INTRODUCTION
Recent studies indicate that severe plastic deformation (SPD) of aluminum alloys leads to the formation of nano- and ultrafine-grained structures with a regulated distribution of alloying elements, including secondary strengthening phases, nanoclusters, and grain boundary segregations or films, resulting in a unique combination of properties [1‐9]. Several works have demonstrated that refining grain size to ultrafine (UFG) or nano-sized ranges in various alloys enables superplasticity (SP) to be achieved at low temperatures [10‐16]. The 7000 and 2000 series alloys nanostructured by SPD techniques have recently been found to retain their high strength properties after deformation under low temperature SP conditions within a certain temperature–strain rate range [13, 14].
The possibility of achieving record strength levels through the formation of a UFG structure with nanostructured second-phase precipitates was first demonstrated using the example of commercial 7475 and 7075 alloys of the Al–Zn–Cu–Mg system [13, 14]. SPD-formed UFG structure with grain boundary precipitates of alloying elements in the form of interlayers has been shown to promote SP behavior at temperatures below 200°C (0.5Tm) [13, 14]. This reduction in forming temperature, combined with ideal strain rate conditions, allows the alloy to retain its strength after forming, typically 20–30% higher than analogous alloys hardened by conventional treatments [8].
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In [9], Al–Cu–Mg 2024 UFG alloy with an average grain size of 150 nm, which was produced by high-pressure torsion (HPT) at room temperature (RT), exhibited SP behavior at 400°C (0.74Tm) and a strain rate of 10–3 s–1. The maximum relative elongation to failure was 400%. The study [14] indicates that 2024 alloy in NS state, having a grain size of 100 nm, demonstrates an equivalent maximum elongation (about 400%) when deformed at a reduced temperature of 270°C (0.56Tm). The study found that the strength of the NS alloy remained high even after deformation under low-temperature SP conditions.
Alloys 1421 and 1570 of the Al–Mg–Li–Sc and Al–Mg–Sc systems, respectively, showed high-strain rate and low-temperature SP in UFG and NS states at temperatures ranging from 200 to 400°C and strain rates between (5 × 10–3–2) s–1 [15]. Ultimate elongations to failure of 1500 and 1100% were achieved in the UFG and NS states, respectively, at a strain rate of 0.1 s–1 at 400°C in the 1421 alloy. High plasticity (300%) of UFG 1570 alloy with 200-nm grain size at 200°C was observed [15]. Increasing the deformation temperature in the range of 250–300°C shifted the plasticity maximum to the high strain rate region. The authors have identified significant characteristics regarding the deformation behavior of these alloys, including the absence of porosity formation, an extended phase of strain hardening, and a weak correlation between the strain rate sensitivity coefficient and the strain rate. These factors suggest that the main contribution to specimen deformation is provided by cooperative grain boundary sliding along the fragments spanning groups of grains [15]. High-speed forming at low temperatures avoids grain growth and results in improved mechanical characteristics of the products [8, 15].
From the above, it can be concluded that investigation of UFG and NS aluminum alloys at lower SP temperatures holds scientific and practical importance. The identified characteristics of the mechanical properties of alloys with this microstructure open up new prospects for the production of high-strength and lightweight products and structures, such as composite materials (CM) using aluminum and its alloys [17, 18].
More recently, based on the 5182 alloy, a promising aluminum–magnesium 1565ch alloy (high in Mg and alloyed with Zn and Zr) has been developed and is now widely used in the transportation industry. This alloy maintains high strength, corrosion resistance, and ductility despite containing a relatively high magnesium content [19]. Zirconium is necessary to form an Al3Zr phase in the 1565ch alloy, serving as nuclei for solid solution decomposition [20].
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Studying the aluminum 1565ch alloy in both UFG and NS states is of scientific and practical interest. The former relates to the extension of ideas regarding the nature of SP at lower temperatures, while the latter pertains to the potential use of this material as a matrix for producing aluminomatrix fiber composites [17, 18]. The exceptional ability of SP materials to effortlessly fill micron-sized gaps (or gaps of about ten microns in size) is expected to significantly reduce the temperature and strength requirements for compacting CM, thereby simplifying the manufacturing process and improving the quality of the resulting products.
The primary objective of this study is to investigate the mechanical behavior of commercial 1565ch alloy from the Al–Mg system in NS and UFG states, specifically at room and elevated temperatures. The study aims to identify the stability threshold of its strength properties after thermal and thermal mechanical treatment and to determine the temperature–strain rate range of SP manifestation.
EXPERIMENTAL
Commercially available formable, non-age hardenable Al–Mg 1565ch alloy (Al–5.66Mg–0.81Mn–0.67Zn–0.09Zr–0.07Cr–0.04Ti–0.001Be–0.3(Fe + Si) (wt%) according to GOST 4784–2019) was chosen as the material of the study.
To achieve precipitation hardening and maximum refinement of the grain structure through SPD and/or work hardening, the initial alloy billets were annealed at 480 ± 5°C for 1 hour and then rapidly quenched in water [21].
SPD deformation of the 1565ch alloy:
To form a nanostructure, some of the billets in the form of discs with a diameter of 20 mm and a thickness of 1.4 mm were subjected to HPT at RT at a specific pressure of 6 GPa and a strain rate of 1 rpm. The total number of revolutions of the movable anvil is ten. For a description and schematic of the HPT, see [22].
Some of the billets used to form the UFG structure underwent six cycles of equal-channel angular pressing via the Conform (ECAP-Conform) method at 200°C, in accordance with the Bc mode on the ECAP-C-01NM laboratory machine. Specimens with a 10 × 10 mm square cross section and a length of 110 mm were employed. A detailed ECAP-Conform scheme and description can be found in [23].
The microstructure was analyzed using transmission electron microscopy (TEM) employing a JEOL JEM-2100 electron microscope with an accelerating voltage of 200 kV.
Metallographic analysis was carried out using a Carl Zeiss Axio Observer A1m microscope.
A JEOL JSM-6490LV scanning electron microscope (SEM) equipped with an INCA chemical analysis attachment for energy dispersive X-ray spectroscopy (EDX) was used to study the initial state, fracture surfaces, and strain patterns after mechanical testing.
X-ray phase and structural analysis (XRD) of alloy specimens were conducted based on results of measurements taken with the Bruker D2 Phaser diffractometer using CuKα radiation (voltage and current set to 30 kV and 20 mA, respectively). X-ray analysis was performed using the Rietveld full profile refinement method implemented in the MAUD software [24]. The instrumental broadening was accounted for by analyzing the Al2O3 Standard sample. The lattice parameter values (a), the size of coherent scattering domains (dXRD), and the microstrains in the crystal lattice \(\left( {{{{\langle {{\varepsilon }^{2}}\rangle }}^{{1/2}}}} \right)\) were calculated. The dislocation number density (ρ) was calculated from these values according to equation (1) [25]where b = √2/2a is the Burgers vector modulus.
$$\rho = 2\surd 3{{\langle {{\varepsilon }^{2}}\rangle }^{{1/2}}}/({{d}_{{{\text{XRD}}}}}b),$$
(1)
Mechanical Properties
Mechanical testing was performed on an INSTRON 5982 universal testing machine and using Bluehill 3 software. The testing was performed at RT, 200, 250, and 300°C in the strain rate range 5 × 10–5–10–2 s–1 to determine the flow stress, elongation at fracture, and strain rate sensitivity coefficient.
$$m = \frac{{d\ln \sigma }}{{d\ln \dot {\varepsilon }}},$$
(2)
where σ is the flow stress at 100% strain and \(\dot {\varepsilon }\) is the strain rate.
The size of the gauge portion of the specimens was 1.0 × 0.8 × 3.2 mm3.
Vickers microhardness (HV) was measured on the EMCO-TestDuraScan 50 instrument under a load of 1 N for 15 s.
RESULTS
Microstructure
The microstructure of the 1565ch alloy in its initial coarse-grained (CG) state after rolling is characterized by elongated grains resembling fibers with an average width of 2.3 ± 0.2 μm. Additionally, there are coarse casting-induced particles up to 20 μm in sizes in the microstructure (see Fig. 1) [26]. Annealing at 480°C for one hour resulted in the formation of a fully recrystallized structure [26‐28]. Two types of particles were observed in the aluminum grains, with particle sizes ranging from 2–10 μm. According to EDX analysis (Fig. 1b) and data found in the literature [26‐29], the light-gray particles are identified as Al6(Mn,Fe) phase and the black particles are identified as Mg2Si phase.
Fig. 1.
Microstructure of the 1565ch alloy: (a) in the initial CG state after rolling and (b) after heterogenization annealing. The white-vizualized particles are identified as Al6(Mn,Fe) particles, while the black-vizualized particles are Mg2Si particles.
To determine the optimal temperature and strain-rate conditions for the occurrence of SP deformation at reduced temperatures in the 1565ch alloy, two structural states were formed using different SPD techniques.
After HPT processing, a homogeneous nanostructure with an average grain size of 95 ± 5 nm and a grain elongation coefficient of 1.1 was formed in the 1565ch alloy specimens (Fig. 2). Individual Al3Zr phase precipitates are observed within and along the boundaries of aluminum grains, as is indicated by arrows in Fig. 2 [26, 27].
Fig. 2.
Microstructure of the 1565ch alloy after HPT treatment at room temperature. Arrows indicate the Al3Zr phase: (a) bright field with diffraction pattern, (b) dark field.
The alloy 1565ch achieved a homogeneous UFG state with an average grain size of 200 ± 5 nm and a grain elongation coefficient of 1.3 (Figs. 3a, 3b) and 200 ± 3 nm, grain elongation coefficient of 1.1 in the longitudinal and cross-section of the billet (Figs. 3c, 3d) after ECAP-Conform processing at 200°C. Multiple precipitates of Al6Mn and Al3Zr phases are observed within and along aluminum grain boundaries according to [26, 27], as is indicated by arrows in Fig. 3. Al3Mg2 β-phase particles, judging by morphological features, form in the structure due to dynamic decomposition of supersaturated solid solution [16, 29, 30].
Fig. 3.
Microstructure of 1565ch alloy after ECAP-Conform treated at 200°C. Black arrows indicate the β phase and white arrows, Al6Mn: (a), (c) bright field with diffraction pattern, (b), (d) dark field. Microstructure in (a), (b) longitudinal and (c), (d) cross sections.
These observations are in good agreement with the XRD data.
X-ray profiles taken from specimens of the 1565ch alloy in various structural states are presented in Fig. 4. Indexing the peaks indicates that the phase composition of the alloy in its initial state consists of intermetallic particles, including Mg2Si and Al6(Mn,Fe), which is consistent with the findings of SEM/EDX analysis and literature data on the composition of this alloy heat-treated at T > 400°C [19]. In addition, the XRD analysis indicated that the UFG-structured alloy formed through ECAP-Conform treatment at 200°C exhibits diffraction peaks corresponding to Al3Mg2 β-phase particles [19‐21, 29, 30]. The presence of Al3Mg2 β-phase particles has been previously identified in Al–Mg alloys with magnesium content exceeding 4 wt %, such as in the UFG-structured alloy 1560 formed via ECAP under similar temperature conditions [31‐33].
Fig. 4.
X-ray diffraction patterns taken from the 1565ch alloy specimens and analysis of its phase composition in the initial state and after SPD by HPT and ECAP-Conform techniques.
The use of SPD through HPT at RT does not lead to a qualitative change in the initial phase composition. However, the peak intensity corresponding to intermetallic phases such as Mg2Si and Al6(Mn,Fe) is significantly reduced. This may indicate a reduction in the size of large Al6Mn particles presented in the starting material due to SPD. The reliable confirmation of β‑phase formation during HPT at RT is beyond the sensitivity limit of X-ray diffraction.
The results of the structural parameter calculations of the 1565ch alloy in its initial and deformed states show that the lattice parameter of the alloy in the undeformed state (a = 4.0785 Å) is significantly higher than the characteristic value for pure aluminum (a ~ 4.050 Å). This can be explained by the formation of a solid solution of magnesium in the aluminum matrix caused by the thermal treatment. It is known that the dissolution of one atomic percent of magnesium in aluminum leads to a lattice parameter increase of 0.0046 Å [23]. Therefore, the magnesium content in the solid solution before deformation can be estimated to be ~6.2 at %, which agrees well with the magnesium content in the chemical composition of the alloy (5.66 wt % ~ 6.24 at % of magnesium in Al–Mg alloy). Interestingly, HPT performed at RT results in a considerable decrease of the lattice parameter in the alloy (from 4.0785 to 4.0764 Å), indicating that some magnesium (~0.4 at %) has escaped from the solid solution. The lack of β-phase peaks indicates that the decomposition of the solid solution did not lead to the precipitation of intermetallic particles, but instead caused heterogeneous grain boundary segregation of magnesium atoms. This was similarly observed in alloy 1570 after HPT [3, 5]. The above assumption is in line with the very high strength of the 1565ch alloy after HPT (flow strength of 710 MPa). This value significantly exceeds the value predicted for this grain size based on the Hall–Petch relationship [34]. The phenomenon can be explained by the presence of segregations/clusters of alloying elements at grain boundaries in the NS state HPT. These segregations/clusters hinder the emission of dislocations from the grain boundaries, which increases the stress required for plastic flow in this material [35, 36].
SPD using the ECAP-Conform technique at 200°C also leads to a reduction of the lattice parameter of the alloy to 4.0760 Å, lower than that achieved by SPD using HPT (which is 4.0785 Å). In this case it indicates the decomposition of the solid solution, accompanied with the observed formation of the β phase, in agreement with relevant literature data [29, 31‐33]. The β phase was shown [37] to form during annealing of alloy 5083 at 175°C for 10 days on previously existing manganese-rich particles and at grain boundaries, indicating that these sites serve as effective heterogeneous nucleation centers. In addition, ECAP processing of Al–7Mg alloy at RT and subsequent SP deformation at 300°C resulted in the formation of submicron β-phase particles along grain boundaries, which precipitated in the initial phase of the deformation process and coarsened during SP deformation. In addition, β-phase nanoparticles formed inside grains, which were attributed to grain boundary migration [16, 31‐33]. Nano and submicron Al3Mg2 precipitates can restrict dislocation and grain boundary motion, delaying complete recrystallization and inhibiting grain growth.
The coherent domain size determined by X-ray diffraction matches well with the corresponding grain size measured by transmission electron microscopy (TEM) techniques. After HPT, it was found to be dXRD = 69 nm (with a TEM measured value of d = 95 nm) and dXRD = 140 nm (with a TEM measured value of d = 200 nm) after ECAP-Conform. The difference between dislocation number-density values after HPT and ECAP, with a more than two-fold increase in the former, can be explained by the differences in SPD type, temperature, and degree of deformation.
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Mechanical Properties at Room Temperature
Forming an NS state by RT HPT resulted in reaching a flow strength of 715 ± 5 MPa and a tensile strength of 800 ± 6 MPa at an elongation to fracture of 1.2 ± 0.2% (Fig. 5). The strength properties of the NS alloy are twice as high as those of the alloy in a coarse-grained state after standard cold rolling, with a flow strength of 330 ± 3 MPa, a tensile strength of 380 ± 6 MPa at a relative elongation to fracture of 16 ± 0.2%. The microhardness of the NS 1565ch alloy was 237 ± 4 HV. For comparison, the microhardness of the material in the coarse-grained state (CG) after rolling was estimated to be 110 ± 2 HV.
Fig. 5.
Room-temperature tensile curves of UFG and NS specimens.
Forming a UFG state by ECAP-Conform at 200°С resulted in reaching a flow strength of 400 ± 5 MPa and a tensile strength of 450 ± 6 MPa at an elongation to fracture of 18 ± 0.5% (Fig. 5). The microhardness of the UFG 1565ch alloy was 158 ± 4 HV.
To investigate the thermal stability of the NS and UFG 1565ch alloy, the deformed specimens were annealed for one hour in the temperature range of 200–300°C, in which recrystallization processes do not occur [26, 27].
Figure 6 illustrates the dependency of the microhardness on the annealing temperature for the NS and UFG alloy states.
Fig. 6.
Microhardness vs. annealing temperature for 1 hour: NS, UFG, and CG states.
Annealing at 250°C was found to reduce the microhardness of the NS alloy. Annealing at 300°C causes a significant softening of the NS alloy. The reduction in microhardness can be explained by the processes of recovery, structure perfection, and grain growth [1‐8, 14].
The alloy in the UFG state with a grain size of 200 nm is slightly softened during thermal annealing in the temperature range of 200–250°C, and a significant softening of the UFG alloy is observed after annealing at 300°C.
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Mechanical Properties at Increased Temperatures
After establishing the thermal stability range, the examination of SP behavior of the 1565ch alloy was conducted in both NS and UFG states. The investigation was tested at constant rates of 5 × 10–5–10–2 s–1 at temperatures of 200, 250, and 300°C.
Figure 7 shows the flow stress versus strain rate curves of NS (Fig. 7a) and UFG (Fig. 7b) alloy specimens at temperatures of 200, 250, and 300°C, respectively, for a strain rate of 5 × 10–5–10—2 s–1.
Fig. 7.
Flow stress versus strain rate for the alloy with (a) NS and (b) UFG structures.
The curves of flow stress versus strain rate in logarithmic coordinates are typical of superplastic materials (Fig. 7). As can be seen, the flow stress decreases with increasing temperature and decreasing strain rate. At 200°C, the flow stress is weakly dependent on strain rate for both states. The coefficient m does not exceed 0.18. The elongations to fracture are 100–120% in the NS state and 130–160% in the UFG state.
The range of optimal SP strain rate values exhibiting maximum rate sensitivity corresponds to temperatures between 250–300°C. Figure 8 shows the strain curves at these temperatures for various constant strain rates.
Fig. 8.
Tensile curves of alloy specimens in (a), (b) NS and (c), (d) UFG states, the rate sensitivity coefficients at temperatures of (a), (c) 250 and (b), (d) 300°C are indicated in the curves.
The maximum strain rate sensitivity coefficient (m) for the NS alloy at a deformation temperature of 250°C is in the range of 0.32–0.35. The maximum elongation is achieved at a strain rate of 5 × 10–3 s–1, reaching a value of 300%. The elongation increases up to 500% at the same strain rate at a temperature of 300°C. Simultaneously, the temperature increase is followed by a considerable drop in the flow stress (by 50%), and the parameter m equals 0.28.
For the UFG alloy, the highest strain rate-sensitivity coefficient is 0.24 at a deformation temperature of 250°C. It also reaches a maximum ductility of 150% at a strain rate of 5 × 10–4 s–1. The maximum flow stress is exceptionally high. It is about 170 MPa. Increasing the temperature for SP deformation to 300°C enhances ductility up to 560% at a strain rate of 10–2 s–1. This, however, causes the flow stress to decrease by a factor of 2.5 to 60 MPa, with parameter m equal to 0.73. A comparable value of m (approximately 0.75) was observed for Al–7Mg alloy deformed under SP conditions at 300°C in the UFG state [16].
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Mechanical Properties of the UFG Alloy at Room Temperature after SP Deformation
To assess the maintenance of stable high-strength properties in a UFG alloy after SP deformation, we measured its microhardness and conducted RT tensile tests on specimens that had undergone thermal mechanical treatment at temperatures of 250 and 300°C and strain rates of 10–3 s–1. The test material achieved a significant elongation to fracture of 100 and 500% at the parameter m = 0.32–0.73 through the implementation of these thermal-mechanical treatment (TMT) conditions.
Microhardness curves with respect to temperature of SP deformation are depicted in Fig. 9a, and values were obtained from the fractured specimen’s gauge part. Even after deformation at 250°C, the NS 1565ch alloy maintains high microhardness values.
Fig. 9.
(a) Microhardness vs. SP deformation temperature (SPD), (b) Tensile curves at RT for the UFG 1565ch alloy after SP deformation.
The UFG stress-strain curves before and after SP deformation are presented in Fig. 9b. Analysis of the curves indicates that UFG 1565ch alloy retains its high-strength state of 440 ± 3 MPa after SP deformation at 250°C with a ductility of 7.0 ± 0.3%. The results obtained show no hardening during SP deformation in the temperature range of 250–300°C for UFG alloy 1565ch.
Detailed analysis of the fracture surfaces (Fig. 10) enabled us to determine that at areas far from the primary pores, the grains, along with the movable groups of combined grains, maintain their equiaxed shape, implying deformation in the SP regime due to the grain-boundary sliding (GBS) mechanism [38]. The final grain size does not exceed 1 μm. As pores appear, SP flow conditions are perturbed, causing grains to elongate between the main pores, indicating a transition to an intragranular slip mechanism. This is especially noticeable in places where there are fewer than ten grains packed into the space between the main pores of the grain. In some cases, pore enlargement occurs through localized deformation via the GBS mechanism, while the UFG structure sustains the SP regime between the main pores elsewhere. However, local deformation does occur in this scenario. Fractographic studies revealed characteristic pitted fractures in the SP materials without any detectable particle fractures (see Fig. 11).
Fig. 10.
Pores on the deformation relief of the UFG specimen after SP deformation at 300°C, a strain rate of 10–3 s–1; (1) localized deformation by GBS mechanism and (2) transition to intragranular slip.
Fig. 11.
Fracture surface of the NS specimen after SP deformation at: (a) 200 and (b) 300°C.
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DISCUSSION
The formation of NS and UFG structures with predominantly high-angle boundaries containing nanoscale precipitates of alloying elements, both within the aluminum grain interiors and along its grain boundaries, has resulted in achieving high strength of 450–800 MPa in the 1565ch alloy.
The results obtained indicate that the NS and UFG 1565ch alloy shows features of SP behavior at temperatures between 250–300°C and strain rates of 5 × 10–4 to 1 × 10–2 s–1. The alloy exhibits a high strain rate sensitivity coefficient with no hardening during deformation, and elongation reaches 170–560% at low flow stress. A comparison of two different alloy states, NS (95 nm) which were formed at RT and UFG (200 nm) which were formed at 200°C, showed their similar SP behavior at relatively low temperatures. For example, alloy 1565ch with an average grain size of 4–10 μm after multidirectional forging at 350°C and cold rolling showed SP behavior only at temperatures above 500°C [39]. The maximum relative elongation of 530% occurred at a strain rate of 1 × 10–3 s–1 and a temperature of 540°C, with a flow stress not exceeding 7 MPa (m = 0.66). At a constant strain rate of 5 × 10–3 s–1 and a temperature of 520°C, the elongation was 400%.
The remarkable homogeneity and stability of the initial NS and UFG structures, as compared to the structural features shown in research where alloy 1565ch in the fine-grained state exhibited high temperature SP [40, 41], is one of the main reasons for the increased elongation values and decreased SP temperatures. According to [38, 42] fine grains facilitate GBS and its accommodation by other mechanisms, hence superplastic behavior is observed at higher strain rates and lower temperatures. Segregation of dissolved magnesium along the nano/ultrafine grain boundaries in the initial structure may have an additional influence, increasing the resistance to grain growth by reducing the energy of the boundaries and limiting their mobility due to pinning by impurity atoms. In addition, nano and submicron-sized Al3Mg2 precipitates may restrict grain boundary migration, thereby playing an essential role in inhibiting dynamic grain growth and ensuring stable flow when undergoing SP deformation. The NS/UFG alloy exhibits relative elongations to fracture with a distinctive maximum of 500–560% at 300°C, which can be credited to the optimization of a set of conditions required for the SP implementation [7, 43].
The results obtained are highly relevant for the practical application of the SP effect in high strain rate forming technologies for mass production scenarios.
CONCLUSIONS
(1) The NS and UFG states were formed with an average grain size of 95 and 200 nm, respectively, and a controlled distribution of particles both in the grain body and along the Al/Al boundaries, by TMT, including annealing at a temperature of 480°C, severe plastic deformation by HPT at RT, and ECAP-Conform at 200°C.
(2) The formation of UFG and NS states in the 1565ch alloy enabled the achievement of high tensile strength in the range of 450–800 MPa.
(3) Alloy 1565ch exhibits superplastic behavior in the temperature range of 250–300°C and strain rate range of 5 × 10–4 to 1 × 10–2 s–1, in the NS and UFG states.
(4) A detailed study of strain topography features has shown that low-temperature SP deformation occurs due to GBS and intragranular dislocation slip and is accompanied by pore formation.
(5) The analysis of the results shows the optimal TMT modes (providing optimal microstructural parameters) for achieving low temperature SP and maintaining high properties in alloy 1565ch.
ACKNOWLEDGMENTS
All studies were conducted at the Nanotech Collaborative Access Center at the Ufa University of Science and Technology.
CONFLICT OF INTEREST
The authors declare that they have no conflicts of interest.
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Translated by T. Gapontseva