CC BY
17
DOI: 10.17277/amt.2019.04.pp.013-026
The Formation Mechanism of Microduplex Type Structure During Thermomechanical Treatment of Superalloy
V.A. Valitov
Institute for Metal Superplasticity Problems of the Russian Academy of Sciences, 39, Stepan Khalturin ul, Ufa, Republic of Bashkortostan, 450001, Russian Federation
* Corresponding author. Tel.: +7 347 223 64 07. E-mail: Valitov_VA@mail.ru
Abstract
The influence regularities of the deformation processing conditions as well as the size of the strengthening Y -phase on the formation of a fine-grained microduplex type structure in superalloy are studied. It is shown that during hot deformation of alloys with an isomorphic Y -phase, depending on its size, volume fraction and degree of deformation, recrystallization proceeds both by continuous and discontinuous mechanisms. It was established that during hot deformation in the y +y' region (0,7-0,85 Tm) of the superalloy with a precoagulated y'-phase, continuous dynamic recrystallization develops, during which a subgrain structure forms in the initial large grains with the size of subgrains commensurate with the interparticle distance, and its gradual transformation into a fine-grained microduplex type structure with high-angle interphase and intergrain boundaries is observed.
Keywords
Microduplex structure; heat resistant alloy; deformation processing; y'-phase.
© V.A. Valitov, 2019
Introduction
The formation of a microduplex type structure in the precipitation hardening of the superalloy occurs under complex conditions of interaction between the recrystallization and the precipitation processes of the strengthening y'-phase [1-6]. The regularities of such a structure formation were mainly considered for the case of static recrystallization of low alloyed alloys. It should be noted that the development of recrystallization largely depends on the particle size of the y'-phase, the distance between them and the type of coupling with the matrix. The influence of these factors on the formation of a micro duplex during hot deformation, as well as the temperature and speed conditions of the latter, has not been sufficiently studied.
The paper presents the results of systematic studies of the influence of the regimes of cold deformation, annealing, and hot deformation on the transformation of a coarse-grained structure into a fine-grained
microduplex type in the superalloy, hardened by an isomorphic y'-phase based on Ni3(Al, Ti) intermetallic compound.
Materials and research methods
The superalloy, containing 45% of the Y'-phase and having a temperature of its complete dissolution of 1150 °C was selected as the material under study. The chemical composition of the alloy is shown below: C - 0.1; Cr - 13.09; W - 2.82; V - 0.36; Mo - 4.63; Al - 3.22; Ti - 2.64; Fe - 0.58; Co - 10.74; Nb - 3.44; Ni - main (wt.%). Before deformation, the samples cut from the hot-pressed rod were subjected to homogenization annealing at a temperature of 1180 °C for 1 hour. Some of the samples from the homo-genization temperature were cooled in air (state 1), and the others were cooled with an oven at a rate of 20-30 °C / hour up to 900 °C (state 2), then in air. Samples with a diameter of 8 mm and a height of 12 mm were deposited on a 1231U-10 apparatus under
a) b) c) d)
Fig. 1. The microstructure of the superalloy after homogenizing annealing and subsequent cooling in air (state 1)
and with the furnace (state 2):
a, b - state 1; c, d - state 2
isothermal conditions at 1000-1150 °C with degrees from 2 to 75 % and strain rates The exposure of the samples at the test temperature before deformation was 25 min. To establish the features of the microstructure change during hardening and thermal treatment, some of the samples were subjected to cold deformation at 20 °C by 10 and 50 % and subsequent annealing at 1100 °C for 1 h, as well as annealing of undeformed samples at the same temperature for a period of 0.5 to 100 hours. The microstructure of water-quenched samples was studied by optical (Neofot-32, Epiquant) and electron (BS-540, JEM-2000EX, JSM-840) microscopy. Foils were made according to the method described in [7].
Research results and discussion
The original microstructure. Metallographic studies showed that after homogenizing annealing, an equiaxed coarse-grained microstructure with an average y'-phase grain size of 80 (m was formed in the alloy. In this case, state 1 was characterized by the presence of dispersed coherent particles of the y'-phase with a size of 0.1 (m, homogeneously isolated in the bulk of the matrix grains (Fig. 1).
At the same time, in state 2, the intragranular y'-phase was coagulated to a size of 0.5 (m, had a predominantly cuboid shape and partially disturbed coherence, as evidenced by individual misfit dislocations at the interphase boundaries. The average distance between particles was 0.25 microns. At the grain boundaries, the width of the zones free from precipitation reached 1-2 (m. Grain-boundary precipitates of the y'-phase were 2-3 times larger than intragranular ones, had a globular shape, partially coherent border with one grain and high angle with another bordering grain.
It should be noted that the formation of a microduplex structure is a complex process, determined not only by the conditions of heterogeneous structure, but also by phase and microstructural changes that occur during heating and hot deformation. Consider the role of each of these processes sequentially.
Heterogenization. We studied the change in the morphology, amount, and size of y'-phase precipitates upon heating of the alloy to 1100 °C. As can be seen in Fig. 2, the volume fraction of the y'-phase decreased to equilibrium upon exposure, but the size and distance between the particles of the y'-phase increased.
X, dY', |im
1.2
1.0
0.8
0.6
0.4
0.2
xs
/ / f T J / T A
fL , dy V y y
f
Y
f, %
60 50 40 30 20 10 .0
0
1
2
3
4 t, hours
Fig. 2. Dependences of the change in the /-phase parameters on the annealing time at 1100 °C of the superalloy:
x - state 1; O - state 2 (f is volume fraction of the y'-phase; X is the interparticle distance; dY' is the particle size of the y'-phase; t is holding time, hour)
Along with an increase in the particle size of the Y'-phase, a change in the structure of interphase boundaries occurred. Characteristic of state 1 is the alignment of cuboid particles of the Y'-phase into chains forming a spatial "quasiperiodic" lattice typical of nickel alloys [8, 9]. The coherence of cuboid particles of the Y'-phase was maintained after holding for |2-8 hours (Fig. 3), although there was an increase in the size of the Y'-phase to 0.4 ^m. Only after prolonged exposure for 50-100 hours in this state, a partial violation of coherence was observed at the interphase boundary of Y'-phase particles, enlarged to a size of 0.71-0.86 ^m.
In state 2, the processes of spheroidization, coalescence of particles of the Y'-phase, occurred intensively, resulting in a change in the structure of the
Y / y' interfaces. If at the initial moment most of the
Y / y' boundaries were coherent, then after holding for more than 30 min. mismatch dislocations appearred at the y / Y' interfaces (Fig. 4b, c, e, f), causing a partial loss of coherence. Fig. 4 shows that the y / Y' interface contained mismatch dislocations having, according to [10], a Burgers vector of type a/2 <110>. Long-term
annealing of samples (50-100 hours) in state 2 contributed to the enlargement of Y'-phase particles to a size of 0.89-0.98 ^m, but did not lead to the conversion of partially coherent interphase boundaries to incoherent general types.
Fig. 2 shows that in both states a linear
dependence of the particle size dy on the cubic root of the duration of annealing is observed, i.e. complies with the Livshits-Slezov-Wagner law [11, 12]. So, after heating and holding at 1100 °C for 0.5-8 hours, characteristic of hot deformation processes, in state 2 there is a heterogenization of the structure - intensive coagulation of the y'-phase and the formation of partially coherent interphase boundaries.
Mechanical properties. Heterogenization of the microstructure and temperature-velocity conditions of deformation significantly affect the flow stress and the shape of the c-s curves (Fig. 5). At all test temperatures, the flow stress of the alloy in state 1 is noticeably greater than in state 2. The higher the temperature and the lower the strain rate were, the lower the c-s peak in the c-s curves was.
d)
e)
f)
Fig. 3. The microstructure of the superalloy in state 1 after isothermal annealing at 1100 °C:
a, b, c - exposure time of 8 hours; d, e - exposure time of 50 hours; f - exposure time of 100 hours
а) b) c)
d) e) f)
Fig. 4. The microstructure of the superalloy in state 2 after isothermal annealing at 1100 °C:
a, b - exposure time of 2 hours; c - exposure time of 8 hours; d, e - exposure time of 50 hours; f- exposure time of 100 hours
a further increase in the test temperature and a decrease in the strain rate.
The state of the alloy also affected the nature of the c-s curves, which was especially noticeable at a test temperature of 1100 °C. Thus, when the alloy was deformed in state 2, the peak of the flow stress was absent on the c-s curve and, in addition, the oscillation of the flow stress was observed on it. Evaluation of the coefficient of velocity sensitivity of the flow stress showed that in state 2 its greatest value m = 0.31 was
-5 _3 -1
achieved at 1100 °C in the speed range 10 -10 s with a degree of deformation of 50 %, although at s = 5 % the ratio was 0.23. For state 1, the coefficient m in the studied temperature - speed range did not exceed 0.16.
The maximum allowable degree of deformation, above which cracks appear on the lateral surface of the
-3 -1
samples at 1100 °C and a strain rate of 10 s , was 60 % and was 1.5-2 times greater than in state 1. When the temperature was reduced to 1000 °C. With an increase in the strain rate to 10-2-10-1 s-1 in state 2, it decreased by 15-30 %.
a, MPa 400 300 200 100
4
N \
' 4 5
6
r" a ™— _e^r1 9 1-Х- 10
10
20
30 40 50 E, %
Fig. 5. Dependence of the flow stress of the superalloy on temperature and strain rate:
x - state 1; О - state 2 (10-3 s-1: 1, 3 - 1000 °C; 2, 5 - 1050 °C; 6, 7 - 1100 °С;
9 - 1150 °С; 10-1 s-1: 4 - 1100 °C; 10-5 s-1: 8,10 - 1100 °С)
Upon reaching a temperature of 1050 °C and a speed of
_3 -1
10 s , the steady-state stage of the flow was observed on the a-s curves, the length of which increased with
Cold deformation and post-deformation annealing. Heterogenization has a significant impact on the nature of the plastic flow and the type of microstructure formed during recrystallization.
Optical metallography showed that after cold deformation by 50 % in state 1, coarse bands were observed; they were most likely strain localization bands, which, according to electron microscopic studies, corresponded to foil sections with an increased dislocation density (Fig. 6). Meanwhile, despite the rather high dislocation density, a substructure in this state did not form.
Annealing of cold-deformed samples led to the development of recrystallization, the mode of which substantially depended on the degree of preliminary deformation and the state of the y'-phase and was consistent with the results of studies reported in [3]. In state 1, recrystallization occurred extremely nonuniformly, mainly near the initial grain boundaries. In this case, it was characteristic that particles of the y'-phase were isolated in recrystallized grains mainly in
lamellar form and are coherently connected with the matrix (Fig. 6e, f).
In state 2, the character of the metallographic picture was different than in state 1. No traces of coarse slip bands were observed (Fig. 7). An increase in the degree of preliminary deformation from 10 to 50 % leads not only to a significant increase in the fraction of recrystallized volume, but also to a significant change in the morphology of y' -phase precipitates. In this case, along with the lamellar form, the microstructure contains particles of the y ' -phase of a globular shape with partially disturbed coherence. The microstructure was heterogeneous: along with y'-phase grains of 3-10 ^m in size, larger grains up to 40 ^m in size were present. In unrecrystallized areas, the dispersed Y' -phase was preserved.
A study of the fine microstructure showed that a substructure was formed at a degree of deformation of 50 %. A confirmation of this can be the blurring of reflections on the electron diffraction pattern (Fig. 7d).
■jm
iïaÈ&rsSk, m?
M
« Ipf!^
b)
* , /
d)
e)
f)
Fig. 6. Microstructure of the superalloy state 1 after cold deformation and annealing:
a,, d - e = 50 %; b, e - e = 10 %, +1100 °C, 1 hour; f- e = 50 %, +1100 °C, 1 hour
Fig. 7. Microstructure of the superalloy state 2 after cold deformation and annealing:
a - e = 10 %; d, e - s = 50 %; b, c - s = 10 %, +1100 °C, 1 hour; f- s = 50 %; +1100 °C, 1 hour
At a lower degree of deformation (10 %), particles of the y'-phase were not cut by dislocations whose density was highest near the y/y ' interface. At the same time, misorientations between microregions were not detected.
As a result of annealing of the state 2 samples deformed by 10 %, recrystallization developed mainly near the initial grain boundaries. In this case, particles of the Y -phase of a globular shape were present in recrystallized grains up to 20 ^m in size. It is important to note that the boundaries of recrystallized grains migrating into the deformed matrix have a tortuous shape that bent between coagulated particles of the y' -phase in contact with the boundary (Fig. 7b).
A completely recrystallized microstructure of the microduplex type with grain sizes of y and y' -phase, respectively 4.2 and 1.3 ^m, was formed in samples deformed by 50 %. In the resulting microstructure, almost the entire y' -phase is distinguished along the boundaries of recrystallized grains of the y' -phase and has incoherent interphase boundaries of a general type.
Hot deformation. Heterogenization of the microstructure also significantly affects the processes of structure formation during hot deformation. In state 1, after hot deformation by 5 % at 1100 °C with a speed of 10-5 s-1, separate grain cells with lamellar y' -phase particles are detected in the border regions of the initial grains. In this case, inside the grains there is a tendency
a) b) c) d)
Fig. 8. The microstructure of the superalloy in state 1 after hot deformation at 1100 °C at a speed of 10 s-5:
a, b - s = 5 %; c, d - s = 50 %
to align y' -phase particles in chains longer than during annealing, and their oriented growth is predominantly in the same crystallographic direction within individual grains of the y' -phase (Fig. 8). With an increase in the degree of deformation up to 50 %, the orientation of the particles of the y' -phase and a change in their shape manifest themselves more clearly. So, the average size of intragranular particles of the y'-phase in the longitudinal direction was 1.5, and in the transverse direction 0.4 ^m.
At the same time, in the zones of excretion-free zones, the formation of new recrystallized y'-phase grains with a size of 3.7 ^m occurs, containing у' -phase precipitates of both lamellar and globular shapes.
-3 -1
At strain rates of 10 s and higher, lamellar particles coherently associated with the matrix do not form. After intense deformation by 55-75 % at 1100 °С at a rate of 10-3 s-1, new recrystallized grains 3-4 pm in size are formed predominantly along the boundaries of the initial grains (Fig. 9a). According to morphological features, such a structure can be attributed to a structure of the "necklace" type [13].
Deformation with lower degrees (20-30 %) and a
-3 -1
rate of 10 s does not lead to recrystallization (Fig. 9b). At the same time, short-term post-deformation annealing (1100 °С, 1 hour) leads to tortuosity of the initial grain boundaries.
The development of dynamic recrystallization with the formation of a microduplex type structure is observed during hot deformation of the alloy in state 2 (Fig. 10). In this case, the average grain size of у and у ' -phase in the micro duplex, as well as the fraction of the recrystallized volume, decrease with decreasing temperature and a corresponding increase in the ratio of the volume fraction of the y' -phase to its size after heterogeneous annealing f / dy' (Tables 1 and 2).
a) b)
Fig. 9. Microstructure of the superalloy in state 1after hot deformation at 1100 °C at a speed of 10-4 s-1:
a - e = 75 %; b - e = 25 %, +1100 °С, 1 hour
In studying the evolution of the microstructure, the following was established. After deformation at 1100 °C by 5 % at a rate of 10-5 s-1, the formation of weakly oriented (less than 1°) subgrain fragments in the initial large grains of the matrix occurs, the size of which is determined by the distance between the particles of the y' -phase. It is important to note that the border subgrains are larger than the intragranular ones and have a misorientation close to high angle. As can be seen in Figs 10b and 10c, the misfit dislocations on the y / y ' interface are intertwined with the dislocations that make up the subboundaries in the matrix. An increase in the degree of deformation up to 10-20 % leads to the formation of small recrystallized grains along the initial boundaries of the chain, and the formation of thinner walls of subboundaries inside the deformed grains. In this case, the misorientation between subgrains, the average size of which was
m ■ ti-7 % ■ *'. ■-•
a)
6)
• A I
0.5 |am
c)
Fig. 10. Microstructure of the superalloy in state 2 after hot deformations at 1100 °C at a speed of 10 s
a, b, c - e = 5 %; d, e, f - s = 50 %
2.2 ^m, increased markedly. Separate dislocations are no longer detected on the interphase surface of y'-phase particles, enlarged to a value of (1 ± 0.2) ^m, which obviously indicates an almost complete loss of coherence and the formation of a common type interface. Deformation with a degree of 50 % contributed to an increase in the recrystallized volume to 43 %. Along with this, in non-crystallized regions there is a further increase in the misorientation angle between subgrains up to high-angle in some areas where the largest subgrains of size 3-0.3 ^m are present. The average grain size of y and y' phases in the recrystallized regions of the microduplex structure formed during deformation was 6.8 and 1.6 ^m, respectively.
Schematically, the mechanism of transformation of the initial coarse-grained structure into the microduplex one in the process of hot deformation of the superalloy was presented by the author in [14]. It should also be noted that such a mechanism of microduplex structure formation is also characteristic
of other alloys hardened by precipitates of the y'-phase based on the Ni3(Al, Ti) intermetallic compound, including in powder superalloy EP741NP made by powder metallurgy [15].
Table 1
The sizes of recrystallized grains (dY) and their specific volume (V) after hot deformation at 1100 °C by 50 % at a speed of 10-3 s1
Deformation temperature, °C 1000 1050 1100 1125
dy V f/dYha 2.4 7 0.73 3.5 16 0.57 4.6 33 0.36 6.7 69 0.2
Note: f is the volume fraction of the y'-phase; dha is particle size of the y'-phase after heterogeneous annealing.
Table 2
Change in the volume fraction (V) of recrystallized grains from the degree and strain rate at 1100 °C
Deformation Deformation degree, %
rate, s-1
5 10 20 40 50 75 75*
10-3 0 4 11 26 33 53 95
10-5 4 9 18 39 43 76 -
* Fractional deformation for 2 transitions with postdeformation annealing at 1100° C for 2 and 8 hours.
EBSD (Electron Back Scattered Diffraction) analysis of the change in boundary misorientations in the superalloy. Studies of changes in the angles of misorientations of boundaries in the alloy after deformation at 1100 °C with a speed of 7 • 10-4 s-1 with various degrees of deformation (30, 55, 75 %) were carried out.
With an increase in the deformation degree, an insignificant change in the average size of recrystallized grains during deformation is observed under the indicated deformation modes (Table 3). At the same time, the fraction of recrystallized volume increases to 77 %. When analyzing the change in specific misorientations of microregions within elongated initial grains (point to origin misorientation), the following was discovered. The initial grains in the process of deformation undergo less bending than in nichrome (the specific misorientation in grains even after deformation by s = 75 %, the specific misorientation is 0.12°/^m (Table 4). Compared to nichrome, this is a low value. Perhaps, in the superalloy the presence of coagulated particles of the y'-phase prevents the strong bending of the matrix grains and contributes to a more uniform distribution of deformation.
In the process of plastic deformation, the initial grains were divided into regions misoriented relative to each other. With an increase in the degree of deformation from 30 to 75%, the size of these regions decreased from 4.6 ^m to 3.1 ^m (Table 5). With an increase in the degree of deformation, the misorientation between regions slightly increased from 2.4 to 3 deg. It is interesting to note that the size of the microregions approximately corresponded to the size of the subgrains at the corresponding degrees of deformation, the values of which were obtained from electron microscopic studies.
Table 3
Volume fraction (Vrec) of recrystallized grains and their average size after deformation at 1100 °C with a strain rate of 7 • 10-4 s1
e, % 30 55 75
Free, % 28 52 77
4ec, ^m 4 3.76 3
Table 4
Change in specific misorientations within elongated initial grains (point to origin misorientation)
e, % 30 55 75
©/D, °/pm 0.26 0.1 0.12
Note: ©/D, (°/^m), where © is the general misorientation at the end of the initial grain with respect to the beginning of the grain, D is the length of the elongated grain, ^m.
Table 5
The effect of the deformation degree on the size of microregions and the average misorientation between microregions
Deformation degree e, % 30 55 75
The average misorientation 2.4 2.5 3.0 between micro regions inside deformed grains, degrees.
The average size of the 4.6 4.4 3.1 microregions, ^m
No data are available on the change in the proportion of small-angle boundaries (SAB) and highangle boundaries (HAB) at the initial stage of deformation (10-30 %). It is assumed that at this stage the formation of a subgrain structure occurs, which leads to an increase in the fraction of SAB and, accordingly, to a decrease in the share of HAB of the initial grains, as evidenced by the previously cited electron microscopic studies. With an increase in the degree of deformation to 30 and further to 50 and then to 75 %, the proportion of high-angle misorientations gradually increases (Fig. 11). At the same time, the number of small-angle misorientations decreases.
Boundaries share, %
1 |_im
605040302 0-
SAB 01-15°
HAB 15°-65°
10 20 30 440 5 0 (50 70 s, %
a)
b)
c)
d)
Fig. 11. Change in the microstructure (a, b, c) and the share of SAB and HAB (d) in deformation of the superalloy (1100 °C, s = 7 • 10-4 s-1):
a - 30 %; b, c - 75 %
An analysis of the angle misalignment distributions showed (see Fig. 11) that, with an increase in the degree of deformation from 30 to 75 %, the conversion of SAB to HAB is observed. This is evidenced by a decrease in the share of SAB (up to 45 %) and an increase in the share of HAB (up to 48 %). This indicates the formation in the process of hot deformation of small-angle boundaries (sub-boundaries) and their subsequent transformation into high-angle boundaries of new recrystallized grains. In this alloy, a much smaller fraction of special boundaries (4 %) was observed than in nichrome. But the number of special boundaries with an increase in the degree of deformation from 30 to 75 % increased slightly to 7.8 %.
The influence of temperature and strain rate. Similar microstructural changes occurred at higher strain rates. Thus, at strain rates of 10-3 s-1 and higher, a subgrain structure was formed with an increased dislocation density, in which, along with equiaxed subgrains, needle-shaped subgrains were revealed (Fig. 12b).
It should be noted that the appearance of needle-shaped subgrains in the studied temperature and speed range corresponded to the state of the material
—3 —1
deformed at speeds of 10 s and higher and with degrees exceeding the maximum allowable before the formation of first cracks on the surface of the samples.
With an increase in the degree of deformation, the misorientation between subgrains increases, although less intensively than at 1100 °C. This is probably because, even after intense deformation at 1000 °C, by 72 % at a rate of 10-3 s-1, small-angle misorientation remained between subgrains. Only at a lower strain rate of 10-5 s-1 does the subgrain structure transformed into a grain structure with high-angle interphase and grain boundaries, however, the fraction of the recrystallized volume was insignificant (Fig. 13).
A decrease in the deformation temperature from 1100 to 1000 °C promoted the formation of a less perfect subgrain structure. An increase in the deformation temperature to 1125 °C led to a significant decrease in the volume fraction of the y ' -phase and the intensive development of recrystallization. At the same time, the resulting microstructure was heterogeneous. In addition to grains of the y ' -phase of size 5 pm, grains up to 40 pm are found in the microstructure.
Since, as a result of hot deformation, 75 % of the samples in state 2 did not recrystallize the entire volume of the material, it was important to evaluate the effect of post-deformation annealing, as well as the fractional deformation and changes in the strain rate. It turned out that even conducting short-term annealing for 30 minutes contributed to the occurrence of subgrain coalescence processes, a decrease in the dislocation density, and an increase in the recrystallized volume.
Conducting deformation in 2 transitions, for
—3 —1
example, at a speed of 10 s and a total degree of
75 % at 1100 °C and subsequent annealing for 2
and 8 hours made it possible to form a micro duplex
microstructure in almost the entire volume of the
material (Table 2, Fig. 12). In addition, if the initial
strain rate is increased by 2 orders of magnitude at
1100 °C after the degree of s = 50 % (Vrec = 39 %) to —2 —1
10 s , the recrystallized volume can be increased to 80 %, thereby exceeding it by 20 % compared with a material deformed to the same degree of deformation at a speed of 10-4 s-1.
The results of the study of the process of transformation of a coarse-grained matrix microstructure into the microduplex-type structure make it possible to reveal its laws and features caused by the state of the Y -phase and processing conditions.
a) b) c) d)
Fig. 12. The microstructure of the superalloy after deformation at 1100 °C at a speed of 10-3 s-1 at s = 75%:
a, b - sediment for 1 transition; c, d - sediment for 2 transitions with post-deformation annealing at 1100 °C for 2 and 6 hours
A comparative analysis of the mechanical properties and microstructural changes of the superalloy in two states indicates that heterogenization of the microstructure, leading to the enlargement of y' -phase particles, an increase in the distance between them and the transformation of y/y' coherent boundaries into partially coherent ones, significantly improves plasticity and reduces voltage flow and creates the most favorable conditions for the formation of microduplex microstructure.
The positive effect of heterogenization of the microstructure is manifested primarily in a change in the nature of dislocation slip under both cold and hot deformation. The material with dispersed y'-phase particles completely coherent with the matrix (state 1) is characterized by rough slip, which, according to [3], is caused by cutting of particles by dislocations and a small number of slip systems involved [12]. This is primarily evidenced by the rough slip traces in the cold-
deformed material, and in the hot-deformed one, the oriented growth of y'-phase particles, observed according to [16] in the case of slip development mainly in the same crystallographic plane. Meanwhile, in an alloy in state 2 with coagulated y' -phase particles, when the distance between the particles increases noticeably and the interphase boundaries become partially coherent, multiple slip develops, which contributes to the formation of a disoriented cellular (after cold deformation) or dynamically polygonized (after hot deformation) structure.
Such significant differences in the nature of dislocation slip in the alloy in states 1 and 2 lead to significant differences in the kinetics of recrystallization processes and, accordingly, the production of various microstructures. It should be noted that the formation of the microstructure occurs under the conditions of the interaction of the processes of recrystallization and the release of the y' -phase, the
ratio of which is significantly affected by the state of the y'-phase. In the alloy in state 1, when the formation of recrystallization nuclei is difficult, the leading role is played by precipitation processes, which at low strain rates (10-5 s-1) lead to the formation of grain cells with coherent particles of the y' -phase lamellar shape in the boundary regions, and at higher speeds they cause the development of recrystallization only in zones free of precipitates. At the same time, in state 2, the development of multiple slip contributes to the formation of numerous recrystallization nuclei, as a result of which a microduplex microstructure is formed in the material.
The transformation of a coarse-grained microstructure into a microduplex one during hot deformation observed in a material with coagulated y -phase precipitates occurs by a mechanism that differs significantly from that described in [1, 3, 4, 17]. As follows from the experimental results, in contrast to static during dynamic recrystallization, the formation of high-angle interphase and grain boundaries occurs through the interaction of lattice dislocations with grain and phase boundaries and is due to the action of specific mechanisms, including superplastic deformation.
In the process of hot deformation in the SP temperature-velocity regime of the alloy under study [18], the dislocation is introduced and accumulated on the partially coherent interface y' -phase (see Fig. 10), which is an effective barrier to the thermally activated dislocation motion [19]. This leads to a further violation of coherence up to its complete loss and to an increase in matrix distortion near y'-phase particles. Inhibition of lattice dislocations by partially coherent precipitates of the y -phase, significant local lattice stresses near the particles of the y -phase activate multiple intracranial dislocation slip, which causes the formation of a subgrain structure in the matrix. Simultaneously with the conversion of the initial low-energy partially coherent y -phase interfaces to incoherent ones with an increase in the degree of deformation, the misorientation angle between the subgrains also increases and gradually transforms into grains with high-angle boundaries. It is important to emphasize that the introduction of lattice dislocations into the interphase boundary, as well as the large extent of small-angle boundaries, apparently contribute to the activation of diffusion processes [1, 20], as a result of which rearrangement of partially coherent interphase boundaries to incoherent and enlargement of y -phase particles occurs more rapidly. This assumption is supported by quantitative analysis, according to by which the growth rate of y -phase particles in a hot-
deformed (or upon annealing cold-deformed) material is more than 2 orders of magnitude higher than the growth of particles in an undeformed matrix.
The formation of a chain of small recrystallized grains along the initial grain boundaries of the matrix and the deformation conditions corresponding to the SP temperature-velocity regime of the material stimulate the action of a specific deformation mechanism, the grain boundary sliding [1, 21]. Indeed, at small degrees of deformation (10-20 %), when the recrystallized volume is small (< 18 %), the grain boundary sliding is difficult. With an increase in the degree of deformation, the grain boundary sliding develops more intensively, as evidenced by the growth of the coefficient m to 0.31. This gives rise to local stresses at high angle boundaries, which are effective sources and sinks of lattice dislocations [1, 21], and, apparently, can contribute to the generation of lattice dislocations in subgrains. An increase in flow stress on the a-s curves is probably associated with an increase in the dislocation density in subgrains, and a decrease in stress with the transformation of small-angle boundaries of subgrains into high-angle ones, as a result of which the dissociation of lattice dislocations is sharply accelerated, which in turn lead to the development of stimulated grain boundary sliding [1, 21].
It should be noted that the above-mentioned features of the formation of a microduplex structure were observed under certain processing conditions -in the temperature range of 1000-1125 °C with a ratio of the volume fraction of the y' -phase to its size after heterogenizing annealing from 0.2 to 0.73.
Meanwhile, the mechanism of structure formation considered above during deformation in the same temperature range was almost completely suppressed when a dispersed coherent y' -phase (state 1) with a ratio fldy > 0.73 was present in the material. In this case, the formation of a microduplex occurs only after intense deformation (80-99 %) by a different mechanism, which consists in the dissolution of small particles of the y' -phase in migrating boundaries, followed by the release of (y' forming elements on its growing grains [1, 3, 4, 22]. At temperatures close to Tp.r у, when fldy < 0.2, dynamic recrystallization developed extremely intensively, which was consistent with the data of [4], in which a transition from slow to accelerated recrystallization, in biphasic alloys with less than 2 times ratio fid = 0.1. Obviously, due to the small volume fraction of the у' -phase at the deformation temperature, a subgrain structure with very large cells will form, which will transform it into a micro-structure of the matrix type in which most of the
y ' -phase it will stand out inside recrystallized grains during cooling from the deformation temperature.
Along with the state of the y' -phase, the conditions of hot deformation have a significant effect on the formation of a microduplex structure. It is interesting to note that after intense deformation by 75 %, recrystallization did not cover the entire volume of the material (see Tables 1 and 2), also obviously due to the localization of the SP deformation in the fine-grained zone.
As the experiments showed, in order to obtain a completely recrystallized microstructure in the material, deformation at the final stage (> 50 %) was
preferably carried out in the SP third velocity range
-2 -1 -1
(10 -10 s ). This is explained by the fact that during the transition to the SP third velocity interval, the action of the main mechanism of the SP deformation -the grain boundary sliding becomes more difficult, but the role of the Burgers dislocation increases, which provides additional pumping of the dislocation into subboundaries and their subsequent transformation into
high angle boundaries. At the same time, at increased
-2 -1 -1
strain rates of 10 -10 s , the intensity of the development of dynamic recrystallization substantially decreases, while in the non-crystallized regions a substructure with an increased dislocation density and an unfavorable needle-shaped subgrain is formed. In this case, fractional deformation with post-deformation annealing is very effective, as a result of which there is a further development of processes that began during dynamic recrystallization, which is consistent with the ideas presented in [1, 23]. Since the processes occurring in fractional processing, which are characteristic of both static and dynamic recrystallization, simultaneously with the mechanism of micro duplex formation considered above by transforming the subgrain structure into a grain and partially coherent particles of the y -phase into incoherent particles, another mechanism can act in which the formation of recrystallization nuclei due to the splitting of the interphase boundary between the coagulated particle of the y' -phase and the y-phase.
Thus, the studies showed that the development of recrystallization in a nickel alloy is significantly affected by the size of the y -phase, the distance between particles of the y -phase and the state of the y/y' interfaces. Preliminary heterogenization of the microstructure, which leads to an increase in the particle size of the y' -phase, interparticle distance, and the transformation of coherent interphase boundaries into partially coherent subsequent processing under temperature-velocity conditions of the joint venture, provide the most favorable conditions for obtaining a
microduplex structure in a highly alloyed nickel alloy. In the course of dynamic recrystallization, a subgrain structure is formed and it gradually transforms into grains with high-angle boundaries, while partially coherent precipitates of y -phase particles are transformed into incoherent grain particles with an arbitrary orientation relative to the matrix grains.
The study of the microstructure and properties of the nickel alloy after hot deformation and EBSD analysis was supported by the Russian Science Foundation (grant RSN No. 18-19-00685), as well as the study of the structure and properties after cold deformation and annealing, fractional deformation was also performed as part of the IPSM RAS State Assignment No. AAAA-A17-117041310215-4. Experimental studies were performed on the basis of the Center for the collective use of scientific equipment of the IPSM RAS.
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