CN117660809A - high-B low-P nickel-based alloy and preparation method thereof - Google Patents
high-B low-P nickel-based alloy and preparation method thereof Download PDFInfo
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- 238000000034 method Methods 0.000 claims description 20
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Abstract
The invention discloses a high-B low-P nickel-based alloy and a preparation method thereof, belongs to the technical field of high-temperature alloys, and solves the problem that the nickel-based high-temperature alloy in the prior art is difficult to simultaneously meet the comprehensive requirements of components on long service life of the alloy and low crack growth rate. The components of the high-B low-P nickel-base alloy comprise the following components in percentage by mass: c:0.02% -0.06%, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.90 percent of Al:1.3 to 1.6 percent of Ti:2.80 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.25 to 2.38, (Al+Ti): 4.35 to 4.58 percent of Fe:0.05% -2.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance. The high-B low-P nickel-base alloy has high strength, good low-cycle fatigue property, low crack growth rate and excellent comprehensive performance.
Description
Technical Field
The invention relates to the technical field of high-temperature alloy, in particular to a high-B low-P nickel-based alloy and a preparation method thereof.
Background
The GH4738 alloy in China is mainly applied to the ground gas turbine and the blades at present, and is gradually applied to aerospace annular members and disc forgings along with the update of smelting technology and domestic equipment. However, because of different service conditions, the performance requirements of the materials are emphasized. The turbine disk and the blades of the gas turbine have long-term service requirements, and the high-temperature durability of the materials is required; aircraft engine turbine disks require materials that take into account both high temperature durability and creep fatigue. Different service requirements of the material ensure a certain lower crack expansion rate, and have higher damage tolerance limit so as to ensure the service reliability of the material; however, the existing GH4738 alloy has a relatively high crack growth rate, and cannot meet the comprehensive requirements of some aerospace components on high alloy durability and low crack growth rate.
Disclosure of Invention
In view of the above, the present invention aims to provide a high-B low-P nickel-based alloy and a method for preparing the same, which are used for solving the problem that the existing nickel-based superalloy is difficult to simultaneously meet the comprehensive requirements of components on high alloy durability and low crack growth rate.
The aim of the invention is mainly realized by the following technical scheme:
in one aspect, the invention provides a high-B low-P nickel-base alloy, which comprises the following components in percentage by mass: c:0.02% -0.06%, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.90 percent of Al:1.3 to 1.6 percent of Ti:2.80 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.25 to 2.38, al+Ti:4.35 to 4.58 percent of Fe:0.05% -2.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
Further, the components of the high-B low-P nickel-base alloy comprise the following components in percentage by mass: c:0.04 to 0.06 percent, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.8 percent of Al:1.3 to 1.5 percent of Ti:3.0 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.3 to 2.38, al+Ti:4.4 to 4.58 percent of Fe:0.5% -1.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
Further, of the components of the high-B low-P nickel-base alloy, ti/Al:2.3 to 2.38.
Further, among the components of the high-B low-P nickel-base alloy, al+ti:4.4 to 4.58 percent.
Further, the microstructure of the high-B low-P nickel-base alloy mainly comprises equiaxed austenite grains, uniformly dispersed boride and carbide, and gamma' phase dispersed in the grains.
Further, in the microstructure of the high-B low-P nickel-based alloy, the carbide mainly comprises M 23 C 6 And MC.
Further, the boride content in the microstructure of the high-B low-P nickel-based alloy is 0.35-0.5%.
Further, in the microstructure of the high-B low-P nickel-base alloy, the size of the gamma' -phase is about 30-50 nm.
The invention also provides a preparation method of the high-B low-P nickel-based alloy, which comprises the following steps:
step 1: smelting to obtain an ingot;
step 2: homogenizing heat treatment is carried out on the cast ingot;
step 3: forging to prepare a bar blank and a forging;
step 4: solution treatment;
step 5: stabilizing and performing time-efficient treatment to obtain the high-B low-P nickel-based alloy.
Further, in step 2, the process steps of homogenizing heat treatment include:
s201, heating to 1165-1170 ℃, and preserving heat for 35-37 hours;
s202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, and discharging and air cooling.
Compared with the prior art, the invention can at least realize one of the following beneficial effects:
a) According to the high-B low-P nickel-based alloy, the solid solution strengthening effect of the alloy and the grain boundary strength of the alloy are improved by accurately controlling the content of B, zr, fe, C, cr, co, al, ti and other single elements in the alloy; and by cooperatively controlling the values of B, zr and C, the best matching of the morphology and distribution of carbides in the alloy and the contents and the sizes of other second phases is ensured, the risk crack sources in the alloy are reduced on the premise of not reducing the strength of the alloy, the carbides are uniformly distributed in the alloy grain boundaries, the strength of the grain boundaries is improved, the dislocation slip speed is delayed, the purposes of improving the strength of the alloy, reducing the cracking tendency and the crack expansion rate are achieved, and the comprehensive performance of the alloy is ensured.
b) The preparation method of the high-B low-P nickel-base alloy avoids deformation by precisely controlling the technological parameters of each step, ensures that crystal grains of the alloy are uniform and fine, and ensures that the morphology and distribution of carbides in the alloy and other second phase content and sizes are optimally matched, thereby achieving the purposes of improving the alloy strength, reducing the cracking tendency and the crack expansion rate and ensuring the comprehensive performance of the alloy.
c) The performance of the high-B low-P nickel-base alloy of the invention is as follows: room temperature performance: tensile strength sigma b More than or equal to 1300 MPa (e.g. 1330-1360 MPa); yield strength sigma 0.2 Not less than 1050 MPa (for example, 1052-1080 MPa); elongation after break delta 5 More than or equal to 27 percent (for example, 27 percent to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (for example, 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 h (for example, 72-88 h); elongation after break delta 5 More than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83-95 hours); elongation after break delta 5 More than or equal to 26 percent (such as 26 percent to 30 percent); low cycle fatigue performance: 500 ℃/strain control 0-0.7%/0.33 Hz > 5 x 10 4 Week (e.g., 57532-63645 weeks). The high-B low-P nickel-base alloy has high strength, good low-cycle fatigue property, low crack growth rate and excellent comprehensive performance.
Additional features and advantages of the invention will be set forth in the description which follows, and in part will be obvious from the description, or may be learned by practice of the invention. The objectives and other advantages of the invention will be realized and attained by the structure particularly pointed out in the written description and claims thereof as well as the appended drawings.
Drawings
The drawings are only for purposes of illustrating particular embodiments and are not to be construed as limiting the invention, like reference numerals being used to designate like parts throughout the drawings;
FIG. 1 is a graph of crack growth rates for example 1 and comparative example 4;
FIG. 2a is one of the microstructure diagrams of example 1 of the present invention;
FIG. 2B is a B-element electron probe diagram of example 1 of the present invention;
FIG. 2C is a diagram of an electron probe for the C element according to example 1 of the present invention;
FIG. 2d is a diagram of an electron probe for Ti element according to example 1 of the present invention;
FIG. 2e is a diagram of an Al element electron probe according to example 1 of the present invention;
FIG. 2f is a diagram of an electron probe for Cr element according to example 1 of the present invention;
FIG. 2g is a Co element electron probe diagram of example 1 of the present invention;
FIG. 2h is a diagram of an electron probe of Mo element in example 1 of this invention;
FIG. 3 is a second microstructure chart of example 1 of the present invention;
FIG. 4 is a microstructure of comparative example 4;
fig. 5 is a grain diagram of comparative example 4.
Detailed Description
Preferred embodiments of the present invention are described in detail below with reference to the attached drawing figures, which form a part of the present invention and are used in conjunction with embodiments of the present invention to illustrate the principles of the present invention.
The invention provides a high-B low-P nickel-base alloy, which comprises the following components in percentage by mass: c:0.02% -0.06%, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.90 percent of Al:1.3 to 1.6 percent of Ti:2.80 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.25 to 2.38, al+Ti:4.35 to 4.58 percent of Fe:0.05% -2.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
Specifically, O+N+S is less than or equal to 50PPm, for example O+N+S is less than or equal to 40PPm.
The following is a specific description of the action and the selection of the amounts of the components contained in the invention:
c: the C content is increased, and carbide (MC, M) 23 C 6 ) Will also increase, but M 23 C 6 The precipitation temperature of (2) does not vary greatly. In the invention, the beneficial phase M is separated out at the grain boundary by controlling the content of C and matching with B element 23 C 6 The alloy is continuously or necklace-shaped, film-shaped precipitation is avoided, the film-shaped continuous precipitation is harmful to alloy performance, and the continuous or necklace-shaped precipitation is beneficial to alloy performance; by properly reducing the content of C and simultaneously controlling the content of alloy impurity elements, the contents of primary carbide and nitride in the crystal are effectively reduced, the source of fatigue cracking cracks is reduced, and the fatigue performance of the alloy is improved. In general, when the C content is more than 0.06%, the alloy life is reduced, but the grain boundary segregation of B element causes grain boundary membranous M by increasing the content of B, zr and Mg element 23 C 6 Is not easy to separate out, and improves grain boundary. Therefore, the content of C is controlled to be 0.02-0.06%.
Cr: the Cr element can improve the oxidation resistance and corrosion resistance of the alloy. At the same time Cr as M 23 C 6 Is a main forming element of (2) in the content corresponding to M 23 C 6 The content of (2) has a smaller influence but a significant influence on M 23 C 6 The more Cr element content, M 23 C 6 When the Cr content exceeds 20%, M 23 C 6 The precipitation temperature of (2) exceeds 960 ℃. Therefore, the Cr content is controlled to be 18.5-20.0 percent in the invention.
Co: the Co element mainly plays a solid solution strengthening role, and the addition of Co to the gamma matrix can reduce the stacking fault energy of the matrix, reduce the stacking fault energy, increase the occurrence probability of the stacking fault and make the intersection and sliding of dislocation more difficult, so that the deformation requires larger external force and is shown as the improvement of the strength; and the stacking fault energy is reduced, the creep rate is reduced, and the creep resistance is increased. In addition, the Co element can also reduce the solubility of the gamma prime forming elements Ti and Al in the matrix, so that the quantity of gamma prime precipitated phases in the alloy is increased, and the service temperature of the alloy is increased. However, co belongs to a scarce resource in China, and the cost factor of the alloy is considered, so that the content of Co element is controlled to be 13.0-14.0%.
Mo: mo enters the nickel-based alloy matrix and mainly plays a role in solid solution strengthening. Meanwhile, the lattice constant of the matrix is increased, the gamma-gamma' mismatching degree is reduced, and the alloy stability and high-temperature creep property are improved. However, the increase in Mo element lowers the initial and final melting temperatures of the alloy. Therefore, the content of Mo element is controlled to be 4.0-4.9%.
Al: the Al element is the main forming element of the gamma' phase, and about 20% of Al added into the alloy enters a gamma solid solution to play a role in solid solution strengthening; 80% of Al forms Ni with Ni 3 Al plays a role in precipitation strengthening. In order to ensure that the alloy has the necessary gamma prime phase to maintain the high temperature strength at 700 c. The invention controls the Al content to be 1.3-1.6%.
Ti: about 10% of the Ti added into the nickel-based alloy enters a gamma solid solution to play a certain solid solution strengthening role, and about 90% enters a gamma' phase. Under the condition of a certain Al content, as the quantity of the gamma 'phase is increased and the high-temperature strength of the alloy is increased along with the increase of the Ti content, in order to ensure that the gamma' phase which is necessary for maintaining the high-temperature strength of 700 ℃ exists in the alloy, the invention controls the Ti content to be 2.80-3.25 percent and simultaneously controls the Ti/Al:2.25 to 2.38, al+Ti:4.35 to 4.58 percent.
Increasing the content of Al or Ti alone or simultaneously does not necessarily increase the content of the secondary γ' phase at the corresponding solid solution temperature. However, the high (Al+Ti) content has obvious advantages that the primary gamma 'phase is completely dissolved in the low (Al+Ti) group due to the improvement of the total dissolution temperature of the primary gamma' phase, the secondary gamma 'phase of the high (Al+Ti) group has more content in the temperature range where the primary gamma' phase is not completely dissolved in the high (Al+Ti) group, and the reasonable temperature control can prevent the grain size from growing due to the pinning effect of the primary gamma 'phase, and meanwhile, the gamma' phases of the two sizes can jointly play a reinforcing role.
B: boride (mainly including M 3 B 2 ，M 4 B 3 ，MB 12 ) The boride is uniformly distributed in the crystal by homogenization treatment and cogging. By increasing the B content, the segregation of B atoms in the grain boundary is increased, the binding force of the grain boundary is increased, and the activation energy of grain boundary diffusion is increased, so that beneficial continuous or necklace-shaped M is formed in the grain boundary by utilizing the thermodynamic transformation 23 C 6 Carbide can increase the surface energy of grain boundary cracks while preventing grain boundary sliding, improve the grain boundary strength of alloy and improve the heightThe temperature lasting creep property has obvious effect on improving creep limit, especially the lasting strength of the material, and the higher the content of B, the better the high temperature lasting property. However, when the B content is higher, the formation of a more multilayered low-melting phase is unfavorable for the improvement of the alloy properties but leads to the decrease of the properties, so that the B content is not excessively high. Comprehensively considering the invention, the content of B is controlled to be 0.015% -0.03%.
Fe: the lattice constant of Fe element is 3% different from that of Ni element, and its lattice expansion causes long-range stress field to block dislocation movement; meanwhile, fe can also reduce stacking fault energy, improve yield strength and play a solid solution strengthening role. The majority of elements of the high-B low-P nickel-base alloy are grain boundary strengthening elements, and when the grain boundary strengthening elements are added, a part of Fe elements are added, so that the strength of the base is improved and matched; thus, in the present invention, fe:0.05% -2.0%.
P: according to the invention, the beneficial precipitated phase of the grain boundary P element is obtained by reasonably controlling the content of the P element, so that the high-temperature durable creep property of the alloy is improved, and meanwhile, the plasticity of the alloy is not damaged; but high levels of P have significant notch sensitivity for the high temperature durability of some alloys. In the invention, a proper amount of B, zr elements are added, so that the control of P is required in the invention due to comprehensive consideration: 0.004% -0.009%.
Zr: zr element is used as a crystal boundary strengthening element, which can effectively reduce the solidus temperature of the nickel-based alloy so as to enlarge the solidification zone of the alloy, and simultaneously promote C, S element to gather in the crystal boundary together with Mg element, purify the alloy, and the addition of proper amount of Zr element can improve the creep resistance of the alloy, but the excessive content of Zr element can form more Ni 5 Zr, etc., and to the alloy Jin Suxing. The invention improves the content of the element B and the content of the element Zr, and improves the strength of the material through the segregation of the Zr at the interface of the eutectic structure and the matrix and the grain boundary. For the problem of material plasticity reduction, the eutectic phase is eliminated by homogenization treatment regulation and control, and the eutectic phase is prevented from becoming a crack source. Considering comprehensively, zr needs to be controlled in the invention: 0.05% -0.15%.
Mg: the Mg element has high affinity with S and is easy to form sulfide, so that the harmful risk of S is reduced; in addition, mg element is strongly segregated in the grain boundary and the phase boundary, and the carbide is thinned through the action of the grain boundary and the phase boundary, so that the grain boundary energy and the phase boundary energy are reduced, the interface binding force is improved, the creep strength and the plasticity of the alloy are improved, and the crack nucleation and the crack expansion are difficult to carry out. In the invention, the Mg content is controlled to be less than or equal to 0.005 percent by comprehensive consideration.
O: the reduction of the O, N element content can reduce the inclusion quantity in the material, and is beneficial to improving the plasticity and toughness of the alloy. During melting of the alloy, N readily forms Ti (C, N) with Ti, and the increase in Ti (C, N) increases the likelihood of fatigue source formation. Therefore, in the present invention, it is necessary to control O: less than or equal to 20PPm, N: and the volume of the mixture is less than or equal to 20PPm.
S: higher S content has an effect on the plasticity and long-term properties of the alloy. The higher the S content at the end of alloy smelting solidification, the easier the sulfide is to precipitate. The S element has obvious influence on the nickel-based alloy at the temperature of more than 800 ℃, the S element is obvious in the steel ingot smelting and cogging process, and the inventor finds that in the research: 100 The S-content in ppm fails to smelt, while the S-content in 56 ppm group steel ingot has serious cracking in the cogging process, and after the S-content exceeds 10ppm, the durability and plasticity at 730 ℃ are reduced to different degrees. Therefore, in the present invention, it is necessary to control S: less than or equal to 10PPm.
In order to further improve the comprehensive performance of the high-B low-P nickel-base alloy, the components of the high-B low-P nickel-base alloy comprise the following components in percentage by mass: c:0.04 to 0.06 percent, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.8 percent of Al:1.3 to 1.5 percent of Ti:3.0 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.3 to 2.38, (Al+Ti): 4.4 to 4.58 percent of Fe:0.5% -1.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
Specifically, the microstructure of the high-B low-P nickel-based alloy mainly comprises equiaxed austenite grains, boride and carbide which are uniformly dispersed and distributed, and gamma' phase which is dispersed and distributed in the grains; wherein the carbide mainly comprises M 23 C 6 And MC; wherein M is 23 C 6 The content is 0.85% -1.17%, the MC content is 0.11% -0.33%, and the boride content is 0.35% -0.5%; gamma' phase rulerSmall, the size of the gamma' phase is about 30-50 nm; the MC carbide has a size of about 5-10 μm and M 23 C 6 The size is about 0.3 to 1 μm.
In particular, M in the microstructure of the high-B low-P nickel-base alloy 23 C 6 Is distributed in a continuous shape or necklace shape and short bars shape, and is not precipitated in a film shape; MC is often in the form of a block with rounded boundaries.
Specifically, the crystal grains of the high-B low-P nickel-base alloy are uniform and fine, and the average grain size can reach about 8-9 levels uniformly.
The invention also provides a preparation method of the high-B low-P nickel-based alloy, which comprises the following steps:
step 1: smelting to obtain an ingot;
step 2: homogenizing heat treatment is carried out on the cast ingot;
step 3: forging to prepare a bar blank and a forging;
step 4: solution treatment;
step 5: stabilizing and performing time-efficient treatment to obtain the high-B low-P nickel-based alloy.
Specifically, in the step 1, the ingot is formed by adopting VIM+ESR+VAR triple smelting or VIM+VAR duplex smelting.
Specifically, in the step 2, the purpose of the homogenization heat treatment is to eliminate the low melting point phase and eutectic phase in the ingot and to reduce element segregation. The homogenization heat treatment is provided with 2 sections of heat preservation, and in the 1 st section of heat preservation, the temperature is lower, and the main effect is to eliminate a low-melting-point phase and a eutectic phase in the alloy; the heat preservation in the 2 nd stage can promote the uniform diffusion of segregation elements. Specifically, the homogenizing heat treatment process comprises the following steps:
s201, heating to 1165-1170 ℃, and preserving heat for 35-37 hours;
s202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, and discharging and air cooling.
Specifically, in S201, the temperature is slowly increased from the furnace temperature of less than or equal to 400 ℃ to 1165-1170 ℃ for 10-15 h.
Specifically, in S201, considering that the low-melting-point phase and the eutectic phase formed by the B element in the as-cast structure are remelted due to the excessively high temperature and the excessively long heat preservation time, on one hand, holes are formed after homogenization to affect the subsequent alloy cogging and even cause ingot cracking to be scrapped, on the other hand, the B element cannot be fully diffused into the matrix to prepare for the subsequent formation of the grain boundary strengthening phase, and finally the grain boundary cannot be strengthened to improve the lasting creep property of the alloy. Therefore, the temperature is controlled to be raised to 1165-1170 ℃ and kept for 35-37 hours.
Specifically, in S202, the higher the temperature and the longer the heat preservation time, the more fully the remelting of the segregation element, but when the temperature and the time reach the matching balance point, the remelting of the segregation element will remain stable, so that the temperature is too high and the heat preservation time is too long, the positive effect on the remelting of the segregation element is smaller, and on the contrary, coarse grains, energy waste and production efficiency are caused; and the low temperature and the short heat preservation time can not ensure that most of segregation elements are remelted, and dendrite segregation reduces the hot working plasticity in the forging process. Therefore, the temperature is controlled to be raised to 1200-1210 ℃, and the temperature is kept for 38-42 h.
Specifically, in S202, considering that more grain boundary strengthening elements and carbide forming elements are added into the alloy of the invention, the strength of the alloy is improved, and meanwhile, the second phase is increased, and the risk of cracking of the alloy in the ingot casting and cogging process is increased, so that furnace cooling is controlled to 1000±10 ℃ and then furnace discharging and air cooling are performed, the furnace cooling treatment can enable the bent grain boundary to have more proportion and better high-temperature plasticity, so that the furnace cooling is controlled to be increased by controlling homogenization treatment, the toughness of the ingot casting in the cogging process is improved, and cracking is avoided.
Specifically, in the step 3, the forging method includes, but is not limited to, free forging or upsetting-drawing+radial forging combined cogging and subsequent hot rolling and cold drawing, wherein the fast forging method includes, but is not limited to, a combined method of single drawing and upsetting-drawing. The bar specification of cogging forging or subsequent rolling and cold drawing can reach the range of phi 5-phi 600 mm.
Specifically, in the step 3, for the single drawing, the drawing deformation amount is 30-60% per time, and the target size can be achieved through three to four times of drawing; for upsetting and drawing, the upsetting deformation is 30-60% per time, and the drawing deformation is 30-50% per time. The temperature range for preparing the bar is 1060-1200 ℃. The temperature of each firing is reduced by 40-80 ℃ until the temperature of the last firing is lower than 1080 ℃, wherein upsetting and drawing are carried out in a group, and then the temperature is reduced.
Specifically, in the step 3, the rod blank may be subjected to die forging to prepare a turbine disc with a disc shaft integrally or to prepare an annular member by ring rolling.
Specifically, in the step 4, the solution treatment process is as follows: heating to 995-1020 ℃, preserving heat for 3-5 h, discharging and cooling.
Specifically, in the step 4, the heat preservation time is insufficient because the alloy core does not reach the solid solution temperature due to the excessively high heating speed; therefore, the temperature is slowly raised, and the temperature raising speed is controlled to be 3-6 ℃/min. Too high a temperature or too long a holding time can result in grain growth that is detrimental to alloy performance. The temperature is low or the heat preservation time is too short, and the preparation for obtaining a proper gamma' phase in the subsequent aging process cannot be made; therefore, the temperature is controlled to be increased to 995-1020 ℃ and the heat is preserved for 3-5 hours.
Specifically, in the step 4, in order to prevent the excessive stress caused by the too fast cooling speed from deforming the parts, different cooling modes are adopted for the parts with different wall thicknesses. For example, for disc shaft parts with the wall thickness of 20-200 mm, discharging the disc shaft parts into cooling oil for cooling; and (3) discharging air cooling is adopted for the annular piece with the wall thickness of 10-100 mm.
Specifically, in the step 5, the stabilizing and time-efficient treatment process comprises: heating to 820-850 ℃, preserving heat for 5-7 h, air-cooling to 500+/-10 ℃, continuously heating to 740-770 ℃, preserving heat for 8-10 h, and air-cooling to room temperature.
Specifically, in the step 5, considering that the temperature rising speed is too high, the alloy core does not reach the stabilization and aging treatment temperature, the heat preservation time is insufficient, the precipitated gamma' phase is insufficient, and the quantity and the size do not reach the optimal matching value; too high a temperature or too long a holding time can lead to larger gamma' -phase size, deviate from the optimal size-to-number ratio and affect the alloy strength. Therefore, the temperature is slowly increased, the temperature is increased from the temperature of the furnace to be less than or equal to 400 ℃, and the temperature increasing speed is controlled to be 2-3.5 ℃/min.
Specifically, in the step 5, the temperature is kept at 820-850 ℃ for 5-7 h, and the air cooling is carried out to 500+/-10 ℃ instead of the air cooling to room temperature, so that the efficiency can be improvedThe ratio and the size of carbide can be effectively ensured. For example, MC carbide has a size of 5-10 μm; m is M 23 C 6 The size of the carbide is about 0.3-1 μm.
Specifically, the performance of the high-B low-P nickel-base alloy obtained in the above step 5 is as follows: room temperature performance: tensile strength sigma b More than or equal to 1300 MPa (e.g. 1330-1360 MPa); yield strength sigma 0.2 Not less than 1050 MPa (for example, 1052-1080 MPa); elongation after break delta 5 More than or equal to 27 percent (for example, 27 percent to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (for example, 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 h (for example, 72-88 h); elongation after break delta 5 More than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83-95 hours); elongation after break delta 5 More than or equal to 26 percent (such as 26 percent to 30 percent); low cycle fatigue performance: 500 ℃/strain control 0-0.7%/0.33 Hz > 5 x 10 4 Week (e.g., 57532-63645 weeks).
The invention also provides application of the high-B low-P nickel-base alloy, and the high-B low-P nickel-base alloy can be used for rotating parts of aeroengines in service at the temperature of more than 700 ℃. Such as for alloy disk shaft integrated turbine disks.
Compared with the prior art, the high-B low-P nickel-based alloy improves the solid solution strengthening effect of the alloy and the grain boundary strength of the alloy by accurately controlling the content of B, zr, fe, C, cr, co, al, ti and other single elements in the alloy; and by cooperatively controlling the values of B, zr and C, the best matching of the morphology and distribution of carbides in the alloy and the contents and the sizes of other second phases is ensured, the risk crack sources in the alloy are reduced on the premise of not reducing the strength of the alloy, the carbides are uniformly distributed in the alloy grain boundaries, the strength of the grain boundaries is improved, the dislocation slip speed is delayed, the purposes of improving the strength of the alloy, reducing the cracking tendency and the crack expansion rate are achieved, and the comprehensive performance of the alloy is ensured.
The preparation method of the high-B low-P nickel-base alloy avoids deformation by precisely controlling the technological parameters of each step, ensures that crystal grains of the alloy are uniform and fine, and ensures that the morphology and distribution of carbides in the alloy and other second phase content and sizes are optimally matched, thereby achieving the purposes of improving the alloy strength, reducing the cracking tendency and the crack expansion rate and ensuring the comprehensive performance of the alloy.
The performance of the high-B low-P nickel-base alloy of the invention is as follows: room temperature performance: tensile strength sigma b More than or equal to 1300 MPa (e.g. 1330-1360 MPa); yield strength sigma 0.2 Not less than 1050 MPa (for example, 1052-1080 MPa); elongation after break delta 5 More than or equal to 27 percent (for example, 27 percent to 28 percent); the area reduction ratio psi is more than or equal to 35 percent (for example, 35 to 40 percent); 730 ℃/550MPa durability performance: the lasting time tau is more than or equal to 70 h (for example, 72-88 h); elongation after break delta 5 More than or equal to 17.5 percent (for example, 17.7 to 20 percent); 815 ℃/295MPa durability: the lasting time tau is more than or equal to 83 hours (for example, 83-95 hours); elongation after break delta 5 More than or equal to 26 percent (such as 26 percent to 30 percent); low cycle fatigue performance: 500 ℃/strain control 0-0.7%/0.33 Hz > 5 x 10 4 Week (e.g., 57532-63645 weeks).
Examples 1 to 5
The following specific examples and comparative examples demonstrate the advantages of the present invention for precise control of the composition and process parameters of the high B low P nickel-base alloys. Embodiments 1-5 of the present invention provide a high B low P nickel-based alloy and a method of making the same.
The compositions of the high B low P nickel-base alloys of examples 1-5 are shown in Table 1 below.
The method for preparing the high-B low-P nickel-base alloys of examples 1-5 comprises:
example 1
Step 1: sequentially carrying out vacuum induction melting, electroslag remelting and vacuum arc remelting to obtain an ingot;
step 2: homogenizing and annealing the cast ingot: heating from 400 ℃ to 1165 ℃ through 11h, and preserving heat for 37h; continuously heating to 1200 ℃, preserving heat for 38 hours, cooling to 1000+/-10 ℃, discharging and air cooling;
step 3: preparing a bar blank by adopting quick forging and radial forging: sequentially upsetting and drawing out the alloy rod blank with the temperature of 1200 ℃ for three times, and radially forging; wherein, the deformation amount of upsetting each time is 30 percent, and the deformation amount of drawing each time is 30 percent; cooling to 40 ℃ after upsetting and drawing out each time, and cogging on a radial forging machine for multiple times until the heat preservation temperature is reduced to 1080 ℃ to obtain a finished bar with phi 180 mm; forging the bar blank by adopting a die forging press to prepare a forging piece: orderly upsetting cakes and die forging the alloy rod blank section with the temperature kept at 1060 ℃; the deformation amounts of the upsetting cakes and the die forging are respectively 50% and 35%, so that a forging piece is obtained;
step 4: solution treatment: heating from the furnace temperature of 400 ℃ at a heating speed of 5 ℃/min to 1015 ℃, preserving heat for 3.5h, and cooling with oil;
step 5: stabilizing and performing time-efficient treatment to obtain a nickel-based superalloy disc member: heating from the furnace temperature of 400 ℃ at a heating speed of 5 ℃/min, heating to 850 ℃ and preserving heat for 5h, air-cooling to 500+/-10 ℃, continuously heating to 760 ℃ and preserving heat for 8h, and air-cooling to room temperature to obtain the nickel-based superalloy disc.
Example 2
The preparation method of this example is substantially the same as that of example 1, except that:
step 1: sequentially carrying out vacuum induction melting and vacuum arc remelting to obtain an ingot;
step 3: upsetting, punching and ring rolling alloy rod blank sections with the temperature kept at 1060 ℃ in sequence; upsetting and ring rolling deformation are respectively 70% and 40%, so that a forging is obtained;
step 4: cooling the forge piece after solid solution by adopting air cooling;
the remaining steps are the same.
Example 3
The preparation method of this example is substantially the same as that of example 1, except that:
step 1: sequentially carrying out vacuum induction melting, electroslag remelting and vacuum arc remelting to obtain an ingot;
step 3: preparing a bar blank by adopting a quick forging mode: sequentially upsetting and drawing out the alloy rod blank with the temperature of 1180 ℃ for three times, and radially forging; wherein, the deformation amount of upsetting each time is 50 percent, and the deformation amount of drawing each time is 50 percent; cooling to 40 ℃ after upsetting and drawing out each time until the heat preservation temperature is reduced to 1060 ℃ and then cogging to obtain a finished bar with phi 240 mm; forging the bar blank by adopting a die forging press to prepare a forging piece: orderly upsetting cakes and die forging the alloy rod blank section with the temperature kept at 1060 ℃; the deformation amounts of the upsetting cakes and the die forging are respectively 50% and 35%, so that a forging piece is obtained;
the remaining steps are the same.
Examples 4-5 differ from example 3 in the chemical composition of the alloy, see Table 1, with the remaining parameters being the same.
The invention also provides 6 comparative examples, examples 1-5 and comparative examples 1-5 steel with chemical compositions shown in Table 1. The metallographic structures of the examples and comparative examples are shown in Table 2.
Comparative examples 1 to 4 were prepared in the same manner as in example 1, except that the components were different.
Comparative example 5 has the same components as in example 1, and the homogenization heat treatment mode in the preparation method is different, wherein the homogenization method in comparative example 5 is to heat up to 1160 ℃ from 400 ℃ for 13 hours, and heat preservation is carried out for 48 hours; continuously heating to 1190 ℃, preserving heat for 66h, and then air cooling.
Comparative example 6 has the same components as example 1, and the homogenization heat treatment mode in the preparation method is different, wherein the homogenization method in comparative example 6 is that the homogenization is carried out after heat preservation for 38 hours, and then the homogenization is carried out in an air cooling furnace; cracks appear during the subsequent forging process.
Table 3 shows the room temperature mechanical properties of the examples and comparative examples of the present invention, table 4 shows the durability properties of the examples and comparative examples, and Table 5 shows the low cycle fatigue properties of the examples and comparative examples.
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FIG. 1 is a graph of crack growth rates for example 1 and comparative example 4, from which it can be seen that the present method has smaller da/dN values (fatigue crack growth rates) at the same ΔK (stress intensity factor), and it can be seen that the high-B low-P nickel-base alloy of the present invention has a low crack growth rate.
FIGS. 2a and 3 are each a microstructure of example 1, in which the blocks MC with rounded boundaries are distributed in the crystal as seen in FIG. 2a, and in which the short bars M are seen in FIG. 3 23 C 6 Is continuously and uniformly distributed on the grain boundary. FIGS. 2b to 2h are electron probe diagrams of example 1; FIG. 4 is a microstructure of comparative example 4 in which M films are distributed on grain boundaries 23 C 6 Fig. 5 is a grain diagram of comparative example 4, and it can be seen that the grain structure of comparative example 4 is a mixed crystal.
The present invention is not limited to the above-mentioned embodiments, and any changes or substitutions that can be easily understood by those skilled in the art within the technical scope of the present invention are intended to be included in the scope of the present invention.
Claims (10)
1. A high-B low-P nickel-base alloy, characterized in that the components of the high-B low-P nickel-base alloy comprise, in mass percent: c:0.02% -0.06%, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.90 percent of Al:1.3 to 1.6 percent of Ti:2.80 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.25 to 2.38, al+Ti:4.35 to 4.58 percent of Fe:0.05% -2.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
2. The high-B low-P nickel-base alloy according to claim 1, wherein the components of the high-B low-P nickel-base alloy comprise, in mass percent: c:0.04 to 0.06 percent, cr:18.5 to 20.0 percent, co:13.0 to 14.0 percent, mo:4.0 to 4.8 percent of Al:1.3 to 1.5 percent of Ti:3.0 to 3.25 percent, B:0.015% -0.03%, zr:0.05% -0.15%, ti/Al:2.3 to 2.38, al+Ti:4.4 to 4.58 percent of Fe:0.5% -1.0%, mg is less than or equal to 0.005%, P:0.004% -0.009%, O: less than or equal to 20PPm, N: less than or equal to 20PPm, S less than or equal to 10PPm, nickel: the balance.
3. The high B low P nickel-base alloy according to claim 1, wherein the high B low P nickel-base alloy comprises, in its composition, ti/Al:2.3 to 2.38.
4. The high-B low-P nickel-base alloy according to claim 1, wherein, among the components of the high-B low-P nickel-base alloy, al+ti:4.4 to 4.58 percent.
5. The high B low P nickel-base alloy according to claim 1, wherein the microstructure of the high B low P nickel-base alloy consists essentially of equiaxed austenite grains and uniformly dispersed borides and carbides, and a gamma prime phase dispersed within the grains.
6. The high-B low-P nickel-base alloy according to claim 5, wherein in the microstructure of said high-B low-P nickel-base alloy, carbides mainly comprise M 23 C 6 And MC.
7. The high-B low-P nickel-base alloy according to claim 5, wherein the boride content in the microstructure of the high-B low-P nickel-base alloy is 0.35-0.5%.
8. The high B low P nickel-base alloy according to claim 5, wherein the gamma prime phase has a size of about 30-50 nm in the microstructure of the high B low P nickel-base alloy.
9. A method of making the high B low P nickel-base alloy of any of claims 1-8, comprising:
step 1: smelting to obtain an ingot;
step 2: homogenizing heat treatment is carried out on the cast ingot;
step 3: forging to prepare a bar blank and a forging;
step 4: solution treatment;
step 5: stabilizing and performing time-efficient treatment to obtain the high-B low-P nickel-based alloy.
10. The method according to claim 9, wherein in the step 2, the homogenizing heat treatment comprises the steps of:
s201, heating to 1165-1170 ℃, and preserving heat for 35-37 hours;
s202, continuously heating to 1200-1210 ℃, preserving heat for 38-42 h, cooling to below 1000+/-10 ℃, and discharging and air cooling.
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