulative misorientations inside the recrystallization grains when the annealing time is increased from 10 to 30 min, as shown in Figs. 7e–f and 8c–d. However, the increase is not caused by the increase of annealing time. It only attributes to the wave of misorientation inside each recrystallization grain. The difference between the values shown in Figs. 7e–f and 8c–d results from the uncertainty of the position of underlines.
3.4.1. Influences of Annealing Temperature Fig. 9 shows the SEM micrographs of the deformed samples annealed at different temperatures with the annealing time of 60 min. The needle-like delta phases become rod-like and spherical during annealing treatment, as shown in Fig. 9. The fraction of needle-like delta phase decreases with the increase of temperature. When the temperature reaches 980 °C, all the needle-like delta phases are dissolved. The change of morphology in delta phase is a phase boundary migration process, which is controlled by the atomic diffusion. There is more energy for the atomic diffusion in high temperature. Therefore, the proportion of needle-like delta phase decreases with the increase of temperature. The effect of annealing temperature on the fraction of delta phase is shown in Fig. 10. It can be found that the fraction of delta 219
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
The delta phase which is perpendicular to the direction of compression is subjected to tensile stress, which leads to the phenomenon of tension fracture during deformation. While the delta phases distributed along the direction of compression are bent or broken to adapt to the deformation. Thus, there are a lot of delta phases that have been broken after deformation. The morphology of these delta phases has changed from needle-like to rod-like and spherical. It means that the decomposition of these phase have finished in deformation. During the annealing, these delta phases will dissolve with a high rate, which lead to the dissolution rate being higher than the precipitation one. Therefore, the fraction of delta phase decreases at the early stage of annealing. It indicates that the deformation accelerates the dissolution of delta phase. However, all of these delta phases have dissolved as the annealing time further increases. Thus, the precipitation becomes stronger than dissolution. So, the fraction of delta phase increases. Although the fraction of delta phase at 30 min is lower than the one at 10 min, the grains are heterogeneous after annealing at 980 °C for 30 min, as shown in Fig. 5c. All things considered, the annealing time should be chosen as 10 min.
Fig. 10. Effects of annealing temperature on the fraction of delta phase (The annealing time is 60 min).
phase decreases from 3.9% to 0.52% when the annealing temperature is increased from 950 to 980 °C. This is because the annealing temperature has a decisive effect on the diffusion process of delta phase. With the increased temperature, the dissolution rate increases owing to the high diffusion coefficient at high temperature [61]. The excessive content of delta phase can promote crack propagation rate, and reduce the fatigue life of workpiece [35].Therefore, both considering the fatigue life and the uniformity of microstructure, the annealing temperature should be chosen as 980 °C.
4. Conclusions (1) It has been found that adopting an annealing treatment after deformation can effectively improve the homogeneity of the microstructure. The optimum parameters are the annealing temperature of 980 °C and time of 10 min. (2) The grains become homogeneous due to the occurrence of static recrystallization. There is an incubation period which is not less than 5 min during the static recrystallization. The initial heterogeneous grains become homogeneous firstly, and then heterogeneous again when the annealing time is increased from 5 to 30 min. (3) The deformation accelerates the dissolution of delta phase in annealing treatment. The fraction of delta phase decreases at the early stage of annealing whether the annealing temperature exceeds the solution temperature.
Acknowledgements This work was supported by Hunan Provincial Natural Science Foundation of China (2017JJ3380), National Natural Science Foundation of China (No. 51775564), the Science and Technology Leading Talent in Hunan Province (No. 2016RS2006), Program of Chang Jiang Scholars of Ministry of Education (No. Q2015140), the Open-End Fund for the Valuable and Precision Instruments of Central South University (No. CSUZC201821)，and Hebei Iron and Steel Joint Funds (No. E2015209243).
7
7
6
6 Fraction (%)
Fraction (%)
3.4.2. Influences of Annealing Time Figs. 11 and 12 show the effects of annealing time on the fraction and morphology of delta phase, respectively. It can be seen that the annealing time has a great effect on the fraction of delta phase. As the annealing time is increased, the fraction of delta phase decreases gradually at the annealing temperature of 980 °C. The fraction of delta phase decreases from 6.8% to 0.52% when the annealing time is increased to 60 min, as shown in Fig. 11(a). Meanwhile, the needle-like delta phase becomes spherical, as shown in Fig. 12. However, at the annealing temperature of 950 °C which is lower than the solution temperature of delta phase, the fraction of delta phase firstly decreases and then increases with the increasing annealing time. The reasons are as follows. According to the existing literature [61,62], the static dissolution of delta phase mainly consists of three parts: (1) the decomposition of delta phase; (2) the short-range diffusional process of atoms; (3) a longer-range diffusional process of atoms. But, the time for the decomposition of delta phase is much more than the short-range diffusional process of atoms and longer-range diffusional process of atoms.
5 4 3 2 1 0
5 4 3 2 1
0
10
20
30
40
50
0
60
Annealing time (min)
(a)
0
10 60 Annealing time (min)
180
(b)
Fig. 11. Effects of annealing time on the fraction of delta phase at the annealing temperature of: (a) 980 °C; (b) 950 °C. 220
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
Fig. 12. SEM micrographs of the deformed specimens at the annealing temperature of 980 °C, and the time of: (a) 10 min; (b) 30 min.
References
[20] Y. Wang, Q. Pan, Y. Song, C. Li, Z. Li, Hot deformation and processing maps of X750 nickel-based superalloy, Mater. Des. 1 (2013) 154–160. [21] P.M. Mignanelli, N.G. Jones, K.M. Perkins, M.C. Hardy, H.J. Stone, Microstructural evolution of a delta containing nickel-base superalloy during heat treatment and isothermal forging, Mater. Sci. Eng. A 621 (2015) 265–271. [22] M. Retima, S. Bouyegh, H. Chadli, Effect of the heat treatment on the microstructural evolution of the nickel based superalloy, Meta 17 (2011) 71–77. [23] X.M. Chen, Y.C. Lin, D.X. Wen, J.L. Zhang, M. He, Dynamic recrystallization behavior of a typical nickel-based superalloy during hot deformation, Mater. Des. 57 (2014) 568–577. [24] M.S. Chen, Y.C. Lin, K.K. Li, Y. Zhou, A new method to establish dynamic recrystallization kinetics model of a typical solution-treated Ni-based superalloy, Comput. Mater. Sci. 122 (2016) 150–158. [25] Y.X. Liu, Y.C. Lin, Y. Zhou, 2D cellular automaton simulation of hot deformation behavior in a Ni-based superalloy under varying thermal-mechanical conditions, Mater. Sci. Eng. A 691 (2017) 88–99. [26] L.A. Reyes, P. Páramo, A. Salas Zamarripa, M. de la Garza, M.P. Guerrero Mata, Grain size modeling of a Ni-base superalloy using cellular automata algorithm, Mater. Des. 83 (2015) 301–307. [27] T.S. Prithiv, P. Bhuyan, S.K. Pradhan, V.S. Sarma, S. Mandal, A critical evaluation on efficacy of recrystallization vs. strain induced boundary migration in achieving grain boundary engineered microstructure in a Ni-base superalloy, Acta Mater. 146 (2018) 187–201. [28] C. Zhang, L. Zhang, M. Li, W. Shen, S. Gu, Effects of microstructure and γ′ distribution on the hot deformation behavior for a powder metallurgy superalloy FGH96, J. Mater. Res. 29 (2014) 2799–2808. [29] Y.H. Liu, Y.Q. Ning, X.M. Yang, Z.K. Yao, H.Z. Guo, Effect of temperature and strain rate on the workability of FGH4096 superalloy in hot deformation, Mater. Des. 95 (2016) 669–676. [30] J.T. Yeom, C.S. Lee, J.H. Kim, N.K. Park, Finite-element analysis of microstructure evolution in the cogging of an alloy 718 ingot, Mater. Sci. Eng. A 449 (2007) 722–726. [31] A.J. Brand, K. Karhausen, R. Kopp, Microstructural simulation of nickel base alloy inconel 718 in production of turbine discs, Mater. Sci. Technol. 12 (2014) 963–969. [32] Y.C. Lin, X.M. Chen, A critical review of experimental results and constitutive descriptions for metals and alloys in hot working, Mater. Des. 32 (2011) 1733–1759. [33] X. You, Y. Tan, J. Li, P. Li, C. Dong, S. Shi, S. Qin, Effects of solution heat treatment on the microstructure and hardness of inconel 740 superalloy prepared by electron beam smelting, J. Alloys Compd. 638 (2015) 239–248. [34] A. Ghosh, N.P. Gurao, Effect of crystallographic texture on ratcheting response of commercially pure titanium, Mater. Des. 115 (2017) 121–132. [35] J. An, L. Wang, Y. Liu, W. Cai, X. Song, The role of δ phase for fatigue crack propagation behavior in a Ni base superalloy at room temperature, Mater. Sci. Eng. A 684 (2017) 312–317. [36] D. Hu, Q. Yang, H. Liu, J. Mao, F. Meng, Y. Wang, R. Wang, Crack closure effect and crack growth behavior in GH2036 superalloy plates under combined high and low cycle fatigue, Int. J. Fatigue 95 (2017) 90–103. [37] G.Z. Quan, Y.L. Li, L. Zhang, X. Wang, Evolution of grain refinement degree induced by dynamic recrystallization for Nimonic 80A during hot compression process and its FEM analysis, Vacuum 139 (2017) 51–63. [38] J. Kang, D.S. Wilkinson, R.K. Mishra, W. Yuan, R.S. Mishra, Effect of inhomogeneous deformation on anisotropy of AZ31 magnesium sheet, Mater. Sci. Eng. A 567 (2013) 101–109. [39] Z. Sun, H. Yang, N. Sun, Effects of parameters on inhomogeneous deformation and damage in isothermal local loading forming of Ti-alloy component, J. Mater. Eng. Perform. 21 (2012) 313–323.
[1] H.Y. Li, J.D. Hu, D.D. Wei, X.F. Wang, Y.H. Li, Artificial neural network and constitutive equations to predict the hot deformation behavior of modified 2.25 Cr–1Mo steel, Mater. Des. 42 (2012) 192–197. [2] D.X. Wen, Y.C. Lin, H.B. Li, X.M. Chen, J. Deng, L.T. Li, Hot deformation behavior and processing map of a typical Ni-based superalloy, Mater. Sci. Eng. A 591 (2014) 183–192. [3] M.R. Jahangiri, H. Arabi, S.M.A. Boutorabi, Comparison of microstructural stability of IN939 superalloy with two different manufacturing routes during long-time aging, Trans. Nonferrous Metals Soc. China 24 (2014) 1717–1729. [4] Z.H. Yao, Q.Y. Wang, M.C. Zhang, J.X. Dong, Microstructure control and prediction of GH738 superalloy during hot deformation, Acta Metall. Sin. 47 (2011) 1591–1599. [5] Y.C. Lin, D.X. Wen, J. Deng, G. Liu, J. Chen, Constitutive models for high-temperature flow behaviors of a Ni-based superalloy, Mater. Des. 59( (2014) 115–123. [6] Y. Liu, Z. Yao, Y. Ning, Y. Nan, H. Guo, C. Qin, Z. Shi, The flow behavior and constitutive equation in isothermal compression of FGH4096-GH4133B dual alloy, Mater. Des. 63( (2014) 829–837. [7] F.M. Shore, M. Morakabati, S.M. Abbasi, A. Momeni, Hot deformation behavior of incoloy 901 through hot tensile testing, J. Mater. Eng. Perform. 23 (2014) 1424–1433. [8] Y.C. Lin, K.K. Li, H.B. Li, J. Chen, X.M. Chen, D.X. Wen, New constitutive model for high-temperature deformation behavior of inconel 718 superalloy, Mater. Des. 74( (2015) 108–118. [9] Y.C. Lin, X.M. Chen, D.X. Wen, M.S. Chen, A physically-based constitutive model for a typical nickel-based superalloy, Comput. Mater. Sci. 83 (2014) 282–289. [10] Y. Ning, Z. Yao, X. Liang, Y. Liu, Flow behavior and constitutive model for Ni–20.0 Cr–2.5 Ti–1.5 Nb–1.0 Al superalloy compressed below γ′-transus temperature, Mater. Sci. Eng. A 551 (2012) 7–12. [11] P. Zhang, C. Hu, Q. Zhu, C.G. Ding, H.Y. Qin, Hot compression deformation and constitutive modeling of GH4698 alloy, Mater. Des. 65( (2015) 1153–1160. [12] H.Y. Wu, F.J. Zhu, S.C. Wang, W.R. Wang, C.C. Wang, C.H. Chiu, Hot deformation characteristics and strain-dependent constitutive analysis of inconel 600 superalloy, J. Mater. Sci. 47 (2012) 3971–3981. [13] A. Etaati, K. Dehghani, G.R. Ebrahimi, H. Wang, Predicting the flow stress behavior of Ni-42.5 Ti-3Cu during hot deformation using constitutive equations, Met. Mater. Int. 19 (2013) 5–9. [14] Q. Yu, J. Liu, W. Zhang, L. Yu, F. Liu, W. Sun, Modification of Johnson-Cook model for hot deformation behavior of inconel 625 alloy, Rare Metal Mater. Eng. 41 (2013) 1679–1684. [15] A. Momeni, S.M. Abbasi, M. Morakabati, H. Badri, X. Wang, Dynamic recrystallization behavior and constitutive analysis of Incoloy 901 under hot working condition, Mater. Sci. Eng. A 615( (2014) 51–60. [16] D.X. Wen, Y.C. Lin, J. Chen, J. Deng, X.M. Chen, J.L. Zhang, M. He, Effects of initial aging time on processing map and microstructures of a nickel-based superalloy, Mater. Sci. Eng. A 620 (2015) 319–332. [17] P. Zhang, C. Hu, C.G. Ding, Q. Zhu, H.Y. Qin, Plastic deformation behavior and processing maps of a Ni-based superalloy, Mater. Des. 65 (2015) 575–584. [18] Y. Liu, Y. Ning, Z. Yao, H. Guo, Y. Nan, Effect of true strains on processing map for isothermal compression of Ni–20.0 Cr–2.5 Ti–1.5 Nb–1.0 Al Ni-base superalloy, J. Alloys Compd. 612 (2014) 56–63. [19] D.G. He, Y.C. Lin, M.S. Chen, J. Chen, D.X. Wen, X.M. Chen, Effect of pre-treatment on hot deformation behavior and processing map of an aged nickel-based superalloy, J. Alloys Compd. 649 (2015) 1075–1084.
221
Materials Characterization 141 (2018) 212–222
M.-S. Chen et al.
[51] Y.C. Lin, D.G. He, M.S. Chen, X.M. Chen, C.Y. Zhao, X. Ma, Z.L. Long, EBSD analysis of evolution of dynamic recrystallization grains and δ phase in a nickel-based superalloy during hot compressive deformation, Mater. Des. 97 (2016) 13–24. [52] X.M. Chen, Y.C. Lin, F. Wu, EBSD study of grain growth behavior and annealing twin evolution after full recrystallization in a nickel-based superalloy, J. Alloys Compd. 724 (2017) 198–207. [53] R.L. Goetz, Particle stimulated nucleation during dynamic recrystallization using a cellular automata model, Scr. Mater. 52 (2005) 851–856. [54] F.J. Humphreys, Particle stimulated nucleation of recrystallization at silica particles in nickel, Scr. Mater. 43 (2000) 591–596. [55] X. Wei, W. Zheng, Z. Song, T. Lei, Q. Yong, Q. Xie, Static recrystallization behavior of inconel 718 alloy during thermal deformation, J. Wuhan Univ. Technol. 29 (2014) 379–383. [56] Y.S. Jang, D.C. Ko, B.M. Kim, Application of the finite element method to predict microstructure evolution in the hot forging of steel, J. Mater. Process. Technol. 101 (2000) 85–94. [57] S.P. Coryell, K.O. Findley, M.C. Mataya, E. Brown, Evolution of microstructure and texture during hot compression of a Ni-Fe-Cr superalloy, Metall. Mater. Trans. A 43 (2012) 633–649. [58] X.M. Chen, Y.C. Lin, M.S. Chen, H.B. Li, D.X. Wen, J.L. Zhang, Microstructural evolution of a nickel-based superalloy during hot deformation, Mater. Des. 77 (2015) 41–49. [59] H. Zhang, K. Zhang, S. Jiang, H. Zhou, C. Zhao, X. Yang, Dynamic recrystallization behavior of a γ′-hardened nickel-based superalloy during hot deformation, J. Alloys Compd. 623 (2015) 374–385. [60] V. Randle, P.R. Rios, Y. Hu, Grain growth and twinning in nickel, Scr. Mater. 58 (2008) 130–133. [61] D. Cai, W. Zhang, P. Nie, W. Liu, M. Yao, Dissolution kinetics of δ phase and its influence on the notch sensitivity of inconel 718, Mater. Charact. 58 (2007) 220–225. [62] V. Beaubois, J. Huez, S. Coste, O. Brucelle, J. Lacaze, Short term precipitation kinetics of delta phase in strain free Inconel 718 alloy, Mater. Sci. Technol. 20 (2004) 1019–1026.
[40] H. Mirzadeh, M.H. Parsa, Hot deformation and dynamic recrystallization of NiTi intermetallic compound, J. Alloys Compd. 614 (2014) 56–59. [41] S. Saadatkia, H. Mirzadeh, J. Cabrera, Hot deformation behavior, dynamic recrystallization, and physically-based constitutive modeling of plain carbon steels, Mater. Sci. Eng. A 636 (2015) 196–202. [42] G.R. Ebrahimi, A. Momeni, H.R. Ezatpour, Modeling the viscoplastic behavior and grain size in a hot worked Nb-bearing high-Mn steel, Mater. Sci. Eng. A 714 (2018) 25–35. [43] C. Zhang, L.W. Zhang, W.F. Shen, Q.H. Xu, Y. Cui, The processing map and microstructure evolution of Ni-Cr-Mo-based C276 superalloy during hot compression, J. Alloys Compd. 728 (2017) 1269–1278. [44] Y.Q. Ning, C. Zhou, H.Q. Liang, M.W. Fu, Abnormal flow behavior and necklace microstructure of powder metallurgy superalloys with previous particle boundaries (PPBs), Mater. Sci. Eng. A 652 (2016) 84–91. [45] S.K. Pradhan, S. Mandal, C.N. Athreya, K.A. Babu, B. de Boer, V.S. Sarma, Influence of processing parameters on dynamic recrystallization and the associated annealing twin boundary evolution in a nickel base superalloy, Mater. Sci. Eng. A 700 (2017) 49–58. [46] X.F. Tang, B.Y. Wang, H.C. Ji, X.B. Fu, W.C. Xiao, Behavior and modeling of microstructure evolution during metadynamic recrystallization of a Ni-based superalloy, Mater. Sci. Eng. A 675 (2016) 192–203. [47] G.Z. Quan, J. Pan, X. Wang, Z.H. Zhang, L. Zhang, T. Wang, Correspondence between grain refinements and flow softening behaviors at Nimonic 80A superalloy under different strain rates, temperatures and strains, Mater. Sci. Eng. A 679 (2017) 358–371. [48] A. Sarkar, S. Sanyal, T.K. Bandyopadhyay, S. Mandal, Influence of annealing parameters on phase evolution and recrystallization kinetics of a Mn-Al-Si alloyed duplex steel, Mater. Charact. 134 (2017) 213–224. [49] Y.C. Lin, X.Y. Wu, X.M. Chen, J. Chen, D.X. Wen, J.L. Zhang, L.T. Li, EBSD study of a hot deformed nickel-based superalloy, J. Alloys Compd. 640 (2015) 101–113. [50] H. Jiang, J. Dong, M. Zhang, Z. Yao, Evolution of twins and substructures during low strain rate hot deformation and contribution to dynamic recrystallization in alloy 617B, Mater. Sci. Eng. A 649 (2016) 369–381.
222