Ultrahigh-strength double-refined high-temperature bearing steel and preparation method thereof

Abstract

The invention relates to ultrahigh-strength double-refined high-temperature bearing steel and a preparation method thereof. The invention relates to ultrahigh-strength double-refined high-temperature bearing steel, which comprises the following chemical components in percentage by mass: 0.75 to 0.85 percent of C, 3.75 to 4.30 percent of Cr, 4.00 to 4.50 percent of Mo, 0.90 to 1.10 percent of V, less than or equal to 0.35 percent of Si, less than or equal to 0.35 percent of Mn, 0.05 to 0.15 percent of Nb, 0.5 to 1.0 percent of Al, less than or equal to 9ppm of O, less than or equal to 0.002 percent of Ti, 0.03 to 0.05 percent of Ce + La and the balance of iron and other inevitable impurities. The method has the beneficial effects that by adding Nb and Al for composite alloying, the problems of large-size and non-uniformly distributed eutectic carbides and coarse original austenite grains in the M50 high-temperature bearing steel are solved, and a new breakthrough in tensile strength and rotational bending fatigue strength of the M50 steel is expected to be realized through structure refinement; and a new alloying scheme is provided for developing the high-temperature bearing steel with long service life and high reliability.

Description

Technical Field

[0001] This invention relates to the field of high-temperature bearing steel materials, and in particular to an ultra-high strength double-refined high-temperature bearing steel and its preparation method. Background Technology

[0002] M50 steel is a molybdenum-chromium-vanadium high-carbon alloy steel, belonging to the category of fully hardening martensitic high-speed steel. Due to the significant secondary hardening effect caused by the dispersed precipitation of nanoscale alloy carbides during tempering, it maintains high hardness at both room temperature and high temperature. Because of its excellent hot strength, wear resistance, dimensional stability, and high-temperature contact fatigue resistance, M50 steel is considered an ideal material for manufacturing aero-engine bearings and spacecraft transmission system components. Aero-engine bearings are subjected to the combined effects of high-speed rotation, extremely high contact stress, and high temperatures for extended periods. With the development of aero-engines towards higher thrust-to-weight ratios and longer service lives, the existing properties of M50 steel are insufficient to meet the requirements for superior durability of bearing materials under extreme operating conditions. In-depth analysis shows that its performance bottleneck is related to two inherent microstructural defects. First, due to its high-carbon and high-alloy composition, M50 steel experiences segregation of alloying elements between dendrites during the final solidification stage, leading to the precipitation of large-sized eutectic carbides. Although these brittle phases undergo partial fragmentation and refinement during subsequent large deformation forging and high-temperature diffusion, their excellent high-temperature thermal stability means that large carbide particles of 30–50 μm remain after final quenching and tempering. Under cyclic stress, stress concentration induced at the carbide-matrix interface becomes the initiation source of fatigue cracks. Second, at conventional quenching temperatures, the original austenite grain size in M50 steel tends to grow excessively, with a grain size level of only ASTM grade 8 (approximately 15 μm). This coarse-grained structure results in insufficient strength and toughness reserves, failing to effectively suppress fatigue crack propagation and limiting the durability and reliability of bearings under ultra-high stress. In summary, to break through the performance limits of M50 steel and achieve long service life and high reliability of high-temperature bearings, it is necessary to focus on the synergistic refinement of microstructure, that is, to simultaneously achieve the fragmentation and uniform distribution of eutectic carbides, as well as the significant refinement of the original austenite grains.

[0003] In existing technological cases, there are mainly the following paths to achieve microstructure refinement in M50 steel. Firstly, the complex heat treatment path; this path refines the microstructure through complex solid-state phase transformation processes. A typical example is patent CN118531189A, which employs a complex process combining three high-temperature homogenization treatments with two deep cryogenic treatments to significantly eliminate primary carbides and improve grain size to level 9. However, the entire process involves at least seven main steps, resulting in low production efficiency, increased energy consumption, and difficulty in precise control during actual large-scale production. Secondly, the thermomechanical processing path. This path combines complex large plastic deformation with specific thermal processes to physically break down and modify carbides. For example, patent CN114457212A employs a multi-stage thermomechanical process including "two-stage pretreatment of ingots," "upsetting and drawing," and "step-cooling secondary heat treatment of forged billets" to refine the average carbide size to 1.37–1.66 μm. However, its core relies heavily on precise control of the heat treatment process, resulting in a narrow process window. Summary of the Invention

[0004] Existing technological processes (complex heat treatment and thermomechanical processing) are often accompanied by high costs, high energy consumption, and complex control challenges. This invention innovatively proposes a technical solution using Nb and Al composite alloying, aiming to circumvent complex processes and achieve more stable and efficient synergistic refinement of austenite grains and primary carbides through compositional design, providing a new approach to achieving breakthroughs in the microstructure and properties of M50 steel.

[0005] The technical solution of the ultra-high strength double-refined high-temperature bearing steel provided by this invention is as follows:

[0006] An ultra-high strength double-refined high-temperature bearing steel, wherein the chemical composition of the bearing steel by mass percentage includes C: 0.75-0.85%, Cr: 3.75-4.30%, Mo: 4.00-4.50%, V: 0.90-1.10%, Si: ≤0.35%, Mn: ≤0.35%, Nb: 0.05-0.15%, Al: 0.5-1.0%, O ≤9ppm, Ti ≤0.002%, Ce+La: 0.03-0.05%, with the balance being iron and other unavoidable impurities.

[0007] Furthermore, the bearing steel has a grain size ≥10, a maximum carbide size ≤20μm, a tensile strength at room temperature ≥2800MPa, a yield strength ≥2400MPa, and a rotational bending fatigue strength (R=-1, 10) 7 ≥1100MPa.

[0008] Furthermore, it includes the following steps:

[0009] (1) Smelting and solidification

[0010] The casting was carried out in a 50kg vacuum induction furnace. After casting and demolding, the casting was placed in a 700℃ furnace for annealing for 5 hours and then cooled in the furnace to 500℃ for air cooling to eliminate casting stress.

[0011] (2) High-temperature diffusion and forging

[0012] The billet is heated to 1200-1250℃ and held for no less than 6 hours for high-temperature homogenization treatment to fully eliminate element segregation. Then it is forged in the range of 900-1200℃ into a long bar of Φ15. After forging, it is quickly placed in a 680℃ heat treatment furnace and cooled to 500℃ before being air-cooled to room temperature.

[0013] (3) Normalizing treatment

[0014] Forged bars are normalized to eliminate network carbides by holding at 1000–1100℃ for 0.5–1h, and then air-cooled after being removed from the furnace to refine the grains and improve the uniformity of the microstructure.

[0015] (4) Spheroidizing annealing

[0016] The normalized bars are subjected to spheroidizing annealing. The specific process is as follows: the temperature is raised to 850-890℃ in the furnace and held for 6-8 hours, then the furnace is cooled to 550℃ at a rate of 20℃ / h and then air-cooled to obtain a uniform spheroid pearlite structure.

[0017] (5) Quenching and tempering

[0018] The spheroidized annealed bearing steel is subjected to quenching and tempering treatment. The quenching process is a furnace-stage heating process. The specific quenching and tempering process is as follows: first, the temperature is raised to 800-850℃ and held for 0.5-1h, then the temperature is raised to 1050-1150℃ and held for 15-40min to complete austenitization, followed by oil quenching. Finally, the steel is tempered 2-4 times at 530-550℃ for 2h each time to obtain a stable tempered martensite structure.

[0019] The functions and proportions of each element in this invention are based on the following:

[0020] Carbon (C): Carbon is the most fundamental element in high-temperature bearing steel. Its role is to ensure the formation of a high-carbon martensitic matrix after quenching, thereby achieving the required matrix hardness (≥60HRC). Simultaneously, it combines with strong carbide-forming elements to form various alloy carbides, providing wear resistance and secondary hardening. Its composition range is 0.75–0.85%. If it is below the lower limit, wear resistance and hardness cannot be guaranteed. If it is above the upper limit, it leads to an excessive number of eutectic carbides with excessively large sizes, forming a network structure that impairs toughness and fatigue performance. Furthermore, carbon content exceeding this value does not adequately improve hardness, resulting in more disadvantages than advantages.

[0021] Chromium (Cr): Chromium primarily provides hardenability, corrosion resistance, and secondary hardening. After quenching, chromium dissolves largely in austenite, significantly improving hardenability and enabling complete hardening of large bearing rings and rolling elements. Chromium also provides some resistance to tempering, oxidation, and corrosion. Its composition range is 3.75%–4.30%. If it is below the lower limit, hardenability cannot be guaranteed; if it is above the upper limit, the residual austenite content and the hardening properties of large bearing rings and rolling elements will be compromised. 23 An increase in C6 impairs toughness and dimensional stability.

[0022] Molybdenum (Mo): Molybdenum is the core element that imparts red hardness. During tempering, fine Mo2C alloy carbides are dispersed and precipitated, producing a strong secondary hardening effect, allowing M50 steel to maintain high hardness (≥58HRC) even at 316℃. In addition, molybdenum also has advantages such as refining grain size, improving matrix toughness, enhancing tempering stability, and suppressing or reducing temper brittleness. When the molybdenum content is between 4.0% and 4.5%, its secondary hardening effect, toughness, and tempering stability are optimally balanced. Too low a content results in insufficient high-temperature performance, while too high a content leads to a significant increase in cost without a significant performance gain.

[0023] Vanadium (V): Vanadium primarily provides wear resistance and refines grain size. Vanadium mainly forms high-hardness, highly stable VC or V4C3 carbides. These dispersed, fine carbides are the hardest phase in steel, effectively resisting abrasive wear. Simultaneously, during heating, these carbides pin grain boundaries, inhibiting the high-temperature growth of austenite grains. Its composition is set between 0.90% and 1.10%. Too low a content results in insufficient VC, leading to inadequate wear resistance; too high a content causes a dramatic increase in VC, resulting in poor machinability. Within this range, a balance between wear resistance and machinability is achieved.

[0024] Silicon (Si): Silicon strengthens the matrix by solid solution and improves its strength and elastic limit, and inhibits the formation of cementite during low-temperature tempering. Excessive Si will impair machinability and purity. Its content is generally ≤0.35%.

[0025] Manganese (Mn): Manganese mainly improves hardenability and assists in deoxidation. Excessive Mn will segregate at grain boundaries and stabilize residual austenite, increasing the difficulty of microstructure control; its content is generally ≤0.35%.

[0026] Rare earth elements (Ce+La): Rare earth elements play a role in modifying carbides, refining microstructure, purifying molten steel, and strengthening grain boundaries. During solidification, highly surface-active rare earth elements preferentially segregate at the carbide-liquid interface, inhibiting the directional growth of carbides and transforming the network-like, lamellar, coarse eutectic ledeburite into fine, isolated, spherical or short rod-shaped carbides. Furthermore, rare earth elements have a strong affinity for oxygen and sulfur, forming high-melting-point spherical rare earth oxysulfides. This reduces the formation of harmful inclusions in steel, such as MnS, and the spherical rare earth compounds can pin grain boundaries, refining the grains. Simultaneously, rare earth elements can reduce the segregation of phosphorus and sulfur at grain boundaries, strengthening them. As a "morphology modifier" and "purifier," adding trace amounts of rare earth elements to steel produces significant effects, with the content controlled at 0.03–0.05%.

[0027] Niobium (Nb): Niobium is one of the core elements of the steel in this invention. Niobium is often used as a "nucleation core" and "pinning agent," playing a crucial role in refining the microstructure. Niobium has the highest affinity for carbon, preferentially forming high-melting-point NbC carbides during solidification, providing heterogeneous nucleation cores for primary carbides and increasing the nucleation rate. Simultaneously, NbC precipitation reduces carbon in the molten steel, decreasing the driving force for eutectic carbide formation, thus refining the size and improving the distribution. Furthermore, during solid-state phase transformation and heat treatment, fine NbC particles strongly pin grain boundaries, significantly refining the original austenite grains. This dual refinement by carbides and austenite is key to improving the material's strength, toughness, and fatigue resistance. To achieve the best results, the Nb content is set between 0.05% and 0.15%.

[0028] Aluminum (Al): Aluminum is another core element of the steel of this invention. Aluminum regulates the precipitation and decomposition of eutectic carbides by influencing solidification and high-temperature phase transformation. Al increases the carbon content at the eutectic point. To obtain this eutectic composition, more primary austenite precipitation is required during solidification to enrich the carbon content in the residual liquid phase, thereby reducing the liquid phase fraction that ultimately undergoes the eutectic transformation. Simultaneously, Al has a higher concentration in the matrix. Compared to Fe, Al has a greater difference in electronegativity with elements such as Mo and V, enhancing the solid solution of alloying elements in the matrix and reducing alloy segregation in the molten steel at the end of solidification, thus reducing the formation of eutectic carbides. Furthermore, a small amount of Al incorporated into the M2C lattice reduces its thermal stability and decomposition barrier. Al also increases the activity and diffusion coefficient of carbon in austenite. On the one hand, it enhances the diffusion of carbon from eutectic carbides to the matrix at high temperatures, promoting eutectoid decomposition; on the other hand, it promotes the precipitation of more fine carbides in the matrix. The grain boundary pinning effect of these fine carbides significantly refines the grain size. In addition, Al reacts with free nitrogen in steel to form AlN particles, which can also provide a small amount of grain boundary pinning force. The aluminum content is set between 0.5% and 1.0%; if the content is too low, the refining ability will be insufficient, and if the content is too high, a harmful high-temperature ferrite phase will be formed.

[0029] The implementation of this invention has the following technical effects:

[0030] This invention discloses an ultra-high strength, double-refined high-temperature bearing steel and its preparation method. By adding Nb and Al composite alloying, it solves the problems of large-sized, unevenly distributed eutectic carbides and coarse austenite grains in M50 high-temperature bearing steel. It is expected that the microstructure refinement will achieve new breakthroughs in the tensile strength and rotational bending fatigue strength of M50 steel, providing a new alloying strategy for developing long-life, high-reliability high-temperature bearing steel. After heat treatment, the grain size of the alloyed bearing steel of this invention is improved from grade 9.0 in traditional 8Cr4Mo4V bearing steel to grade 10.0 or higher in the invented steel; the largest single primary carbide is reduced from 30-50 μm in traditional steel to below 20 μm; the tensile strength at room temperature of the invented steel is ≥2800 MPa, the yield strength is ≥2400 MPa, and the rotational bending fatigue strength (R = -1, 10) is... 7 ≥1100MPa. Attached Figure Description

[0031] Figure 1 This is a carbide size distribution diagram of Example 2-F of the φ15mm bar of the present invention.

[0032] Figure 2 The carbide size distribution diagram is for Comparative Example 1 of a φ15mm bar.

[0033] Figure 3 This is a grain size distribution diagram of Example 2-F of the φ15mm bar of the present invention.

[0034] Figure 4 This is a grain size distribution diagram of a φ15mm bar as shown in Comparative Example 1.

[0035] Figure 5 This is a rotational bending fatigue curve of Example 2-F of the φ15mm bar of the present invention. Detailed Implementation

[0036] To verify the beneficial effects of the technical solution of this invention, specific embodiments and comparative examples are provided below. Examples 1-3 employ the Nb-Al composite alloying design of this invention. Nb-Al alloying exhibits significant potential for microstructure refinement by influencing the solidification and solid-state phase transformation processes. Studies have shown that Nb, as a carbide modifier and grain refiner, has been applied in high-speed steel, cast iron, and other steel grades. In the early stages of solidification, the strong affinity of Nb for carbon preferentially leads to the formation of fine NbC particles. These particles can act as heterogeneous nucleation sites for primary carbides, increasing the nucleation rate and refining their size, while also consuming carbon in the molten steel, effectively reducing the volume fraction of coarse eutectic carbides and disrupting their continuous network structure. Furthermore, nanoscale NbC particles strongly pin grain boundaries, while dissolved Nb atoms generate a strong solute dragging effect; the two synergistically and significantly inhibit the growth trend of austenite grains at high temperatures. During high-temperature diffusion, Nb weakens the thermal stability of M2C eutectic carbides by reducing their W and V content, promoting diffusion decomposition at high temperatures. Al has a similar effect to Nb, promoting the precipitation of AlN particles and providing grain boundary pinning force at high temperatures. In gear carburizing steel, the combined addition of Nb and Al can form Nb(C,N) composite precipitates with higher thermal stability than single AlN or NbC, widening the temperature window for grain boundary pinning, thereby achieving more stable austenite grain refinement. Furthermore, Al reduces the volume fraction of eutectic carbides by increasing the carbon content at the eutectic point and promotes their uniform distribution. Due to the strong affinity between Al and C, a small amount of Al replaces some Mo in the M2C lattice, and lattice distortion leads to reduced thermal stability and promotes high-temperature decomposition. Importantly, Al, as a non-carbide-forming element, is mainly dissolved in the matrix. Compared to Fe, Al has a greater difference in electronegativity with alloying elements Mo and V, increasing the solid solubility of alloying elements in the matrix during solidification, while reducing their concentration in the residual molten steel at the end of solidification, inhibiting the eutectic reaction and thus reducing the size and volume fraction of eutectic carbides. Al can also increase the activity and diffusion coefficient of carbon in the matrix, promote the precipitation of fine carbides, and enhance the diffusion of carbon from eutectic carbides to the matrix at high temperatures, thus accelerating their eutectoid decomposition.

[0037] The comparison ratio uses the traditional M50 steel composition, and the smelting process for both is consistent, as detailed below:

[0038] 1)Chemical composition

[0039] Table 1. Actual test composition (wt%) of the steels used in Examples 1-3 and the control steel.

[0040]

[0041] 2) Steelmaking and heat treatment processes

[0042] Examples 1-3 and the comparative examples of this invention all employ the same steelmaking process: including smelting and solidification, high-temperature diffusion and forging, normalizing treatment, spheroidizing annealing heat treatment, quenching and tempering heat treatment. The specific steps are as follows: First, smelting and solidification are carried out using a 50kg vacuum induction furnace. After the ingot is demolded, it is placed in a 700℃ furnace for annealing for 5 hours, followed by furnace cooling to 500℃ and air cooling to eliminate casting stress. Subsequently, the billet undergoes high-temperature diffusion treatment, holding it at a temperature range of 1200℃ for no less than 8 hours to fully eliminate elemental segregation. Next, the forging process is carried out, forging into a Φ15 diameter bar within a temperature range of 900-1200℃. The forging start temperature is no less than 1150℃, and the final forging temperature is no less than 900℃. After forging, it is quickly placed in a 680℃ heat treatment furnace for furnace cooling to 500℃. After reaching 00℃, the material is air-cooled to room temperature. Then, normalizing is performed by heating the forging to 1050℃ and holding for 0.5 hours, followed by air cooling to refine the grains and improve the uniformity of the microstructure. Next, the normalized steel undergoes spheroidizing annealing, holding at 890℃ for 8 hours to fully spheroidize, followed by furnace cooling at 20℃ / h to 550℃ before air cooling to obtain a uniform spheroidized pearlite microstructure. Finally, segmented heating, quenching, and tempering are performed. The material is preheated at 800–850℃ for 0.5–1 hour, then heated to 1050–1150℃ and held for 15–40 minutes to complete austenitization, followed by oil quenching. Finally, it is tempered 2–4 times at 530–550℃ for 2 hours to obtain a stable tempered martensite microstructure. The heat treatment process of this embodiment is shown in Table 2.

[0043] Table 2. Heat treatment processes used for Examples 1-3 and comparative steels.

[0044]

[0045] Samples were taken from spheroidized annealed bars. Examples 1-3 and the comparative example of this invention underwent different quenching and tempering processes to study the effects of different quenching temperatures, quenching holding times, and tempering cycles. Example 1 had 4 groups (A-D), Example 2 had 2 groups (E-F), and Example 3 (G) and Comparative Example 1 (H) each had one group. Groups A-H all used a furnace-assisted staged heating method: first, the temperature was raised to 850℃ within 30 minutes and isothermally heated for 30 minutes; then, the temperature was raised to the quenching temperature within 30 minutes for solution treatment followed by oil quenching; finally, multiple tempering treatments were performed at 550℃. Specific parameters are shown in Table 2 above.

[0046] 3) Performance Characterization

[0047] The hardness of the tempered steel was tested using a Rockwell hardness tester; the mechanical properties were evaluated using the room temperature standard tensile test according to GB / T228.1-2021; and the room temperature twisting fatigue strength was evaluated using a QBWP-6000 simply supported rotating bending fatigue testing machine under a stress ratio R = -1. The performance results for groups A to H are listed in Table 3. The test results show that the maximum carbide size decreased from 34.30 μm (Comparative Example 1) in conventional steel composition to a minimum of 13.46 μm (Example 1-B); the grain size increased from grade 9 to grade 10.94 (Example 1-A); the maximum tensile strength reached 2924 MPa (Example 2-E); the maximum yield strength reached 2504 MPa (Example 2-F); and the room temperature twisting fatigue strength reached 1153 MPa (Example 2-F).

[0048] Table 3 Properties of the Invented Steel and Comparative Steels

[0049]

[0050] The microstructures of Example 2-F and Comparative Example 1 were compared under the same quenching and tempering regimes.

[0051] Figure 1 and Figure 2 These are primary carbide size distribution diagrams for Examples 2-F and Comparative Example 1, respectively. Figure 3 and Figure 4 The images show the original austenite grain morphology of Example 2-F and Comparative Example 1 after oil quenching at 1110℃. The images show that Example 2-F, after Nb and Al alloying, exhibits a more uniform and finer carbide distribution, with significantly more grains and a finer average grain size. According to national standard GB / T 6394-2017, its average grain size is 8.55 μm and 14.13 μm, respectively. After reasonable composition and process control, all the invented steels (A-G) achieved excellent microstructure properties with a grain size ≥10, a maximum carbide size ≤20 μm, a tensile strength ≥2800 MPa at room temperature, and a yield strength ≥2400 MPa. In this example, "ultra-high strength" refers to both high tensile strength and yield strength, and "double refinement" refers to both small grain size and maximum carbide size. Figure 5 The image shows the swirl bending fatigue curve of Group 2-F in Example 2.

[0052] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of the present invention, and are not intended to limit the scope of protection of the present invention. Although the present invention has been described in detail with reference to preferred embodiments, those skilled in the art should understand that modifications or equivalent substitutions can be made to the technical solutions of the present invention without departing from the essence and scope of the technical solutions of the present invention.

Claims

1. A high-strength, double-refined high-temperature bearing steel, characterized in that: The chemical composition of the bearing steel, by mass percentage, includes C: 0.75–0.85%, Cr: 3.75–4.30%, Mo: 4.00–4.50%, V: 0.90–1.10%, Si: ≤0.35%, Mn: ≤0.35%, Nb: 0.05–0.15%, Al: 0.5–1.0%, O ≤9ppm, Ti ≤0.002%, Ce+La: 0.03–0.05%, with the balance being iron and other unavoidable impurities.

2. The ultra-high strength double-refined high-temperature bearing steel according to claim 1, characterized in that: The bearing steel has a grain size ≥10, a maximum carbide size ≤20μm, a tensile strength ≥2800MPa at room temperature, a yield strength ≥2400MPa, and a rotational bending fatigue strength (R=-1、10 7 ≥1100MPa.

3. The method for preparing ultra-high strength double-refined high-temperature bearing steel according to claim 1 or 2, characterized in that, Includes the following steps: (1) Smelting and solidification After smelting and casting, the castings are placed in a 700℃ furnace for annealing for 5 hours, and then cooled in the furnace to 500℃ and air-cooled to eliminate casting stress. (2) High-temperature diffusion and forging The billet is heated to 1200-1250℃ and held for no less than 6 hours for high-temperature homogenization treatment to fully eliminate element segregation. Then it is forged in the range of 900-1200℃ to form a long bar. After forging, it is quickly placed in a 680℃ heat treatment furnace and cooled to 500℃ before being air-cooled to room temperature. (3) Normalizing treatment Forged bars are normalized to eliminate network carbides by holding at 1000–1100℃ for 0.5–1h, and then air-cooled after being removed from the furnace to refine the grains and improve the uniformity of the microstructure. (4) Spheroidizing annealing The normalized bars are subjected to spheroidizing annealing. The specific process is as follows: the temperature is raised to 850-890℃ in the furnace and held for 6-8 hours, then the furnace is cooled to 550℃ at a rate of 20℃ / h and then air-cooled to obtain a uniform spheroid pearlite structure. (5) Quenching and tempering The spheroidized annealed bearing steel is subjected to quenching and tempering treatment. The quenching process is a furnace-stage heating process. The specific quenching and tempering process is as follows: first, the temperature is raised to 800-850℃ and held for 0.5-1h, then the temperature is raised to 1050-1150℃ and held for 15-40min to complete austenitization, followed by oil quenching. Finally, the steel is tempered 2-4 times at 530-550℃ for 2h each time to obtain a stable tempered martensite structure.

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