A heat treatment method for synergistically improving the strength-ductility-toughness of ultra-high strength stainless steel

Abstract

The application provides a heat treatment method for synergistically improving the strength-plasticity-ductility of ultrahigh-strength stainless steel, and belongs to the technical field of metal material heat treatment, and comprises the following steps: 1) high-temperature solid solution treatment; 2) quenching treatment; 3) first cold treatment; 4) first aging treatment; 5) second cold treatment; and 6) second aging treatment. Through the multistage temperature gradient and time gradient characteristic'solid solution-quenching-deep cooling-aging-deep cooling again-aging again' composite heat treatment method, the continuous heat-cooling cycle is applied in multiple temperature zones and different holding time, the whole process regulation of defect configuration, element distribution and phase change behavior in the ultrahigh-strength stainless steel is realized, and thus the optimal microstructure with coordinated strength-plasticity-ductility is obtained.

Description

Technical Field

[0001] This invention belongs to the field of heat treatment technology for metallic materials, and particularly relates to a heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel. Background Technology

[0002] The heat treatment processes for precipitation-hardening ultra-high strength stainless steel and martensitic aging stainless steel are currently limited to the traditional three-step method: solution treatment (initial), cold treatment (intermediate), and aging treatment (final). Current efforts to optimize the service performance (mechanical properties, corrosion resistance, etc.) of existing steel grades by controlling the microstructure through heat treatment processes focus on the effects of heat treatment parameters—holding temperature, time, and heating / cooling rate—on matrix supersaturation, lath martensitic matrix substructure size, and the spatial distribution and composition of residual / reverse austenite and nanoprecipitates. However, the aforementioned work struggles to significantly improve the overall mechanical properties of materials through the control of single parameters, often resulting in a mismatch between the matrix substructure, strengthening phases, and toughening phases. This leads to a severe trade-off between strength and plasticity, and toughness and corrosion resistance in ultra-high strength stainless steel. Summary of the Invention

[0003] In view of this, the purpose of this invention is to provide a heat treatment method that synergistically enhances the strength, plasticity, and toughness of ultra-high strength stainless steel. This invention introduces a gradient, multi-stage hot-cold cycle to synergistically regulate the content, spatial distribution, and composition of the martensitic hierarchical structure (original austenite grains, martensite lath bundles, martensite blocks and lath sizes, dislocation density), nano-precipitated strengthening phases, and thin-film austenitic toughening phases in ultra-high strength stainless steel. The method provided by this invention not only significantly refines the hierarchical structure of the lath martensite matrix but also effectively retains a high dislocation density. It induces the formation of high-density, diffusely distributed nanoscale reinforcing particles and a suitable amount of fine, uniform thin-film austenite phase in the supersaturated matrix. The nanoprecipitates, acting as the main obstacle to dislocation movement, can produce a significant age-hardening effect. The ductile austenite phase can absorb crack propagation energy, deflect crack paths, and provide additional plasticity through the TRIP effect. The retention of a high dislocation density enhances work hardening capacity, promotes dislocation cross-slip, makes plastic deformation more uniform, and delays necking and subsequent rapid instability fracture. Through the synergistic optimization of these multiple microstructures, this invention achieves a simultaneous improvement in the strength, plasticity, and toughness of ultra-high-strength stainless steel, overcoming the unavoidable performance trade-off between strength, plasticity, and toughness in traditional heat treatment processes.

[0004] To achieve the above-mentioned objectives, the present invention provides the following technical solution: This invention provides a heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel, comprising the following steps: 1) High-temperature solution treatment: Heating ultra-high strength stainless steel to form a supersaturated solid solution; 2) Quenching treatment: After holding the supersaturated solid solution at a certain temperature, quench it. 3) First cold treatment: The quenched ultra-high strength stainless steel is kept at -70℃ to -200℃ for 8 to 12 hours. 4) First aging treatment: After the first cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 500℃~600℃ for 3~6 hours; 5) Second cold treatment: After the first aging treatment, cool to room temperature in air, then keep warm at -73℃ to -78℃ for 1 to 3 hours. 6) Second aging treatment: After the second cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 450℃~600℃ for 3~9 hours.

[0005] Preferably, the heating temperature in step 1) is 950℃~1100℃, and the heating time is 0.8~1.2h.

[0006] Preferably, the quenching medium in step 2) is quenching oil.

[0007] Preferably, the temperature of the first cold treatment in step 3) is -73°C and the time is 8 hours.

[0008] Preferably, the temperature of the first aging treatment in step 4) is 505°C and the time is 4 hours.

[0009] Preferably, the temperature of the second cold treatment in step 5) is -73°C and the time is 2 hours.

[0010] Preferably, the temperature for the second aging treatment in step 6) is 490°C and the time is 8 hours.

[0011] This invention is based on a composite heat treatment method with multi-level temperature and time gradients, consisting of "solution-quenching-deep cryogenic treatment-aging-re-deep cryogenic treatment-re-aging". By applying continuous hot-cold cycles in multiple temperature zones and at different holding times, the entire process of defect configuration, elemental distribution, and phase transformation behavior in ultra-high strength stainless steel can be controlled, thereby obtaining a microstructure with optimal strength, plasticity, and toughness. This invention can further implement multi-step cryogenic and aging treatments based on the properties of the alloy system and the second-phase reinforcing particles.

[0012] This invention heats the sample to above the complete austenitizing temperature, causing primary coarse precipitates (such as...) M6CThe intermetallic compounds (such as Laves phase and χ phase) are fully dissolved to form a homogeneous supersaturated solid solution. By controlling the temperature and time of the solution treatment, it is necessary to ensure the formation of a supersaturated solid solution while preventing excessive grain coarsening caused by excessively high temperature or time. After oil quenching, deep cryogenic treatment is performed using liquid nitrogen (-196℃) or an industrial cryogenic chamber (e.g., WD9-0.4FB) at -73℃ to promote the transformation of residual austenite to martensite in the high alloy system as much as possible. At the same time, high-density dislocations, vacancies and their clusters are formed in the matrix, providing a large number of high-energy nucleation sites for subsequent precipitation.

[0013] This invention activates the moderate diffusion rate of alloying elements through short-time aging in the medium temperature range (approximately 505–530°C), inducing... M2C Rapid precipitation of multiple phases, including Laves, α′-Cr, and reverse-transformed austenite, achieves the first stage of toughening.

[0014] To suppress the coarsening of reverse austenite and further activate the defect structure, this invention introduces a short-term recooling treatment at -73℃ after the first aging process. This causes the reverse austenite and the retained austenite to re-martensitize, while simultaneously promoting the further enrichment of elements such as Cr and Mo around dislocations and vacancies, providing more effective nucleation sites for precipitation in the next stage.

[0015] This invention achieves Cr-rich fine particles through long-term secondary aging in a lower temperature range (450–490℃). M2C The synergistic precipitation of Mo-rich Laves phase and stable thin-film reverse-transformed austenite. The reinforcing phase provides strong resistance to dislocation movement, while the thin-film austenite enhances the material's plasticity and toughness through energy absorption, crack deflection, and the TRIP effect.

[0016] This invention achieves systematic regulation of "defect configuration → precipitation kinetics → ratio of strengthening phase to toughening phase" through temperature gradient and time gradient design, breaking through the traditional trade-off between strength and plasticity, and strength and toughness in heat treatment, and significantly improving the comprehensive performance of ultra-high strength stainless steel. Attached Figure Description

[0017] Figure 1 The heat treatment process diagram for the experimental Fe-C-Cr-Ni-Co-Mo-WV martensitic ultra-high strength stainless steel is shown. Figure 2 Stress-strain curves of experimental steel under different heat treatment states; Figure 3 For test steels in different heat treatment states, room temperature K ⅠCFracture morphology images, including (a) macroscopic fracture morphology of the specimen in the single cold-treated state; (b) pre-crack and post-fracture crack interface of the specimen in the single cold-treated state; (c) shear lip of the specimen in the single cold-treated state; (d) macroscopic fracture morphology of the specimen in the double-aged state; (e) pre-crack and post-fracture crack interface of the specimen in the double-aged state; and (f) shear lip of the specimen in the double-aged state. Figure 4 Optical microscopy (OM) images show: (a) quenched PAGBs; (b) matrix structure of a single-aged sample; (c) matrix structure of a double-aged sample; TEM microscopy: (d) matrix structure of the quenched sample and corresponding diffraction spot calibration results; (e) matrix structure of a single-aged sample; (f) matrix structure of a double-aged sample (electron beam incident direction is [-110)). α ); Figure 5 IPF orientation diagrams of experimental steels in different heat treatment states are shown, where (ab) is the single-aged state, (cd) is the double-aged state, (e) is the IPF color diagram, and (f) is the proportion of large-angle grain boundaries (>15°). Figure 6 The figures show the characterization results of austenite microstructure in steel, including (a) XRD diffraction peaks, (b) austenite volume fraction, (c) EBSD phase diagrams of single-aged and (d) double-aged samples. Figure 7 (a) TEM bright-field image of a low-magnification single-aged sample; (b) high-angle annular dark-field image of the microstructure in the region shown by the green circle in Figure (a); (ce) EDS surface scan: Fe / Cr, Fe / Mo, and C elements; (f) electron beam from

[001] α High-resolution atomic image of a single-time-aged sample obtained by directional incident light; (g) Diffraction spots and calibration results obtained by FFT transformation of the region shown in red box in Figure (f); Figure 8 (a) TEM bright-field image of the low-magnification dual-aged sample, (b) high-angle annular dark-field image of the microstructure in the region shown by the green circle in Figure (a), (ce) EDS surface scan: Fe / Cr, Fe / Mo and C elements, (f) electron beam and

[001] α High-resolution atomic images of the dual-time-aged sample obtained by directional incident light, (g) diffraction spots and calibration results obtained by FFT transformation of the region shown in red box in Figure (f); Figure 9High-magnification TEM dark-field images: (a) single-aged state, (b) double-aged state sample; histograms of size distribution of (c) needle-like and (d) blocky / spherical precipitates in the single-aged state sample; histograms of size distribution of (c) blocky / spherical and (d) needle-like precipitates in the double-aged state sample. Figure 10 Three-dimensional atomic images of a single-time-aged sample: (a) Cyan cylindrical region (within the yellow dashed box, region size: 40 × 20 × 10 nm) 3 (a) Detailed features of the Cr-rich phase; (b) One-dimensional concentration distribution of alloying elements in the cylindrical region along the direction of the yellow arrow in (a) (group spacing 0.3 nm); Figure 11 Three-dimensional atomic images of the dual-time-aged sample: (a) Cyan cylindrical region (within the yellow dashed box, region size: 40 × 20 × 10 nm) 3 (a) Detailed features of the Cr-rich phase; (b) One-dimensional concentration distribution of alloying elements in the cylindrical region along the direction of the yellow arrow in (a) (group spacing 0.3 nm); Figure 12 Three-dimensional atomic images of single-time-aged samples: (a) Mo-rich phase and carbides obtained from the corresponding isoconcentration surfaces; one-dimensional concentration distribution maps passing through regions (b) ROI-1 and (c) ROI-3 along the direction of the yellow arrows; Figure 13 Three-dimensional atomic images of single-time-aged samples: (a) Cr-rich phase obtained by selecting the corresponding isoconcentration surface; (b)-(c) One-dimensional concentration distribution map through the ROI-6 region along the direction of the yellow arrow; Figure 14 Three-dimensional atomic images of the dual-time-aged sample: (a) Mo-rich phase and carbides obtained from the corresponding isoconcentration surfaces; one-dimensional concentration distribution maps passing through regions (b) ROI-7 and (c) ROI-9 along the direction of the yellow arrow; Figure 15 Three-dimensional atomic images of the dual-time-aged sample: (a) Cr-rich phase obtained by selecting the corresponding isoconcentration surface; (b)-(c) One-dimensional concentration distribution map through the ROI-12 region along the direction of the yellow arrow; Figure 16 The theoretical yield strength and measured yield strength values ​​of single and double-aged specimens calculated by various strengthening models are given. Detailed Implementation

[0018] This invention provides a heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel, comprising the following steps: 1) High-temperature solution treatment: Heating ultra-high strength stainless steel to form a supersaturated solid solution; 2) Quenching treatment: After holding the supersaturated solid solution at a certain temperature, quench it. 3) First cold treatment: The quenched ultra-high strength stainless steel is kept at -70℃ to -200℃ for 8 to 12 hours. 4) First aging treatment: After the first cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 500℃~600℃ for 3~6 hours; 5) Second cold treatment: After the first aging treatment, cool to room temperature in air, then keep warm at -73℃ to -78℃ for 1 to 3 hours. 6) Second aging treatment: After the second cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 450℃~600℃ for 3~9 hours.

[0019] The present invention does not specifically limit the type of ultra-high strength stainless steel, but preferably includes martensitic precipitation hardening stainless steel, semi-austenitic precipitation hardening stainless steel and martensitic aging stainless steel, and more preferably Fe-C-Cr-Ni-Co-Mo-WV martensitic ultra-high strength stainless steel.

[0020] The present invention does not have any special limitation on the source of the ultra-high strength stainless steel, and it can be obtained by commercially available products or conventional preparation methods in the field.

[0021] The present invention does not have any particular limitation on the heating method for the high-temperature solution treatment, but preferably adopts a box-type resistance furnace, a tube furnace with an inert protective atmosphere, or a vacuum muffle furnace.

[0022] In this invention, the heating temperature for the high-temperature solution treatment is preferably 950℃~1100℃, and the heating time is preferably 0.8~1.2h, more preferably 1h.

[0023] In this invention, the high-temperature solution treatment is preferably performed by heating to above the complete austenitizing temperature, and ensuring that any large, harmful primary precipitates (e.g., M6C, M7C3, M...) remaining in the as-cast state are eliminated. 23 Carbides such as C6 and intermetallic compounds are completely dissolved into the matrix to form a supersaturated solid solution, while ensuring that the grains do not grow abnormally.

[0024] In this invention, it is preferred to quench after holding the temperature for 1 hour, and it is preferred to use quenching oil as the cooling medium. This can ensure that sufficient phase transformation driving force is provided, and that the volume expansion during the martensitic phase transformation process is not too violent due to excessive cooling rate, which would cause excessive internal stress in the billet and thus cracking.

[0025] In this invention, the ultra-high strength stainless steel belongs to a high alloy system with an increased martensitic transformation point. After quenching, a large amount of metastable austenite remains at room temperature. Preferably, the material needs to be cold-treated at -196℃ for 8 hours after quenching to provide additional phase transformation driving force and promote the transformation of metastable austenite to martensite in order to obtain as much martensite as possible.

[0026] In this invention, after the first cold treatment, the temperature is restored to room temperature, and the first aging treatment is performed. Preferably, the temperature is kept in the medium temperature range of 505-530℃ for 4-5 hours to induce the precipitation of M2C carbides (mainly), Laves phase, α'-Cr phase and reverse-transformed austenite.

[0027] In this invention, after the first aging treatment, the substrate is air-cooled to room temperature, and then subjected to a second short-term cold treatment, preferably at -73°C, for 2 hours, to eliminate the metastable residual austenite or reverse-transformed austenite that grew during the first aging process, increase the matrix dislocation density, and retain high-density vacancies (point defects). Furthermore, Mo and Cr atoms are further enriched at crystal defects due to the phase transformation process, providing more high-energy heterogeneous nucleation sites for the precipitation of precipitated phases during the subsequent second aging process.

[0028] In this invention, after the second cold treatment and after returning to room temperature, a second aging treatment is carried out at a low temperature range, preferably 450-490°C, to further induce the precipitation of Cr-rich M2C, Mo-rich Laves phase intermetallic compound phases, especially α'-Cr phase and reverse-transformed austenite, thereby achieving the optimal strength-ductility and toughness matching of ultra-high strength stainless steel.

[0029] This invention can also be applied to the alloy system and type of reinforcing phase of multiphase precipitation-strengthened ultra-high strength stainless steel by adding a multi-step cold treatment-aging process with gradient holding temperature and time to synergistically precipitate high-density, dispersed second-phase reinforcing particles and austenitic phase toughening units.

[0030] The present invention does not have any particular limitation on the methods of cold treatment and aging treatment, and conventional cold treatment and aging treatment methods in the art can be used.

[0031] The technical solutions provided by the present invention will be described in detail below with reference to the embodiments, but they should not be construed as limiting the scope of protection of the present invention.

[0032] Example 1 (1) The test steel was Fe-C-Cr-Ni-Co-Mo-WV series martensitic ultra-high strength stainless steel, and the following was adopted. Figure 1 The heat treatment method, through the design of temperature gradient and time gradient, realizes the systematic regulation of "defect configuration → precipitation kinetics → strengthening phase – toughening phase ratio", thereby significantly improving the comprehensive performance of ultra-high strength stainless steel.

[0033] (2) Room temperature mechanical properties of high Cr-Ni low carbon ultra-high strength stainless steel Table 1. Room temperature mechanical properties of test steels in various heat-treated states

[0034] As shown in Table 1, the Rockwell hardness (HRC) of the tested steel increased significantly after aging treatment, exhibiting a strong age-hardening effect. Specifically, compared with the solution-quenched (ST) specimen, the hardness values ​​of the single- and double-aged specimens increased by 18% and 29%, respectively. This result is consistent with... Figure 2 The stress-strain curves of the solution-quenched and single / double-aged specimens show consistent characteristics. As can be seen from the figure, aging treatment significantly improves the strength of the tested steel, while simultaneously reducing the corresponding plasticity index (elongation after fracture). Notably, compared to the single-aged specimen, the double-aged specimen exhibits improved strength and plasticity simultaneously. Specifically, the room temperature yield strength (YS) of the double-aged specimen is approximately 359 MPa higher than that of the single-aged specimen, and the tensile strength (UTS) (see Table 1) also increases from 2115±3.5 MPa to 2259±1.0 MPa. Meanwhile, compared to the plasticity indexes of the single-aged specimen, namely elongation after fracture (EL) and reduction of area (RA), the plasticity of the tested steel is significantly improved after double aging treatment. In particular, the reduction of area of ​​the material increases from 35±2.0% to 52±0.1%, an increase of approximately 48.5%. Furthermore, the toughness indices of the experimental steel (including impact toughness and planar fracture toughness) showed a monotonic decreasing trend with the advancement of heat treatment processes and the improvement of strength. It is worth mentioning that there was no significant difference in the planar fracture toughness values ​​between single- and double-aged specimens, meaning that while the strength was significantly improved, the material retained considerable crack resistance.

[0035] Given the similar planar fracture toughness values ​​of single- and double-aged specimens, the following discussion focuses solely on the fracture morphology characteristics of the planar fracture toughness specimens in both the single-cold-treated and double-aged states. The fracture morphologies of the planar fracture toughness specimens in the two heat-treated states are shown below. Figure 3 As shown in the figure, similar to the previously mentioned test steel, both exhibit macroscopic fracture surfaces consisting of a fatigue pre-crack zone, a monotonic fracture zone, and shear lips located at the edges of the specimen. Figure 3 As can be seen from (a) and (d), the fracture surfaces of the NDS steel samples in both heat-treated states are very smooth. This phenomenon indicates that the test steel fractured without undergoing significant deformation under plane strain conditions. Figure 3 (b) and (e) show the K samples in the single-cold-treatment state and the double-aging state, respectively. ⅠCMicroscopic morphology of the fracture surface. As shown in the figure, the stretched zone exists between the fatigue pre-crack fracture zone and the crack monotonic fracture zone. The morphology of the fatigue pre-crack fracture zone is typical cleavage fracture, while the morphology of the crack monotonic fracture zone is mainly dimples, containing secondary cracks and a large number of tear ridges. It is worth noting that there are certain differences in the morphology of the crack monotonic fracture zone of the specimens in the two heat treatment states. Specifically, the crack monotonic fracture zone of the double-aged specimen exhibits a quasi-cleavage fracture morphology. The average width of the crack stretched zone of the cold-treated and double-aged specimens was measured to be 26.65±3.68 μm and 21.29±4.82 μm, respectively, indicating that the crack propagation path from the pre-crack region to the crack monotonic fracture zone is longer in the cold-treated specimen, requiring more energy. Figure 3 (c) and (f) show the K state under single-cooling and double-aging conditions, respectively. ⅠC The shear lip morphology of the fracture surface of the specimens consisted of micropores, dimples, and short tear ridges. Statistical results on the average shear lip width showed that, compared to the average shear lip width of the single-cold-treated specimens (286.06±36.03 μm), the average shear lip width of the double-aged specimens was significantly reduced (89.16±8.76 μm). This result is consistent with the aforementioned Ferrium S53 steel K... ⅠC The fracture characteristics of the specimens were consistent, that is, the planar fracture toughness value of the test steel was positively correlated with the average width of the shear lip.

[0036] Figure 4 The image shows the microstructure of the steel matrix under different heat treatment states. Figure 4 (a) shows the original austenitic grains (PAGBs) of the test steel. The average size of the PAGBs, measured by the conventional intercept method, is 108 ± 1.6 μm. It can be seen that the grains of the test steel became relatively coarse after solution treatment at 1100 ℃. The microstructure of the single and double-aged samples exhibits typical lath martensite characteristics, i.e., the PAGBs are composed of several packets with different orientations (see...). Figure 4 (b)-(c)). Figure 4 Figures (d)-(f) show the TEM microstructure of the steel in the solution-treated, single- and double-aged states, respectively. As can be seen from the figures, a large number of twins with the same angle and orientation as the laths exist within the martensitic laths. The crystallographic orientation of the twins within different laths is not entirely the same (see Figure 1). Figure 4 (c) This may be due to the formation of twins at different crystallographic directions during the phase transformation. Statistical results show that the average widths of the martensitic laths in the solid solution state, single- and double-aged states of the experimental steel are 0.271 ± 0.069 μm, 0.298 ± 0.065 μm, and 0.312 ± 0.125 μm, respectively.

[0037] To further illustrate the evolution of the austenitic toughening phase in NDS test steels under different heat treatment states, multiple methods were used to characterize it, and the results are as follows: Figure 6 As shown. Specifically, Figure 6 Figures (a) and (b) show the characteristic diffraction peaks of the XRD matrix of the test steel and the corresponding calculated austenite volume fraction, respectively. As can be seen from the figures, compared to other heat-treated samples, the quenched (ST) and single-cold-treated (FCT) samples exhibit (111)... γ Austenite characteristic diffraction peaks (see) Figure 6 (a) A magnified view of the area shown in the red dashed box) shows higher strength; its corresponding austenite volume fraction is also significantly higher than that of the specimens under other heat treatment regimes. It is worth noting that, compared to the single-aged specimen, the double-aged specimen has (111) γ The intensity of the characteristic diffraction peaks of austenite was slightly increased. Figure 6 (a) The area circled in green (in dashed circles). Correspondingly, the austenite content in the double-aged sample is approximately four times higher than that in the single-aged steel. Figure 6 As shown in the EBSD phase diagrams (c)-(d), the number of fine spherical and strip-shaped reverse-transformed austenite (red area) in the double-aged sample is significantly increased compared with that in the single-aged sample, and they are dispersed on the martensitic matrix (blue area); this observation is consistent with the austenite volume fraction measured by XRD.

[0038] Figure 7 The image shows the TEM microstructure characterization results of the steel from a single-aging test. Figure 7 (a) (Bright field image) High-angle annular dark field image of a local area of ​​the matrix shown by the green circle ( Figure 7 (b) It can be seen that two nanoscale precipitates with significantly different contrast exist in the single-aged sample: a bright white needle-like precipitate and a black spherical precipitate. The corresponding EDS surface scan results show (see...) Figure 7 (c)-(d)) The single-aged samples showed obvious Cr-rich spherical clusters and a certain amount of Cr- and Mo-rich irregular needle-like and blocky precipitates, with the degree of Mo clustering being less than that of Cr. EDS surface scan results for C (see...) Figure 7 (e) did not show obvious local enrichment of C element. Regarding Figure 5 (f) The FFT transformation results of the area shown in the red box indicate that the diffraction spot calibration results show that there are diffraction spots in the single-aged steel. M 2C carbide precipitate phase.

[0039] Figure 8 The image shows the TEM microstructure characterization results of the steel tested in the double-aged state. Figure 8 (a) (Bright field image) High-angle annular dark field image of a local area of ​​the matrix shown by the green circle ( Figure 8(b) It can be seen that two nanoscale precipitates with significantly different contrasts also exist in the dual-aged sample: a bright white needle-like precipitate and a black blocky precipitate. The corresponding EDS surface scan results show (see...) Figure 8 (c)-(d)) The single-aged samples exhibited obvious Cr-rich spherical and blocky clusters, as well as a certain amount of Cr- and Mo-rich irregular needle-like and blocky precipitates. Furthermore, the degree of Mo clustering was significantly higher in these samples than in the single-aged samples. EDS surface scan results for C (see...) Figure 8 (e) No obvious local enrichment of C element was observed. Regarding... Figure 8 (f) The FFT diffraction spot calibration results of the area shown in the red box indicate that the Laves phase exists in the double-aged steel.

[0040] Figure 9 The figures show the statistical results of the size distribution of dark-field (DF) TEM images and internal precipitates in single- and double-aged steels, respectively. Figure 9 As can be seen in (a)-(b), a large number of second-phase particles are dispersed on the steel matrix in the single-aging test. Among them, there are two precipitates with different morphological characteristics and significantly different contrasts: black spherical precipitates and bright white needle-like precipitates. On the other hand, two second-phase particles with significantly different contrasts are also present in the double-aging sample: black blocky and spherical precipitates and bright white needle-like precipitates. Based on the above characterization results, it can be reasonably inferred that the larger black blocky precipitates are formed by the continuous growth and polymerization of the spherical precipitates during the double-aging process. The size (equivalent circle diameter) of the precipitates with different heat treatment states and morphologies was statistically analyzed, and the results are as follows: Figure 9 As shown in (c)-(f), the average size of the needle-like precipitates gradually coarsens with further aging, with the equivalent circle diameter increasing from 3.9 ± 0.9 nm to 5.2 ± 1.1 nm. In contrast, the blocky / spherical precipitates did not show significant coarsening with the aging process; their average size only slightly increased from 5.5 ± 1.3 nm to 5.8 ± 1.7 nm, which may contribute to improving the tempering softening resistance of the experimental steel.

[0041] The morphology, distribution, and micro-region composition characteristics of various nanoscale precipitates in single- and double-aged samples were analyzed using samples of corresponding isoconcentration. Figure 10 and 11 The figures show the atomic diagrams of single- and double-aged samples, respectively, reconstructed using corresponding isoconcentration surfaces. As can be seen from the figures ( Figure 10(a) The single-aged needle-shaped sample contained a large number of irregularly banded Mo-rich phases and carbides; simultaneously, numerous fine Cr-rich phases with sizes in the range of several nanometers were dispersed throughout the characterization region. Similar to the single-aged sample ( Figure 11 (a) In the dual-aged samples, a large number of carbide precipitates, appearing as bands and large blocks, were also observed; simultaneously, Cr enrichment was observed in the carbides. On the other hand, a large number of isolated, fine Cr-rich phases were also dispersed in the matrix. Examples of such precipitates were selected... Figure 10 In (a) and 11(a), the cyan cylindrical region indicated by the yellow dashed box can be reconstructed using a 30 at.% Cr isoconcentration surface to obtain the atomic image and corresponding cross-sectional diagram of the Cr-rich phase in this region. As shown in the figure, the Cr-rich phase morphology in the single-aged sample is mainly composed of interconnected bands and ellipsoids, while the Cr-rich phase in the double-aged sample is mostly ellipsoidal and tends to precipitate independently. The interfacial area between different Cr-rich phases is significantly smaller than that between Cr-rich phases in the single-aged state. This is attributed to the further growth of large-sized Cr-rich phases and the intensified matrix amplitude modulation decomposition during the double-aging process. To qualitatively describe the matrix amplitude modulation decomposition in single- and double-aged NDS steels, the one-dimensional linear concentration distribution of each major alloying element along the direction of the yellow arrow in the cyan cylindrical region was characterized (group spacing: 0.3 nm). The results are as follows: Figure 10 As shown in (b) and 11(b), the one-dimensional concentration curve of Cr in the aged steel exhibits significant fluctuations and distinct peaks and troughs compared to other elements. This phenomenon indicates that significant amplitude modulation decomposition occurred in the matrix of both single- and double-aged steel samples. Furthermore, since the amplitude of the one-dimensional concentration curve of Cr in the double-aged sample is significantly greater than that in the single-aged sample, the degree of amplitude modulation decomposition in the matrix of the double-aged sample should be higher than that in the single-aged sample.

[0042] During the aging process, the test steel should exhibit the precipitation of two types of second-phase particles within the matrix: a Laves phase, an intermetallic compound rich in Mo and Cr, and a metastable phase rich in Cr, Mo, and C. M 2C carbides. To further clarify the types and micro-region compositional characteristics of nanoscale reinforcing phases in single- and double-aged experimental steels, three-dimensional reconstructions were performed on the above-mentioned heat-treated needle-shaped samples with corresponding isoconcentration surfaces. Figure 14 The image shown is a three-dimensional atomic image of the Mo-rich phase and carbides in a single-aged sample. Figure 12(a) It can be seen that the Mo-rich phase in the single-aged sample exhibits a relatively large, blocky morphology, while the carbide morphology is a smaller, blocky and lamellar morphology. It is noteworthy that some Mo-rich phase and carbide micro-regions overlap. This is because Mo itself is a strong carbide-forming element, therefore, there should be competition between the Mo-rich phase and carbides for Mo in the steel; thus, both types of precipitation co-precipitate within the characteristic micro-regions. A typical Mo-rich phase (ROI-1) and a carbide (ROI-3) are selected from each. Figure 12 (a) The red and red dashed ellipse areas show the approximate histograms of the main constituent elements along the directions indicated by the yellow arrows (see the cross-sectional views of the two types of precipitates shown in the red dashed box in 12(a)). As can be seen from the figures, Cr and Mo are enriched in the Mo-rich phase. Notably, W, V, and Si also show some segregation in the Mo-rich phase. On the other hand, Cr and Mo are mainly enriched in the carbides, as well as W and V, two strong carbide-forming elements, showing segregation. (Selection) Figure 13 (a) The micro-region shown in the yellow dashed box was reconstructed three-dimensionally using a 20 at.% Cr isoconcentration surface. It can be seen that two types of Cr-rich phases of different sizes and morphologies exist in the single-aged sample: large-sized, blocky Cr-rich phases and fine, diffusely distributed ellipsoidal Cr-rich phases. Given the size and morphology of the former, it should be a Mo-rich phase or carbides, while the latter is an α'-Cr phase produced by matrix amplitude modulation decomposition. Therefore, a typical small-sized ellipsoidal α'-Cr phase (ROI-6) (red ellipse in 13(a)) is selected, and a one-dimensional approximate concentration histogram of it along the direction indicated by the yellow arrow in the cross-section (red dashed box in 15(a)) is given. Figure 13 (b)-(c)). As can be seen from the figure, the main constituent elements in the α'-Cr phase are Cr and Fe.

[0043] Figure 14 The image shows a three-dimensional atomic image of the Mo-rich phase and carbides in the dual-aged sample. As can be seen from the image ( Figure 14 (a) The morphology of the Mo-rich phase in the dual-aged state is more complex, mostly appearing as irregular strips and ellipsoids, while the carbides appear as shorter bands or lamellar structures, with the latter being significantly smaller than the former. Simultaneously, a co-precipitation phenomenon of Mo-rich phase and carbides similar to that in the single-aged sample was also observed. Typical microregions of Mo-rich phase and carbides (ROI-7 and ROI-9) were selected from each of the dual-aged samples. Figure 14 (a)), at Figure 14 (b) and (c) respectively give the approximate one-dimensional concentration histograms of each alloying element along the direction of the yellow arrow in the Mo-rich phase and carbide cross-section diagram within the red dashed box shown in 14(a). As can be seen from the figures ( Figure 14(b) In the Mo-rich phase of the dual-aged sample, the elements enriched are mainly Cr and Mo, with some W and Si also present. On the other hand, the carbide second phase is mainly enriched with strong carbide-forming elements such as Cr, Mo, W, and V. Figure 14 (c)). Furthermore, select as follows: Figure 15 (a) The region indicated by the yellow dashed box was reconstructed using a 20 at.% Cr isoconcentration surface to obtain a 3D atomic map of the Cr-rich phase in this region. As shown in the figure, the Cr-rich phase in this region is mostly in the range of several nanometers in size, and predominantly exhibits an independently precipitated ellipsoidal morphology. To further illustrate its elemental composition, a typical Cr-rich phase (ROI-12) was selected. Figure 15 (a) at the red dashed ellipse), along the cross-section ( Figure 15 (a) Create a one-dimensional approximate concentration histogram in the direction of the yellow arrow within the red dashed box. Figure 15 (b)-(c)). As can be seen from the figure, the constituent elements of the Cr-rich single-aged Cr-rich phase are consistent, mainly consisting of Cr and Fe atoms.

[0044] A strengthening model was used to quantitatively estimate the contribution of various strengthening mechanisms—solution strengthening, lath martensitic matrix strengthening (including grain refinement and dislocation strengthening), and second-phase precipitation hardening—to the room-temperature yield strength increment of single- and double-aged NDS steel. Among these, the lattice friction stress of the α-Fe matrix was also considered. σ 0 Still using 50 MPa, the parameters required for the Taylor formula to calculate the strength contribution increment of the martensitic matrix can be found in Table 2. Given that during the aging process, supersaturated solute atoms originally dissolved in the matrix leave the interstitial positions of the matrix lattice during the precipitation of second-phase particles and the formation of reverse austenite, the Fleishcher formula is used to quantitatively calculate the strength increment of the solid solution strengthening mechanism. As can be seen from the table, the martensitic matrix layer structure of the steel samples after double aging treatment is coarsened to a certain extent, and its dislocation density increases slightly, which should be due to the increased matrix dislocation density caused by the secondary cold treatment.

[0045] Table 2. Average size and dislocation density of martensitic hierarchical structure in steel under aging conditions.

[0046] because M Both the 2C and Laves phases have HCP structures and similar elemental compositions (both are Cr, Mo, and W-rich phases). TEM characterization results failed to confirm the presence of these phases in the aged steel. M The co-precipitation of 2C and Laves phases was confirmed by the APT characterization results of the corresponding heat-treated specimens. Multiple microscopic characterization results revealed the presence of three types of nanoscale second-phase reinforcing particles in the aged steel, namely…M The precipitates included 2C carbides, Mo-rich intermetallic compound Laves phase, and Cr-rich α'-Cr phase generated by matrix amplitude modulation decomposition. Measured physical parameters of various precipitated strengthening phases in the aged specimens are shown in Table 3. As can be seen from the table, after double aging treatment, the volume fractions of the three precipitates significantly increased, while the average equivalent circle radius slightly increased. This indicates that during the secondary aging process, the precipitates did not undergo significant coarsening, maintaining a fine spatial size, and the content of the second phase particles further increased, thus providing a greater increment in precipitation strengthening strength. It is noteworthy that the volume fraction of the Laves phase in the experimental steel is close to the volume fraction of the Laves phase predicted by thermodynamic calculations, while... M The measured volume fraction of APT in 2C differs significantly from the thermodynamically calculated value.

[0047] To further illustrate the influence of precipitation behavior on the room temperature yield strength of aged steel, the following quantitative calculations will be performed on the strength increments provided by various types of precipitates. It should be noted that... M The 2C strengthening mechanism was still quantitatively calculated using the Orowan bypass mechanism. Studies on a certain type of composite precipitation-strengthened martensitic ultra-high strength stainless steel showed that the critical transformation sizes for the Laves phase and α'-Cr phase precipitation strengthening mechanisms were 7.5 nm and 6.0 nm, respectively. Given that the average size (equivalent circle radius) of both types of precipitates in the aged specimens was much lower than the critical size, the sizes of the Laves phase and α'-Cr phase were calculated separately. Cr The coherence, modulus, and ordered strengthening contributions of the phases are noteworthy. It is worth noting that the Laves phase in this experimental steel is relatively small in size and has a relatively regular morphology, and its composition is similar to that of the Laves phase. M Since 2C has many identical constituent elements, the lever rule does not apply to the calculation of the volume fraction of the Laves phase in this experimental steel. Therefore, the volume fraction of the corresponding precipitate was obtained by directly using the number of atoms in the precipitate obtained from APT characterization and the volume percentage of the entire characterization region.

[0048] Table 3. Average number density, volume fraction, and equivalent circle radius of precipitates in aged steel.

[0049] Table 4 shows the quantitative calculations of room temperature yield strength increments for each strengthening model. As can be seen from the table, the strength increment in NDS steel mainly originates from precipitation strengthening. Among the three precipitates, the Laves phase exhibits the most significant strengthening effect, followed by [other precipitates]. M 2C carbides, α' Cr The Laves phase contributes the least to the strength increase. After double aging treatment, the newly generated precipitation-strength increase in the experimental steel mainly comes from the Laves phase and α' phase. Cr Mutually. Figure 16The figure shows a comparison between the theoretically calculated room temperature yield strength and the measured value. As can be seen from the figure, the differences between the calculated and measured yield strength values ​​of the single- and double-aged steel are 23 MPa and 24 MPa, respectively, indicating that the error is small. The above model is suitable for predicting the strength increment of this aged steel.

[0050] Table 4. Intensity increments (MPa) calculated by each strengthening model

[0051] The above description is only a preferred embodiment of the present invention. It should be noted that for those skilled in the art, several improvements and modifications can be made without departing from the principle of the present invention, and these improvements and modifications should also be considered within the scope of protection of the present invention.

Claims

1. A heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel, characterized in that, Includes the following steps: 1) High-temperature solution treatment: Heating ultra-high strength stainless steel to form a supersaturated solid solution; 2) Quenching treatment: After holding the supersaturated solid solution at a certain temperature, quench it. 3) First cold treatment: The quenched ultra-high strength stainless steel is kept at -70℃ to -200℃ for 8 to 12 hours. 4) First aging treatment: After the first cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 500℃~600℃ for 3~6 hours; 5) Second cold treatment: After the first aging treatment, cool to room temperature in air, then keep warm at -73℃ to -78℃ for 1 to 3 hours. 6) Second aging treatment: After the second cold treatment, allow the temperature to return to room temperature in the air, and then keep it at 450℃~600℃ for 3~9 hours.

2. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, The heating temperature in step 1) is 950℃~1100℃, and the heating time is 0.8~1.2h.

3. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, In step 2), the quenching medium is quenching oil.

4. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, In step 3), the temperature of the first cold treatment is -73℃ and the time is 8 hours.

5. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, In step 4), the temperature for the first aging treatment is 505℃ and the time is 4 hours.

6. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, In step 5), the second cold treatment is performed at a temperature of -73°C for 2 hours.

7. The heat treatment method for synergistically improving the strength, plasticity, and toughness of ultra-high strength stainless steel according to claim 1, characterized in that, In step 6), the second aging treatment is carried out at a temperature of 490°C for 8 hours.

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