A high-temperature, high-strength TZM alloy mold material, its forming method, and its application.

Abstract

This invention relates to the field of TZM alloy materials, and discloses a high-temperature, high-strength TZM alloy mold material and its forming method, including ball milling, powder mixing, cold isostatic pressing and hydrogen atmosphere sintering, hot isostatic pressing treatment, multi-fire upsetting forging, and stress-relief annealing. In the multi-fire upsetting forging, the deformation exhibits a single-peak change along the forging sequence, first increasing and then decreasing, with lower values ​​at both ends and higher values ​​in the middle section, and the direction is unidirectional forging along the axial direction. This process hardens the material internally, resulting in a high-density substructure deformed microstructure. The invention also includes the application of this material in the preparation of large-size vacuum isothermal forging molds. This invention overcomes the difficulties in preparing large-size TZM alloy molds, producing a synergistic strengthening effect between the substructure and the second phase. The resulting TZM alloy mold material meets the strength and plasticity requirements of mold materials for hot-end components such as powder superalloys and TiAl intermetallic compounds under vacuum isothermal forging conditions of 1000~1200℃.

Description

Technical Field

[0001] This invention relates to the field of TZM alloy materials, specifically to a high-temperature, high-strength TZM alloy mold material, its forming method, and its application. Background Technology

[0002] Molybdenum is a typical refractory metal with a melting point of approximately 2620℃ and poor plasticity during room temperature processing. TZM alloy is a special die material for vacuum isothermal forging of hot-end components of powder metallurgy high-temperature alloys and TiAl intermetallic compounds. When preparing large-size TZM alloy die materials, problems such as cracking, uneven density, and coarse grains are prone to occur during powder metallurgy sintering and forging.

[0003] For large-size vacuum isothermal forging dies, they often need to be used for a long time at 1000℃~1200℃. The requirements for the internal density, microstructure uniformity and high-temperature mechanical properties of the alloy billet are significantly higher than those for ordinary plates or bars. Therefore, the material needs to have high strength and thermal fatigue resistance in this temperature range.

[0004] Existing TZM alloys are mostly concentrated in conventional products such as plates and bars. Their preparation processes are difficult to directly meet the requirements of high-temperature strength and high-temperature stability for large-size and complex molds. Therefore, the development of high-temperature and high-strength TZM alloy materials and their forming methods suitable for large-size vacuum isothermal forging molds has important engineering application value. Summary of the Invention

[0005] The purpose of this invention is to provide a high-temperature, high-strength TZM alloy mold material and its forming method, so as to solve the technical problems of high melting point, easy oxidation, and difficulty in forming large-size TZM alloy molds in the prior art.

[0006] To solve the above-mentioned technical problems, the present invention specifically provides the following technical solution: A high-temperature, high-strength TZM alloy mold material, wherein the material contains both a substructure and a second phase, and the substructure and the second phase synergistically enhance the properties of the material. The material has a tensile strength of 535 MPa, a yield strength of 522 MPa, and an elongation of 12% at 1200℃. After 40 cycles, the strength retention rate exceeds 96%, which meets the requirements of vacuum isothermal forging of the mold at 1000-1200℃.

[0007] Furthermore, by mass percentage, the material comprises 0.020%~0.030% C, 0.50%~0.55% Ti, 0.08%~0.11% Zr, with the balance being Mo and unavoidable impurities.

[0008] Furthermore, the material contains 0.015% to 0.020% O, <0.0015% Al, <0.004% Fe, and <0.0005% Ni.

[0009] To address the aforementioned technical problems, the present invention further provides the following technical solution: Application of a high-temperature, high-strength TZM alloy mold material in the preparation of large-size vacuum isothermal forging molds.

[0010] To address the aforementioned technical problems, the present invention further provides the following technical solution: A forming method for a high-temperature, high-strength TZM alloy mold material includes the following steps: S100. Molybdenum powder, titanium hydride powder, zirconium hydride powder and carbon powder are mixed and then subjected to ball milling, powder mixing, cold isostatic pressing and sintering in a hydrogen atmosphere in sequence to improve the uniformity of powder mixing and the density of sintered blanks to obtain sintered blanks. S200. The sintered billet is subjected to hot isostatic pressing to further improve the material density, promote the diffusion of interfacial elements, improve uniformity, avoid stress concentration caused by pores during subsequent forging, which would lead to billet cracking, and improve subsequent forging efficiency and stability. S300. The billet after hot isostatic pressing is subjected to multiple upsetting forgings. The deformation amount shows a single-peak change with an initial increase followed by a decrease along the forging sequence. The deformation is low at both ends and high in the middle section. The direction is unidirectional forging along the axial direction, which promotes the formation of a high-density substructure inside the material and produces a composite strengthening effect together with the dispersed second phase. After forging, S400 is subjected to stress-relief annealing to obtain a large-size high-temperature and high-strength TZM alloy mold.

[0011] Furthermore, in step S100, the ball mill uses a three-dimensional mixer, with grinding balls of 3mm and 10mm diameter used in combination, and the ball milling time is 6~8h; the powder mixing time is 3~4h.

[0012] Further, in step S100, the pressure of the cold isostatic pressing is 180~250MPa, and the holding time is 1~2h; the temperature of the hydrogen atmosphere sintering is 2000~2100℃, and the sintering time is 4~8h.

[0013] Furthermore, in step S200, the temperature of the hot isostatic pressing is 1200~1300℃, the pressure is 150~170MPa, the time is 2~4h, and the protective atmosphere is argon.

[0014] Furthermore, in step S300, the multi-fire upsetting forging is performed three times, with the material deformation amount in each fire being 10-15%, 25-30%, and 20-25%, and the total deformation amount being controlled at 60%-70%. The initial forging temperature for each fire is 1450-1500℃, and the final forging temperature is 1300-1350℃. After each fire forging, the material is returned to the furnace for reheating, and both the preheating furnace and the reheating furnace are hydrogen-protected heat treatment furnaces.

[0015] Furthermore, in step S400, the stress-relief annealing temperature is 1100~1200℃, the holding time is 5~10h, and the annealing atmosphere is hydrogen protection.

[0016] Compared with the prior art, the present invention has the following advantages: This invention overcomes the difficulties in preparing large-size TZM alloy molds, breaking through the bottlenecks of traditional methods such as easy cracking, uneven density, and coarse grains. The TZM alloy mold material prepared by this invention has a tensile strength of 535MPa, a yield strength of 522MPa, an elongation after fracture of 12%, and a shrinkage after fracture of 39% under vacuum conditions at 1200℃. This invention can meet the strength and plasticity requirements of mold materials for hot-end components such as powder superalloys and TiAl intermetallic compounds under vacuum isothermal forging conditions at 1000~1200℃. This invention employs multi-stage temperature-controlled deformation forging, with the deformation exhibiting a single-peak variation that first increases and then decreases along the forging sequence, being low at both ends and high in the middle, and unidirectional along the axial direction. This introduces high-density subgrain boundaries to promote the formation of a high-density substructure within the material, which, together with the dispersed second phase, produces a composite strengthening effect. Simultaneously, subsequent stress-relief annealing eliminates forging stress without damaging the deformed structure, achieving a balance between high-temperature strength and thermal fatigue resistance. Attached Figure Description

[0017] To more clearly illustrate the embodiments of the present invention or the technical solutions in the prior art, the drawings used in the description of the embodiments or the prior art will be briefly introduced below. Obviously, the drawings in the following description are merely exemplary, and those skilled in the art can derive other embodiments based on the provided drawings without creative effort.

[0018] Figure 1 A schematic flowchart illustrating the forming method of the high-temperature, high-strength TZM alloy mold material provided in this application embodiment; Figure 2 This is a SEM image of the second-phase particles provided in Embodiment 1 of this application; Figure 3 This is a TEM image of the second-phase particles provided in Example 1 of this application; Figure 4This is a large-angle and small-angle grain boundary distribution diagram of the high-temperature and high-strength TZM alloy mold material provided in Embodiment 1 of this application; Figure 5 TEM images of subgrain boundaries and dislocations provided in Embodiment 1 of this application; Figure 6 The grain boundary feature map and grain orientation map with 20% deformation provided in Embodiment 1 of this application; Figure 7 The grain boundary features and grain orientation diagram for 40% deformation provided in Embodiment 1 of this application; Figure 8 The grain boundary features and grain orientation diagram for 60% deformation provided in Embodiment 1 of this application; Figure 9 This is a stress distribution diagram (15%) at the end of the first heat deformation provided in Embodiment 3 of this application; Figure 10 This is a stress distribution diagram at the end of the second heat deformation (45%) provided in Embodiment 3 of this application; Figure 11 The stress distribution diagram at the end of the 3-fire deformation (70%) provided in Embodiment 3 of this application. Detailed Implementation

[0019] To make the objectives, technical solutions, and advantages of this application clearer, the technical solutions of this application will be clearly and completely described below with reference to the accompanying drawings. Obviously, the described embodiments are only some embodiments of this application, not all embodiments. Based on the embodiments of this application, all other embodiments obtained by those skilled in the art without creative effort are within the scope of protection of this application.

[0020] A high-temperature, high-strength TZM alloy mold material is provided. The material has a tensile strength of 535 MPa, a yield strength of 522 MPa, and an elongation of 12% at 1200℃. After 40 cycles, the strength retention rate exceeds 96%, which meets the requirements for vacuum isothermal forging of molds at 1000-1200℃.

[0021] The material contains, by mass percentage, 0.020%~0.030% C, 0.50%~0.55% Ti, 0.08%~0.11% Zr, with the balance being Mo and unavoidable impurities.

[0022] This material contains a small amount of carbon. Excessive carbon content is detrimental to plasticity, as carbides tend to aggregate and grow at grain boundaries, forming coarse carbides. However, by adding an appropriate amount of carbon, a dispersed second phase can be formed, which can play a role in strengthening the material.

[0023] Furthermore, the material contains 0.015% to 0.020% O, <0.0015% Al, <0.004% Fe, and <0.0005% Ni.

[0024] O, Al, Fe, Ni, etc. are all impurities.

[0025] like Figure 1 As shown, the present invention also provides a forming method for high-temperature, high-strength TZM alloy mold material, comprising the following steps: S100. Molybdenum powder, titanium hydride powder, zirconium hydride powder and carbon powder are mixed and then subjected to ball milling, powder mixing, cold isostatic pressing and sintering in a hydrogen atmosphere in sequence to improve the uniformity of powder mixing and the density of sintered blanks to obtain sintered blanks. Mo powder (particle size 1.5~3.0µm), TiH2 powder 0.55~0.70% (particle size 1.5~3.0µm), ZrH2 powder 0.11~0.14% (particle size 1.5~3.0µm), and C powder 0.045~0.060% (particle size 0.5~1.5µm). The purpose of increasing the C content is to form a carbide reinforcing phase. Ball milling: Put the mixed raw material powder into a three-dimensional mixer and ball mill it (using 3mm and 10mm grinding balls together) for 6-8 hours (this step is to initially mix, break up agglomerates, refine the powder, and improve uniformity).

[0026] Powder mixing: The ball-milled powder is sieved to remove the grinding balls, and then placed into a second three-dimensional mixer for homogenization mixing for 3-4 hours (the purpose of this step is: 1. to eliminate interference from the grinding balls and allow for more complete powder flow; 2. to break down some of the particle or flaky structures formed by the impact of the grinding balls during the ball milling process and to re-disperse and mix them evenly). Cold isostatic pressing: pressure 180~250MPa, holding time 1-2h, uniformly pressurizes the powder blank to initially improve the density of the blank; Sintering: Temperature 2000~2100℃, sintering for 4~8h, hydrogen atmosphere sintering furnace to improve density, hydrogen atmosphere to remove oxygen and control the oxygen content of billet.

[0027] S200. The sintered billet is subjected to hot isostatic pressing to further improve the material density, promote the diffusion of interfacial elements, and improve uniformity; avoid stress concentration caused by pores during subsequent forging, which would lead to billet cracking, and improve the efficiency and stability of subsequent forging. Temperature 1200~1300℃, pressure 150~170MPa, time 2~4h, argon protection, argon protection to prevent oxidation and nitriding.

[0028] S300 involves subjecting the hot isostatically pressed billet to multiple upsetting forgings. The deformation amount exhibits a single-peak variation along the forging sequence, first increasing and then decreasing, with lower values ​​at both ends and higher values ​​in the middle. For example, the deformation amount in each forging is controlled successively at 10~15% (for a large aspect ratio, to avoid instability in the first forging), 25~30%, and 20~25%, with unidirectional forging along the axial direction and no special angle. The initial forging temperature for each forging is 1450~1500℃, the final forging temperature is 1300~1350℃, and the total deformation amount is 60%~70%. This process hardens the internal structure of the material and strengthens its substructure, resulting in a high-density substructure deformed microstructure. The furnace is reheated between each heat treatment cycle, and the preheating furnace and the reheating furnace are the same hydrogen-protected heat treatment furnace.

[0029] After forging, S400 is subjected to stress-relief annealing to obtain a large-size high-temperature and high-strength TZM alloy mold.

[0030] The cold isostatic pressing pressure is controlled at 180~250MPa and held for 1~2h, which enables the powder particles to be fully rearranged and in close contact under uniform pressure in all directions, thereby increasing the density of the pressed blank, reducing the density gradient inside the blank, and avoiding the local porosity that is prone to occur in unidirectional pressing.

[0031] Holding the powder under pressure for 1-2 hours promotes more thorough and uniform contact between powder particles, providing a good foundation for the formation, growth, and pore shrinkage of sintering necks during subsequent sintering. This effectively reduces porosity after sintering and increases the density of the billet. If the pressure is too low, the powder will not be sufficiently compacted, resulting in limited contact area between particles, which is not conducive to subsequent densification. If the pressure is too high, it may lead to local over-compaction and stress concentration, affecting the uniformity of the sintered billet.

[0032] Sintering in a hydrogen atmosphere at a temperature controlled between 2000 and 2100°C for 4 to 8 hours effectively promotes atomic diffusion in the Mo matrix, neck growth, and pore closure, thereby improving density. Only at higher temperatures can surface diffusion and grain boundary diffusion significantly drive neck growth and pore shrinkage.

[0033] A sintering time of 4–8 hours ensures sufficient diffusion, neck growth, and pore migration, resulting in a high level of density. If the sintering temperature is too low or the holding time is too short, the Mo atom diffusion rate will be insufficient, the sintering neck development will be limited, and the pores will be difficult to fully shrink and close, which is not conducive to obtaining a highly dense microstructure. If the temperature is too high or the time is too long, although diffusion will be further enhanced, it will easily cause grain growth too quickly and coarsening of the strengthening phase, which is detrimental to the stability of the microstructure and the overall mechanical properties.

[0034] Hot isostatic pressing can further close the pores, improve the bonding state of the internal interface of the sintered billet, and promote the homogenization of local element diffusion, thereby further increasing the density of the billet and helping to reduce the risk of subsequent forging cracks.

[0035] Furthermore, in step S300, the multi-fire upsetting forging is performed three times, with the material deformation amount in each fire being 10-15%, 25-30%, and 20-25%, and the total deformation amount being controlled at 60%-70%. The initial forging temperature for each fire is 1450-1500℃, and the final forging temperature is 1300-1350℃. After each fire forging, the material is returned to the furnace for reheating, and both the preheating furnace and the reheating furnace are hydrogen-protected heat treatment furnaces.

[0036] The first deformation was 15%, mainly because the sintered billet was a long bar and needed to be forged into a round billet. The deformation amount in the first firing needed to be controlled to avoid deformation instability. The hot deformation of this TZM alloy was mainly recovery, and recrystallization was not obvious within the range of deformation temperature and strain rate.

[0037] The first forging process avoids forging instability and prevents incomplete forging in subsequent stages; the intermediate forging processes allow for better deformation, but the greater the deformation, the more difficult it becomes to deform. If the forging time is too long, the temperature drops and accumulated stress can easily cause cracking. Therefore, the deformation in the final forging process is relatively smaller.

[0038] The total deformation reaches 60%~70%, which can further enhance the strengthening effect of subgrain boundaries.

[0039] Furthermore, during the forging process, it is necessary to ensure that the initial forging temperature of each forging pass is 1450-1500℃ and the final forging temperature is 1300-1350℃. If the temperature exceeds 1500℃, the grains are prone to coarsening; if the temperature is below 1300℃, the machinability will decrease (generally, the minimum temperature will not be lower than 1200℃, and the triaxial tensile stress zone is prone to cracking). This improves forging efficiency, reduces surface oxidation during the forging process, and also allows for quick transfer back to the furnace for reheating. Additionally, the proportion of small-angle grain boundaries will decrease slightly (approximately 3%).

[0040] Furthermore, in step S400, the stress-relief annealing temperature is 1100~1200℃, the holding time is 5~10h, and the annealing atmosphere is hydrogen protection. Without changing the characteristics of the hot deformation structure, the residual forging stress is eliminated and the machinability is improved.

[0041] The present invention further provides the application of the forming method of high-temperature and high-strength TZM alloy mold material in the preparation of large-size vacuum isothermal forging molds.

[0042] For large-sized die blanks, deformation uniformity and forging stability can be improved by adjusting the number of forging passes and reducing the deformation per pass. The size range and process relationship for large-sized dies are as follows: The above process parameters are used for molds with a diameter of 250~450mm.

[0043] The mold diameter is 450~550mm and it is forged in 4 heats (the deformation is controlled by 15%+20%+15%+15% in sequence, for a total of 65%).

[0044] For molds with diameters of 550~700mm, a 5-pass forging process is used (the deformation is controlled at 10%+15%+15%+15%+10% in sequence, for a total of 65%). For molds with a diameter of 700~900mm, a 6-pass forging process is used (the controlled deformation is 10%+15%+15%+12%+10%+8% in sequence, with a total deformation of 70%).

[0045] It is worth noting that the larger the size, the greater the deformation, ensuring thorough forging.

[0046] Example 1

[0047] Raw material ratio: By mass percentage, C 0.025%, Ti 0.52%, Zr 0.10%, with the balance being Mo powder (particle size 2.0µm), TiH2 powder (particle size 2.0µm), ZrH2 powder (particle size 2.0µm), and C powder (particle size 1.0µm).

[0048] Preparation process: 1. Ball milling: Three-dimensional mixer, with 3mm and 10mm grinding balls, ball milling for 7 hours.

[0049] 2. Mixing powder: After sieving, homogenize the powder and mix for 3.5 hours.

[0050] 3. Cold isostatic pressing: pressure 220MPa, pressure holding for 1.5h.

[0051] 4. Sintering: hydrogen atmosphere, temperature 2050℃, hold for 6 hours.

[0052] 5. Hot isostatic pressing: temperature 1250℃, pressure 160MPa, time 3h, argon protection.

[0053] 6. Forging: Multi-fire upsetting forging with deformation amounts of 15%, 25%, and 20% (decimal points omitted), for a total deformation amount of 60%; initial forging temperature of 1480℃ and final forging temperature of 1320℃; hydrogen protection for reheating between each forging.

[0054] 7. Annealing: under hydrogen protection, at 1150℃, for 7.5 hours.

[0055] Component analysis (chemical analysis): C: 0.020~0.030%, Ti: 0.50~0.55%, Zr: 0.08~0.11%, O: 0.015~0.020%, Al: <0.0015%, Fe: <0.004%, Ni: <0.0005%.

[0056] Example 2:

[0057] Raw material ratio: The composition is C 0.030%, Ti 0.55%, Zr 0.11%, with the balance being Mo powder (particle size 1.5µm), TiH2 powder (particle size 1.5µm), ZrH2 powder (particle size 1.5µm), and C powder (particle size 0.5µm).

[0058] The preparation process is the same as in Example 1, except that the deformation amount per forging is approximately 15%, 25%, and 20%, with a total deformation of 60%; the initial forging temperature is 1500℃ and the final forging temperature is 1350℃.

[0059] Example 3:

[0060] Raw material ratio: The composition is 0.020% C, 0.50% Ti, 0.08% Zr, with the balance being Mo powder (particle size 3.0µm), TiH2 powder (particle size 3.0µm), ZrH2 powder (particle size 3.0µm), and C powder (particle size 1.5µm).

[0061] The preparation process is the same as in Example 1, except that the deformation amount per forging is approximately 15%, 30%, and 25%, with a total deformation of 70%; the initial forging temperature is 1500℃ and the final forging temperature is 1350℃.

[0062] Comparative Example 1: Raw material ratio: Same as in Example 1.

[0063] Preparation process: Except for the sintering temperature being changed to 1800℃, the other steps are the same as in Example 1.

[0064] Comparative Example 2 Raw material ratio: Same as in Example 1.

[0065] Preparation process: The remaining steps are the same as in Example 1, but hot isostatic pressing is not performed.

[0066] Comparative Example 3: Raw material ratio: Same as in Example 1.

[0067] Preparation process: Forging is carried out in a single heat, with a deformation amount of 20% and a total deformation amount of 20%; initial forging temperature is 1480℃ and final forging temperature is 1320℃; the remaining steps are described in Example 1.

[0068] Example of detection: 1. Material characterization and testing: (1) Element ratio The element ratios in Example 1 are shown in Table 1: Table 1

[0069] (2) Microstructural characteristics analysis of Example 1 ① Second phase particles Tissue photographs of the second-phase particles are shown below. Figure 2 and Figure 3 As shown in the figure, there are two types of second phases: one is a polygonal micron-sized (0.5~3.5µm) Mo-rich carbide, and the other is a near-spherical nano-sized (50~200nm) Mo-rich (Zr) carbon oxide. The volume fraction of the second phase is about 3%, and the ratio of the two second phases is about 4:1.

[0070] ②Grain boundaries and substructure Large-angle and small-angle grain boundary distribution diagrams are as follows Figure 4 As shown, when a certain number of subgrain boundaries are reached, the strengthening effect of the subgrain boundaries is significant.

[0071] TEM images of subgrain boundaries and dislocations are as follows Figure 5 As shown, the TEM field of view is relatively small. Using the EBSD method, it was observed that small-angle grain boundaries account for approximately 70%. The microstructure contains some intact subcells and a large number of dislocation walls. According to calculations, solid solution strengthening, second-phase strengthening, substructure strengthening, and grain boundary strengthening contribute approximately 5.6%, 47.2%, 40.4%, and 6.7% to the total strengthening increment, respectively. Among these, second-phase strengthening mainly originates from nanoscale second-phase particles, while substructure strengthening mainly originates from a high proportion of small-angle grain boundaries and dislocation walls. It is evident that Example 1 achieved strengthening.

[0072] Example 1 shows a microstructure characterized by synergistic reinforcement of substructure and second phase, in which a high proportion of small-angle grain boundaries plays an important role in improving high-temperature strength.

[0073] ③ Segmented deformation hot compression test A segmented deformation hot compression test was conducted on Example 1. The grain boundary characteristics and grain orientation diagrams for deformation amounts of 15%-60% are shown below. Figures 6-8 As shown. The simulated stress diagram of the deformation of Example 1 at the end of the 1-3 fire pattern is shown below. Figure 9-11 As shown.

[0074] Figure 6-8 These correspond to cumulative deformations of 15%, 40%, and 60%, respectively, reflecting the evolution of substructure, grain boundary type, and texture of the material during hot deformation.

[0075] Small-angle grain boundaries (LAGBs) begin to appear after 15% deformation (corresponding to approximately 15% single-stage deformation after the first firing). This 15% deformation can approximately reflect the microstructure characteristics of the low-deformation stage of the first firing. It is evident that the initial deformation is mainly characterized by dislocation tangles and the formation of local subgrain boundaries. The substructure is still unstable and prone to recovery during high-temperature service, resulting in limited strengthening effect. Its main function is to accumulate dislocation density for subsequent deformation.

[0076] It is evident that the first firing process uses a small deformation of 15% to avoid billet instability and to provide "pre-energy storage" for subsequent microstructure evolution.

[0077] With a 40% deformation amount (corresponding to a cumulative deformation of approximately 15% + 25% after the second forging), the number of small-angle grain boundaries increases significantly, and dislocation walls and subgrain boundary networks appear in some areas. Grains begin to elongate along the forging direction, exhibiting slight deformation texture. It is evident that at this point, the deformation enters a rapid substructure development stage, with small-angle grain boundaries gradually evolving into stable subgrain boundaries, enhancing the strengthening effect. Grain elongation also contributes to subsequent recrystallization or substructure stabilization.

[0078] The intermediate firing process employs a larger deformation of 25%, which effectively introduces high-density subgrain boundaries and is the main source of overall strengthening contribution (40.4%).

[0079] With 60% deformation (corresponding to a cumulative deformation of approximately 15% + 25% + 20% after the third firing cycle), the proportion of small-angle grain boundaries reaches approximately 70% (EBSD measurement), forming a complete subcellular structure and numerous dislocation walls. The grains exhibit a strong banded deformation texture, and the proportion of large-angle grain boundaries (HAGBs) decreases. At this point, after deformation, the material's substructure strengthening reaches its peak. The subgrain boundaries are stable and uniformly distributed. The banded structure helps to change the crack propagation path and improve high-temperature fatigue resistance. The formation of the texture means that the grain orientation tends to be consistent, which is beneficial for transferring loads along specific directions.

[0080] It can be seen that the final 20% deformation amount, without excessively increasing the risk of cracking, pushes the total deformation amount to 60%~70%, thus achieving substructure saturation strengthening.

[0081] By increasing the deformation amount in a "small-large-medium" manner through multiple firings, a high-density substructure can be systematically constructed, avoiding instability or cracking caused by a large deformation at once.

[0082] (3) Stress simulation in Example 3 Figures 9-11 The simulation results correspond to the internal stress distribution at the end of the first fire (15%), the second fire (45% cumulative), and the third fire (70% cumulative), respectively, and were obtained using finite element simulation (Deform-3D).

[0083] At the end of the first firing (approximately 15% deformation), stress is mainly concentrated in the central area of ​​the billet, with lower stress on the surface and ends. The material is still in the elastic + slight plastic stage, and the stress distribution is relatively uniform, with no obvious stress concentration, which helps to avoid initial cracking. Therefore, 15% is a reasonable amount for the initial firing deformation.

[0084] After the second forging cycle (approximately 45% cumulative deformation), stress increases significantly and its distribution area expands. The high-stress zone remains primarily in the center but extends towards the edges, highlighting the effectiveness of hot isostatic pressing at this stage. Therefore, it is evident that the intermediate forging cycles require strict control of the final forging temperature (≥1300℃) and reheating in the furnace to soften the material and alleviate stress.

[0085] After the third firing (approximately 70% cumulative deformation), the internal stress distribution of the billet remained within a controllable range, and no obvious localized abnormal stress concentrations were observed. Although Example 3 had the largest final deformation, a deformation of 25% was used in the final firing instead of a larger one precisely to avoid stress exceeding the material limit while ensuring that the total deformation met the standard.

[0086] This invention, through a "small-large-medium" deformation sequence, can gradually strengthen the structure while controlling stress accumulation and avoiding cracking caused by stress concentration.

[0087] In summary, under small deformation conditions, the substructure in the microstructure is mainly composed of discontinuous small-angle grain boundaries, with very few complete subcells. The substructure lacks stability and is prone to recovery under high-temperature service conditions, thus limiting its strengthening effect on the material. Under small deformation conditions, grain orientation does not exhibit obvious deformation texture characteristics, while under large deformation conditions, strong deformation texture and preferred orientation are formed. Furthermore, the banded deformed grains formed under large deformation facilitate changing their direction during crack propagation, thus hindering crack propagation.

[0088] This invention selects a deformation control sequence of 10-15%, 25-30%, and 20-25% to achieve safe billet preparation, avoid instability, efficiently introduce substructures, ultimately achieve saturation strengthening, and control stress. If all three forgings use 20%, although the total deformation remains unchanged, the lack of active deformation control during the first forging may lead to cracking. Furthermore, if the first forging uses 15%, the deformation progressively increases in subsequent forgings, potentially causing problems in the final forging. Figure 11 The peak stress in the material may exceed its strength, leading to cracking.

[0089] As can be seen, the single-peaked change with a rise and fall along the forging sequence adopted in Examples 1 to 3, with low deformation at both ends and high deformation in the middle, is to achieve a balance between substructure strengthening efficiency and stress control safety, and to improve high-temperature strength (1200℃ / 528MPa) and cyclic stability (91.3% retention rate after 40 cycles).

[0090] 2. Room temperature tensile properties test Implementation standard: GB / T228.1 Test method: (1) Measure the original cross-sectional area and original gauge length of the specimen before the test; (3) Install the specimen correctly in the test machine fixture to ensure axial force; (4) Install the extensometer according to the specified method and set the extensometer gauge length; (5) Zero the force value and deformation measurement system; (6) Set the test rate: 0.00025 / s before the end of yielding (strain control), 0.0067 / s after the end of yielding (displacement control); (7) Start loading and continuously record the force-extension curve and stress-strain curve; (8) Determine the yield point, specified plastic extension strength, tensile strength, etc. according to the curve; (9) After stretching to the fracture, remove the specimen; (10) Carefully align the two sections of the fractured specimen and measure the gauge length after fracture; (11) Measure the minimum cross-sectional area at the fracture surface to calculate the reduction of area.

[0091] Instruments and reagents used: In accordance with national standards, the testing instrument used is an electronic universal testing machine.

[0092] Test results The room temperature tensile properties are shown in Table 2: Table 2

[0093] As shown in Table 2, Example 2 performed best: tensile strength 927 MPa, yield strength 812 MPa, and elongation 18%. Examples 1 and 3 were slightly lower. Comparative Example 1 (low-temperature sintering): strength decreased significantly (805 / 690), and elongation was 5%. This is because the sintering temperature was low, resulting in poor density, numerous pores, and insufficient carbide formation.

[0094] Comparative Example 2 (without hot isostatic pressing): Strength 845 / 725, lower than the example, illustrating the importance of hot isostatic pressing in eliminating residual porosity and promoting element diffusion. Comparative Example 3 (single-fired low deformation): Extremely low strength 635 / 485, elongation 2%, indicating the absence of multi-fired cumulative deformation to introduce high-density subgrain boundaries, lack of substructure strengthening, and potential stress concentration.

[0095] In summary, the combined effects of multi-stage forging and hot isostatic pressing resulted in a highly dense microstructure with fine grains and high subgrain boundary density, achieving a good balance between strength and plasticity.

[0096] 3. High-temperature tensile property testing Implementation standard: GB / T228.2 Test steps: (1) Measure the original cross-sectional area and original gauge length of the specimen before the test; (2) Install the specimen correctly in the test machine fixture to ensure axial force; (3) Install the extensometer according to the specified method and set the extensometer gauge length; (4) Zero the force value and deformation measurement system; (5) Start heating the specimen to the specified test temperature of 1200℃ and keep it warm for 15min; (6) Set the test rate: 1mm / min; (7) Start loading and continuously record the force-extension curve and stress-strain curve; (8) Determine the yield point, specified plastic extension strength, tensile strength, etc. according to the curve; (9) Stop the heating device after stretching to fracture; (10) Remove the fractured specimen after the specimen has cooled safely; (11) Carefully align the two sections of the fractured specimen and measure the gauge length after fracture; (12) Measure the minimum cross-sectional area at the fracture point to calculate the reduction of area.

[0097] Instruments and reagents used: In accordance with national standards, the testing instrument adopted is a high-temperature tensile testing device.

[0098] Test results Table 3 shows the statistical table of high temperature tensile properties (1200℃). Table 3

[0099] As shown in Table 3, the trend is similar to that at room temperature. Example 2 is again the best (535 / 522, elongation 12%). Examples 1 and 3 are slightly lower. The performance of Comparative Examples 1, 2, and 3 decreases in that order. The strength of Comparative Example 3 is only 235 / 198.

[0100] This indicates that the substructures (small-angle grain boundaries, dislocation walls) formed during multi-stage forging can hinder dislocation movement at high temperatures, thus contributing to improved high-temperature strength. Simultaneously, the dispersed second-phase particles (Mo-rich carbides, (Mo,Zr)-rich carbon oxides) pin grain boundaries and dislocations at high temperatures, also playing a strengthening role. Without these microstructures, high-temperature strength would decrease dramatically.

[0101] 4. Hardness testing Implementation standard: GB / T4340.1 Test steps: (1) Turn on the hardness tester and the instrument will be automatically calibrated; (2) Select the Vickers test force of 1 kgf; (3) Place the sample on the hardness tester table, find the sample under the low magnification lens, make the surface to be tested clear, and adjust it to a suitable observation and loading position; (4) Select the first test point and ensure that the distance between the indentation position and the edge of the sample and the adjacent indentation meets the requirements; (5) Start loading, hold for a specified time after reaching the specified test force, and unload the test force after the end; (6) Observe the indentation morphology after unloading, and confirm that the indentation boundary is clear and there are no cracks, peeling, abnormal deformation or obvious asymmetry that affect the measurement; (7) Fine adjust the focus to ensure that the indentation boundary is clear and calibrate the position of the indentation diagonal; (8) Calculate the Vickers hardness value of the first test point; (9) Continue to repeat the test to complete the measurement.

[0102] Instruments and reagents used: Vickers hardness tester was used in accordance with national standards. Test results: The statistical table of Vickers hardness results is shown in Table 4: Table 4

[0103] As shown in Table 4, Example 2 had the highest hardness (357 HV1), while Examples 1 and 3 were approximately 315. Comparative Examples 1, 2, and 3 decreased sequentially (276, 292, and 243). Hardness is a comprehensive reflection of strength, density, and tissue state.

[0104] The high hardness of the embodiment originates from fine-grain strengthening, dispersion strengthening, and substructure strengthening.

[0105] 5. High-temperature cycling test Testing method: Vacuum high-temperature cyclic heating and cooling, heating from room temperature at 5℃ / s to 1200℃, holding at that temperature for 8 hours, and then cooling the furnace to room temperature. This is one cycle. Set the cycles to 10, 20, and 40 times.

[0106] Instruments and reagents used: tube furnace The high-temperature tensile properties (1200℃) of the high-temperature cycling tests in Examples 1-3 are shown in Tables 5-7: Table 5

[0107] Table 6

[0108] Table 7

[0109] As shown in Tables 5 to 7, the data from Example 1 shows that the tensile strength gradually decreased from 528 to 482, the yield strength decreased from 486 to 416, but the elongation increased significantly from 10% to 26%, and the shrinkage increased from 35% to 85%. The strength decreased, but the plasticity increased significantly.

[0110] Example 2 is more stable: the strength decreased only slightly from 535 to 519, the yield strength increased from 522 to 509, the elongation increased from 12% to 18%, the plasticity increased and the strength was maintained very well.

[0111] The strength of Example 3 decreased significantly (from 528 to 439), mainly due to the larger deformation, but the plasticity also increased. Comparatively, Example 2 exhibited the most stable microstructure and the best resistance to thermal fatigue. This is related to its superior process window (initial forging at 1500℃, final forging at 1350℃, total deformation 60%).

[0112] During high-temperature cycling, recovery and grain growth occur, leading to changes in the proportion of small-angle grain boundaries, resulting in a slight decrease in strength but an increase in plasticity. However, Example 2 maintains its strength better, indicating high substructural thermal stability.

[0113] In summary, this embodiment has the following advantages: 1. Excellent comprehensive mechanical properties: Whether at room temperature or at a high temperature of 1200℃, the examples all exhibited tensile strength, yield strength, and good plasticity far exceeding those of the comparative examples. In particular, Example 2, while ensuring high strength, also showed good elongation.

[0114] 2. Significant process synergy: The example adopted a combined densification process of "cold isostatic pressing + high-temperature sintering + hot isostatic pressing," which significantly improved the billet density, reduced residual porosity and internal defects, and lowered the risk of cracking during subsequent forging, providing a good billet foundation for the preparation of large-size TZM alloy mold materials.

[0115] 3. Microstructure Optimization: This invention employs multi-stage, temperature-controlled, and deformation-controlled forging, controlling the deformation amount to exhibit a single-peak variation along the forging sequence, initially increasing and then decreasing, with lower values ​​at both ends and higher values ​​in the middle. This reduces the risk of forging instability and cracking while introducing approximately 70% small-angle grain boundaries (subgrain boundaries) and numerous dislocation walls, achieving significant substructure strengthening, which is one of the main sources of the high-temperature strength in this embodiment. Simultaneously, a dispersed second phase is formed, comprising polygonal micron-sized Mo-rich carbides and near-spherical nano-sized (Mo, Zr)-rich carbon oxides. These second-phase particles hinder dislocation movement, enhancing substructure stability and achieving a synergistic effect of substructure strengthening and second-phase dispersion strengthening.

[0116] 4. Excellent high-temperature thermal fatigue performance: The results of high-temperature cycle testing show that the embodiments can still maintain high strength after multiple cycles of heating at 1200℃. In particular, the tensile strength retention rate of Embodiment 2 after 40 cycles exceeds 96%, indicating that its structure has good stability at high temperatures and is suitable for the service conditions of vacuum isothermal forging dies.

[0117] The embodiment 2 provided by the present invention exhibits the best comprehensive mechanical properties and thermal fatigue resistance under room temperature, 1200℃ and high temperature cyclic conditions. Its tensile strength at 1200℃ reaches 535MPa, yield strength 522MPa, elongation 12%, and strength retention rate exceeds 96% after 40 cycles. It can meet the requirements of mold materials for hot-end components such as powder superalloys and TiAl intermetallic compounds under vacuum isothermal forging conditions of 1000-1200℃.

[0118] This invention, through compositional control, multi-step densification treatment, and the introduction of multi-fire substructures into the forging process, achieves a microstructure with high-density subgrain boundaries and dispersed carbides, thus realizing an excellent balance of high-temperature strength, plasticity, and high-temperature stability. This provides an effective technical approach for the stable fabrication of large-size TZM alloy molds.

[0119] Finally, it should be noted that the above embodiments are only used to illustrate the technical solutions of this application, and not to limit them. Although this application has been described in detail with reference to the foregoing embodiments, those skilled in the art should understand that modifications can still be made to the technical solutions described in the foregoing embodiments, or equivalent substitutions can be made to some of the technical features. All such modifications or substitutions should be covered within the protection scope of this application, and should not cause the essence of the corresponding technical solutions to deviate from the spirit and scope of the technical solutions of the embodiments of this application. Therefore, the protection scope of this application should be determined by the protection scope of the claims.

Claims

1. A method for forming a high-temperature, high-strength TZM alloy mold material, characterized in that, Includes the following steps: S100. Molybdenum powder, titanium hydride powder, zirconium hydride powder and carbon powder are mixed and then subjected to ball milling, powder mixing, cold isostatic pressing and sintering in a hydrogen atmosphere in sequence to improve the uniformity of powder mixing and the density of sintered blanks to obtain sintered blanks. S200. The sintered billet is subjected to hot isostatic pressing to further improve the material density, promote the diffusion of interfacial elements, improve uniformity, avoid stress concentration caused by pores during subsequent forging, which would lead to billet cracking, and improve subsequent forging efficiency and stability. S300. The billet after hot isostatic pressing is subjected to multiple upsetting forgings. The deformation amount shows a single-peak change with an initial increase followed by a decrease along the forging sequence. The deformation is low at both ends and high in the middle section. The direction is unidirectional forging along the axial direction, which promotes the formation of a high-density substructure inside the material and produces a composite strengthening effect together with the dispersed second phase. After forging, S400 is subjected to stress-relief annealing to obtain a large-size high-temperature and high-strength TZM alloy mold.

2. The forming method according to claim 1, characterized in that, In step S100, the ball mill uses a three-dimensional mixer with grinding balls of 3mm and 10mm diameter used together, and the ball milling time is 6-8 hours; the powder mixing time is 3-4 hours.

3. The forming method according to claim 1, characterized in that, In step S100, the pressure of the cold isostatic pressing is 180~250MPa, and the holding time is 1~2h; the temperature of the hydrogen atmosphere sintering is 2000~2100℃, and the sintering time is 4~8h.

4. The forming method according to claim 1, characterized in that, In step S200, the temperature of the hot isostatic pressing is 1200~1300℃, the pressure is 150~170MPa, the time is 2~4h, and the protective atmosphere is argon.

5. The forming method according to claim 1, characterized in that, In step S300, the multi-fire upsetting forging is performed three times, with the material deformation amount in each fire being 10-15%, 25-30%, and 20-25%, and the total deformation amount being controlled at 60%-70%. The initial forging temperature for each fire is 1450-1500℃, and the final forging temperature is 1300-1350℃. After each fire forging, the material is returned to the furnace for reheating. Both the preheating furnace and the reheating furnace are hydrogen-protected heat treatment furnaces.

6. The forming method according to claim 1, characterized in that, In step S400, the stress-relief annealing temperature is 1100~1200℃, the holding time is 5~10h, and the annealing atmosphere is hydrogen protection.

7. The application of the method according to any one of claims 2-6 in the preparation of large-size vacuum isothermal forging dies.

8. A high-temperature, high-strength TZM alloy mold material prepared according to any one of claims 1-6, characterized in that, The material contains both a substructure and a second phase, and the substructure and the second phase synergistically enhance the properties of the material. The material has a tensile strength of 535 MPa, a yield strength of 522 MPa, and an elongation of 12% at 1200℃. After 40 cycles, the strength retention rate exceeds 96%, which meets the requirements of vacuum isothermal forging of the mold at 1000-1200℃.

9. The high-temperature, high-strength TZM alloy mold material according to claim 8, characterized in that, The material comprises, by mass percentage, 0.020%~0.030% C, 0.50%~0.55% Ti, 0.08%~0.11% Zr, with the balance being Mo and unavoidable impurities.

10. The high-temperature, high-strength TZM alloy mold material according to claim 9, characterized in that, The material contains 0.015% to 0.020% O, <0.0015% Al, <0.004% Fe, and <0.0005% Ni.

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