JP6485692B2 - Heat resistant alloy with excellent high temperature strength, method for producing the same and heat resistant alloy spring - Google Patents

Description

  The present invention relates to a heat-resistant alloy excellent in high-temperature strength, a manufacturing method thereof, and a heat-resistant alloy spring.

Conventionally, heat-resistant springs have been used for opening and closing exhaust valves of engines, and heat-resistant springs made of heat-resistant alloys such as Waspaloy (registered trademark) and Inconel 718 (Inconel: registered trademark) are known.
Inconel 718 (registered trademark) is known as a Ni-based heat-resistant alloy using a γ ″ (gamma double prime) phase as a reinforcing phase, and Waspaloy (registered trademark) is more stable than γ ″ phase. It is known as a heat-resistant alloy in which about 25% by volume is precipitated using a (gamma prime) phase as a strengthening phase. Udimet 720 (registered trademark) with improved heat resistance is known as a Ni-based heat-resistant alloy in which about 45% by volume of the γ ′ phase is precipitated and tungsten is added for solid solution strengthening of the γ phase.

  Conventionally, as Ni-based alloys having excellent mechanical strength for steam turbines and gas turbines of power plants, C: 0.01 to 0.15%, Cr: 14 to 20%, Co: 10 to 15%, Mo: 6 to 12%, Al: 0.5 to 4%, Ti: 0.5 to 4%, B: 0.001 to 0.006%, from the remaining Ni and inevitable impurities There are known carbides that have a discontinuous portion on at least a part of the crystal grain boundary along the grain boundary and precipitate in a lump shape, and an alloy having a precipitate precipitated in the grain boundary. (See Patent Document 1).

Conventionally, as Ni-base heat-resistant superalloys for heat-resistant members such as engines and gas turbines, Cr, Co, Ti, Al, Ni are included as main elements, annealing twins are included, and the average crystal grain size of the microstructure is 1-60 μm A nickel-based heat-resistant superalloy is known (see Patent Document 2).
Conventionally, 2 to 25% by mass of Cr, 2 to 7% by mass or less of Al, 19.5 to 55% by mass of Co are contained as a basic composition, a specific amount of Ti is added, and a γ ′ solid solution temperature of 93 to A Ni-base heat-resistant alloy that is solid solution at less than 100% is known (see Patent Document 3).

JP 2013-177668 A International Publication No. 2011/138852 International Publication No. 2013/089218

The limit of sag resistance is 650 ° C. for Inconel 718 (registered trademark) and 700 ° C. for Waspalloy (registered trademark).
In recent years, there is a tendency to arrange exhaust valves closer to the engine in order to further increase the efficiency of the engine. That is, this type of heat-resistant alloy is required to further improve heat resistance and mechanical properties. In addition, heat-resistant springs made of this type of heat-resistant alloy are required to have higher temperature and heat resistance than ever before.

  The present invention has been made in view of conventional circumstances, and an object of the present invention is to provide a heat-resistant alloy having a high-temperature strength and sag resistance superior to those of conventional materials, a manufacturing method thereof, and a heat-resistant alloy spring.

  In order to solve the above-described problems, the heat-resistant alloy having excellent high-temperature strength according to the present invention has a mass ratio of Ni: 34 to 46%, Cr: 13.5 to 18%, Mo: 2 to 9%, Nb 2.5 to 3. 5%, Fe: 1-2%, Al: 1.5-2.5%, Ti: 0.6-1.0%, the composition of the balance Co and inevitable impurities, and the structure, A γ parent phase, a γ ′ phase having a particle size of 20 to 50 nm and a volume ratio of 52 to 65%, and a μ phase having a volume ratio of 10% or less are provided.

In the above composition, the heat-resistant alloy having excellent high-temperature strength according to the present invention further has Mn: 0.3 to 0.6%, Si: 0.15 to 0.25%, Zr: 0.04 to 0.06%, C: You may contain 1 type, or 2 or more types of 0.04-0.06%.
In the heat-resistant alloy excellent in high temperature strength of the present invention, the γ ′ phase is represented by a composition formula (Ni, Co) ₃ (Al, Nb, Cr, Co, Ti), and Ni: 57 to 61% by mass ratio. A composition containing Al: 5-6%, Nb: 10-13%, Ti: 3-5%, Cr: 2-3%, and Co 15-17% is preferable.
In the heat-resistant alloy excellent in high temperature strength of the present invention, the Vickers hardness is HV420 or more and less than 560 for the plate material (rolled material), the tensile strength is 1080 MPa or more and less than 1580 MPa, and the wire material (drawn material) is Vickers hardness HV360 or more and 630. Or a tensile strength of 1210 MPa or more and less than 2170 MPa.

The manufacturing method of the heat-resistant alloy excellent in high temperature strength of the present invention is as follows: mass ratio: Ni: 34 to 46%, Cr: 13.5 to 18%, Mo: 2 to 9%, Nb 2.5 to 3.5%, Fe : Ingot containing 1 to 2%, Al: 1.5 to 2.5%, Ti: 0.6 to 1.0%, and having the balance Co and inevitable impurity composition at 1150 to 1250 ° C and 4 to 8 after time homogenization treatment, it was subjected to an aging treatment of 2-4 hours at 700-850 ° C., and molded into a predetermined shape, Ri by the applying 2-4 hours of aging treatment at 700-850 ° C., It is characterized by having a structure including a γ parent phase, a γ ′ phase having a particle size of 20 to 50 nm and a volume ratio of 52 to 65%, and a μ phase having a volume ratio of 10% or less.
The manufacturing method of the heat-resistant alloy excellent in high temperature strength of the present invention is as follows: mass ratio: Ni: 34 to 46%, Cr: 13.5 to 18%, Mo: 2 to 9%, Nb 2.5 to 3.5%, Fe : Ingot containing 1 to 2%, Al: 1.5 to 2.5%, Ti: 0.6 to 1.0%, and having the balance Co and inevitable impurity composition at 1150 to 1250 ° C and 4 to 8 After the time equalization treatment, a cold working with a processing rate of 20% or more and less than 60% is added, formed into a predetermined shape, and subjected to an aging treatment at 700 to 850 ° C. for 1 to 4 hours. And a structure having a γ ′ phase having a particle size of 20 to 50 nm and a volume ratio of 52 to 65% and a μ phase having a volume ratio of 10% or less.

In the method for producing a heat-resistant alloy having excellent high temperature strength according to the present invention, Mn: 0.3-0.6%, Si: 0.15-0.25%, Zr: 0.04-0. An ingot having a composition containing one or more of 06% and C: 0.04 to 0.06% can be used.
In the method for producing a heat-resistant alloy having excellent high-temperature strength according to the present invention, the γ ′ phase is represented by the composition formula (Ni, Co) 3 (Al, Nb, Cr, Co, Ti), and the mass ratio is Ni: 57 to 61%, Al: 5 to 6%, Nb: 10 to 13%, Ti: 3 to 5%, Cr: 2 to 3%, Co 15 to 17%.
Method for producing a heat-resistant alloy excellent in high temperature strength of the present invention, Vickers hardness HV420 above a plate material, less than 560, the tensile strength of 1080MPa or more, less than 1580MPa, the wire Vickers hardness HV360 or more, less than 630, the tensile strength 1210MPa As described above, the pressure is preferably less than 2170 MPa.
In the method for producing a heat-resistant alloy having excellent high-temperature strength according to the present invention, the aging treatment performed after forming into a predetermined shape is performed at 800 ° C. for 2 to 3 hours, so that the bending stress of 550 MPa is obtained at 700 ° C. for 96 hours. It is preferable that the content is 30% or less.
The heat-resistant alloy spring of the present invention is made of a heat-resistant alloy having any one of the above characteristics.

According to the present invention, a fine γ ′ phase having a particle size of 20 to 50 nm hardly grows even at a high temperature and occupies most of the structure. It is possible to provide a heat-resistant alloy having sagability and excellent corrosion resistance. In addition, by containing Co, Ni, Cr, Mo, Nb, Fe, Al, and Ti in a balanced manner as described above, the heat resistant alloy has excellent corrosion resistance, high hardness, excellent workability, and does not easily deteriorate toughness. Can provide.
The fine γ ′ phase effectively acts as a dislocation slip-preventing effect, and since the γ ′ phase hardly grows even at high temperatures, it can provide a heat-resistant alloy that is particularly excellent in high-temperature strength. Moreover, since the generation of cracks during the deformation process can be suppressed by keeping the ratio of the μ phase low, it is possible to provide a heat-resistant alloy that is difficult to cause a decrease in toughness and has excellent workability.
According to the present invention, since the γ ′ phase is a γ ′ phase containing Nb, Cr, Co, and Ti in addition to Ni and Al, even if exposed to a high temperature for a long time, the γ ′ phase hardly grows. It is possible to provide a heat-resistant alloy that can withstand high temperature and long-time heating.

According to the production method of the present invention, it is possible to produce a heat-resistant alloy with improved high-temperature strength due to aging effects, high hardness, excellent workability, and excellent sag resistance.
A heat resistant alloy spring can be formed from the heat resistant alloy of the present invention.

The structure photograph of the sample which was a heat-resistant alloy which concerns on this invention, was drawn at 20% reduction, and was aged at 800 degreeC for 2 hours. The structure photograph of the heat-resistant alloy which concerns on this invention, and was drawn at a rolling reduction of 40% and aged at 800 ° C. for 2 hours. The structure photograph of the heat-resistant alloy which concerns on this invention, and was drawn at 60% reduction, and was aged at 800 degreeC for 2 hours. An example of the compression spring manufactured with the heat-resistant alloy which concerns on this invention is shown, Fig.4 (a) is a side view, FIG.4 (b) is a front view. The other example of the torsion spring manufactured with the heat-resistant alloy which concerns on this invention is shown, Fig.5 (a) is a front view, FIG.5 (b) is a side view. Explanatory drawing which shows the size of the test piece used for the tensile test in an Example. The phase diagram computed using Thermo-Calc about Co-35Ni-17.5Cr-8Mo-3Nb-2Al alloy. The phase diagram calculated using Thermo-Calc about Co-35Ni- (10-25) Cr-4Mo-3Nb-2Al alloy. The phase diagram calculated using Thermo-Calc about Co-35Ni- (10-25) Cr-8Mo-3Nb-2Al alloy. The phase diagram calculated using Thermo-Calc about Co-45Ni- (10-25) Cr-4Mo-3Nb-2Al alloy. The phase diagram computed using Thermo-Calc about Co-45Ni- (10-25) Cr-8Mo-3Nb-2Al alloy. Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C alloy cold-rolled at a processing rate of 20% and aged at 800 ° C for 2 hours Stress strain diagram. Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C alloy cold rolled at a processing rate of 60% and aged at 800 ° C for 2 hours Stress strain diagram. The graph which shows the result of having measured the temperature dependence of the yield stress of each sample produced with the processing rate 20%, 60%, and 66% in the Example. The graph which shows the result of having measured the temperature dependence about the yield stress of each sample produced with the processing rate 20%, 40%, and 66% and SPRON510 (trademark). Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C alloy rolled at 20% and 60%, showing the microstructure of the sample 16A is a structural photograph of a sample with a processing rate of 20%, FIG. 16B is a structural photograph of a part of the structure shown in FIG. 16A, and FIG. 16C is a sample with a processing rate of 60%. FIG. 16 (d) is an enlarged structure photograph of a part of the structure shown in FIG. 16 (c). FIG. 17 (a) shows the metal structure before the heating when the alloy sample was swaged at a processing rate of 20% and aged at 800 ° C. for 3 hours and then heated at 700 ° C. Fig. 17 (b) is a metal structure photograph after heating for 96 hours, Fig. 17 (c) is a metal structure photograph after heating for 276 hours, and Fig. 17 (d) is a metal structure photograph after heating for 380 hours. FIG. 18 (a) shows the metal structure of each alloy sample subjected to swaging processing at a processing rate of 40%, aging at 800 ° C. for 3 hours and heating at 700 ° C. for each elapsed time. FIG. 18B is a metal structure photograph after heating for 96 hours, FIG. 18C is a metal structure photograph after heating for 276 hours, and FIG. 18D is a metal structure photograph after heating for 380 hours. FIG. 19A shows the metal structure for each elapsed time when the alloy sample was swaged at a processing rate of 66% and aged at 800 ° C. for 3 hours and then heated to 700 ° C. FIG. 19 (a) shows the metal structure before heating. FIG. 19B is a metal structure photograph after heating for 96 hours, FIG. 19C is a metal structure photograph after heating for 276 hours, and FIG. 19D is a metal structure photograph after heating for 380 hours. FIG. 20A shows the result of measuring the element distribution with a three-dimensional atom probe in the region where the γ ′ phase is present in the alloy sample, FIG. 20A shows the Ni distribution, and FIG. 20B shows the Al distribution. FIG. 20C shows a distribution of Nb, and FIG. 20D shows a distribution of Ni—Al—Nb. FIG. 21A shows the result of measuring the equi-concentration surface of each element with a three-dimensional atom probe in the region where the γ ′ phase is present in the same alloy sample. FIG. FIG. 21B is a diagram showing the Ni concentration surface, and FIG. 21C is a diagram showing the Nb concentration surface. The graph which shows the result of having conducted the elemental analysis for every position with the three-dimensional atom probe about the area | region where (gamma) 'phase exists about the alloy sample. FIG. 23A shows the result of elemental analysis of the alloy sample by an atom probe using a focused ion beam for the region where the γ ′ phase exists. FIG. 23B shows the distribution of Ni, and FIG. The figure which shows distribution of Al, FIG.23 (c) is a figure which shows distribution of Nb. The graph which shows the micro Vickers hardness about the sample which carried out the swaging process with the processing rate 20%, 40%, and 66% about the alloy sample, the sample before a process, and the sample after a predetermined time heating. The graph which compares and shows the micro Vickers hardness after aging of the sample swage-processed by the processing rate 20% about the alloy sample and a Waspalloy (trademark) sample. FIG. 26 (a) shows a metallographic photograph of a Waspaloy (registered trademark) sample, and FIG. 26 (b) shows a structural photograph of a Waspalloy (registered trademark) sample and an alloy sample obtained in the example. The metal structure photograph after heating an example alloy at 800 degreeC for 3 hours, and heating to 700 degreeC for 96 hours. FIG. 27A shows the result of measuring the true stress-true strain relationship for each temperature of the alloy sample. FIG. 27A shows the measurement result at 500 ° C., FIG. 27B shows the measurement result at 700 ° C., and FIG. (C) is a graph which shows the measurement result of 900 degreeC. FIG. 28A shows a transmission electron micrograph of a metal structure after 20% rolling of the same alloy sample and aged at 800 ° C. for 3 hours. FIG. 28A is a structure photograph showing a state in which stacking faults are arranged in layers. FIG. 28B is a bright-field image of a structure photograph showing a state in which stacking faults intersect, and FIG. 28C is a dark-field image of a stacking fault portion. FIG. 29 (a) is a structural photograph showing a bright field image of a part selected in the γ ′ phase region in the same alloy sample, and FIG. 29 (b) is a view of another part selected in the γ ′ phase region in the alloy sample. A tissue photograph showing a dark field image. The graph which shows how to obtain | require the load loss rate in a compression spring. In an Example, the graph which shows the comparison with the conventional material of the load loss rate with respect to the fastening time in the sag test of a compression spring. In an Example, explanatory drawing which shows the comparison with the conventional material of the load loss rate with respect to the clamping temperature in the sag test of a compression spring. In an Example, the graph which shows an example of the torsion spring used for the sag test. FIG. 34 (a) is a structural photograph showing the state of a precipitate after 40% wire drawing of an example alloy according to the present invention, and FIG. 34 (b) is a precipitate after 60% wire drawing of the alloy. The organization photograph which shows the state of.

The heat-resistant alloy excellent in high temperature strength according to the embodiment of the present invention will be described below.
The heat-resistant alloy of the present embodiment has a mass ratio of Ni: 34 to 46%, Cr: 13.5 to 18%, Mo: 2 to 9%, Nb 2.5 to 3.5%, Fe: 1 to 2%, Al: 1.5 to 2.5%, Ti: 0.6 to 1.0%, the balance is Co and inevitable impurities, and the structure is γ matrix and particle size 20 to 50 nm It is characterized by being mainly composed of a γ 'phase. The γ ′ phase occupying most of the tissue is preferably contained in a volume ratio of 52 to 65%, and a part of the tissue may contain a μ phase having a volume ratio of 10% or less.
In the heat-resistant alloy of the present embodiment, in the above composition, Mn: 0.3 to 0.6%, Si: 0.15 to 0.25%, Zr: 0.04 to 0.06%, C: 0 It is preferable to contain 0.04 to 0.06%.

In the heat-resistant alloy of this embodiment, the γ ′ phase is represented by a composition formula (Ni, Co) ₃ (Al, Nb, Cr, Co, Ti), and has a mass ratio of Ni: 57 to 61%, Al: 5 It is preferable that it is a composition containing -6%, Nb: 10-13%, Ti: 3-5%, Cr: 2-3%, Co15-17%.
In addition, when it describes using-about the content range of each element in this specification, unless otherwise indicated, an upper limit and a lower limit shall be included. For this reason, for example, when expressed as Ni: 59 to 61%, it means that Ni is included in the range of 59% to 61%.

The reason why the composition range is limited in the heat resistant alloy of the present embodiment will be described below.
Cr is an indispensable component for ensuring corrosion resistance and has an effect of strengthening the matrix. However, if it is less than 13.5%, the effect on high temperature sag resistance is weak, and if it exceeds 18%, it is processed during heat treatment. Since a σ phase that affects the toughness and toughness is generated, it is set to 13.5% or more and 18% or less.
Mo has the effect of strengthening the solid solution by dissolving in the matrix, the effect of increasing the work hardening ability, and the effect of increasing the corrosion resistance in the coexistence with Cr. However, if it is less than 2%, the desired effect cannot be obtained. If it exceeds 50%, the μ phase is likely to be generated, so the content was set to 2% or more and 9% or less.

Ni stabilizes the face-centered cubic lattice phase, maintains workability, and improves corrosion resistance. However, in the composition range of Co, Cr, Mo, Nb, and Fe of the alloy of the present invention, when Ni is less than 34% It is difficult to obtain a stable face-centered cubic lattice phase, and if it exceeds 46%, the mechanical strength decreases.
Co itself has a large work-hardening ability, and has the effect of reducing notch brittleness, increasing fatigue strength, and increasing high-temperature strength, but the effect is weak at less than 17%, and when it exceeds 41% in the proposed composition, it is a matrix. Is too hard to process, and the face-centered cubic lattice phase becomes unstable with respect to the close-packed hexagonal lattice phase.

Ti has the effect of strong deoxidation, denitrification, desulfurization, and refinement of the ingot structure, but the effect is weak at less than 0.6%, for example, 1% is not a problem. Inclusions increase, η phase (Ni ₃ Ti) precipitates and the toughness decreases, and Ti is effective for forming a stable γ ′ phase by replacing Al, so 0.6% It was made into 1.0% or less above.
Mn has the effect of deoxidation and desulfurization and the effect of stabilizing the face-centered cubic lattice phase, but if it is too much, the corrosion resistance and oxidation resistance are deteriorated.

Fe has an effect of strengthening the solid solution in the matrix, but if too much, the oxidation resistance is lowered. Therefore, the upper limit is set to 2.0%, and is set to 1% or more and 2% or less.
C forms a carbide with Cr, Mo, Nb, W, etc. in addition to solid solution in the matrix, and has an effect of preventing coarsening of crystal grains. However, if too large, C causes toughness reduction, corrosion resistance deterioration, and the like. 0.04% to 0.06%.
Nb is effective for the formation of a stable γ ′ phase by substituting Al like Ti, and when added in a large amount, the σ phase and δ phase (Ni ₃ Nb) are precipitated and the toughness is lowered. Since it has the effect of strengthening the solution by solid solution and increasing the work hardening ability, it is set to 2.5% or more and 3.5% or less.
Al becomes an element necessary for γ ′ phase precipitation together with effects of improving deoxidation and oxidation resistance. However, when it is contained in a large amount, the precipitation of the μ phase increases and the aimed effect cannot be obtained, so the content was made 1.5% to 2.5%.

Zr has the effect of increasing the grain boundary strength at high temperature and improving the hot workability, but if it is too much, the workability deteriorates conversely, so Zr is 0.04% or more and 0.06% or less. It was.
Si serves to improve the oxidation resistance of the alloy and has a deoxidizing effect in the casting process. When added in a large amount, it makes it easier to form a μ phase and a σ phase, so 0.15% or more and 0.25%. It was as follows.

  As an example, the heat-resistant alloy of this embodiment has a structure including a γ parent phase of 25 to 48% by volume, a γ ′ phase of 52 to 65%, and a μ phase of 10% or less. The μ phase may be a structure not precipitated.

FIG. 1 shows a processing rate of 20 for an alloy of the composition Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C obtained in the examples described later. 2 is a photograph (1000 times) showing an alloy structure after aging treatment at 800 ° C. for 2 hours after cold working at%. FIG. 2 is a photograph (1000 times) showing the alloy structure after aging treatment at 800 ° C. for 2 hours after cold working at a working rate of 40%. FIG. 3 shows the alloy at a working rate of 60%. It is a photograph (1000 times) which shows the alloy structure after aging treatment at 800 ° C for 2 hours after cold working.
The structure of the alloy according to the present invention has a region in which a μ phase that looks like white spots is precipitated as shown in the structure photographs of FIGS. 1 to 3, and in the structure of other dark portions, the particle size is about 20 to 50 nm. The particulate γ 'phase has a fine region in which it is deposited at intervals of about 10 to 20 nm, and the γ parent phase covers the periphery. The detailed structure of the dark color portion will be described in detail in the analysis results of examples described later.

In the alloy of the present embodiment, gamma 'phase usually simple _Ni 3 rather than Al of NiCrCoAl based heat resistant alloy of this type is produced, the L1 ₂ structure represented by the composition formula _Ni 3 (Al, Nb) gamma 'Phase, more specifically, a γ' phase represented by a composition formula (Ni, Co) ₃ (Al, Nb, Cr, Co, Ti).
And this γ 'phase is Ni: 59-61%, Al: 5-6%, Nb: 10-13%, Ti: 3-5%, Cr: 2-3%, Co15-17% by mass ratio It is preferable that it is a composition containing. In the alloy of the present embodiment, Nb is dissolved in the γ ′ phase produced by Ni ₃ Al, and a γ ′ phase containing other additive elements such as Cr, Co, and Ti is produced.

Thus, in addition to Ni ₃ Al, Nb is contained, and in addition, Cr, Co, and Ti are contained, so that the γ ′ phase generated in the alloy of this embodiment is about 20 to 50 nm. Further, by maintaining the particle size within this range even after the fine particles have undergone a thermal history, the characteristics as a heat-resistant alloy excellent in both mechanical properties and heat resistance are exhibited.
In alloys of this system, L1 ₂ structure gamma 'phase represented by the composition formula _Ni 3 without grain growth at high temperatures (Al, Nb), and more formula _{(Ni, Co) 3 (Al} , Nb, Cr, The precipitation of the γ ′ phase represented by (Co, Ti) is a knowledge obtained for the first time this time and shows the specificity of the alloy of the present invention. At the same time, the fact that this type of alloy having the γ ′ phase has developed the Suzuki effect is also the first knowledge obtained this time, indicating the specificity of the alloy of the present invention.

Regarding mechanical properties, for example, Vickers hardness HV 420 or more and less than 560 for a plate material, tensile strength 1080 MPa or more and less than 1580 MPa, for a wire material, Vickers hardness HV 360 or more and less than 630, tensile strength 1210 MPa or more and less than 2170 MPa, and Young's modulus E is Excellent mechanical properties of 210 to 250 GPa, a transverse elastic modulus G of 80 to 95 GPa, and a Poisson's ratio of 0.31 to 0.32 are obtained.
Further, as an example, excellent characteristics as a spring material are obtained such that the load loss rate at 96 ° C. for 96 hours at a torsional stress of 392 MPa is 35% or less.
In the heat resistant alloy of the present embodiment, the average particle size of the μ phase (Mo rich) is 90 to 380 μm, and is preferably 10% or less, for example, less than 9% by volume. The μ phase is a brittle phase in the alloy structure of the present embodiment, and when a large amount of the μ phase precipitates, the toughness decreases and the workability deteriorates.

In the alloy of the present embodiment, the reason why the mechanical characteristics and the spring characteristics are excellent is that the γ ′ phase is a fine particle of about 20 to 50 nm as described above, and has a composition containing the plurality of atoms described above. In addition to the Suzuki effect.
The Suzuki effect is one of the interactions between dislocations and solute atoms. The dislocations of fcc (face-centered cubic lattice) and hcp (close-packed hexagonal lattice) alloys are often extended dislocations and are in a different energy state from the surroundings, and solute atoms segregate in the extended dislocations. To do. When dislocations move in this part, a thermally non-equilibrium segregated part remains and at the same time a part without segregation occurs. Both of these create a new large portion of energy, which resists the movement of dislocations. Its sticking force is comparable to the elastic interaction, but it is more difficult to get out of sticking because of the large dislocation extension. This interaction between dislocations and solute atoms is generally called a chemical interaction or a Suzuki effect. When the Suzuki effect is manifested, mechanical properties such as hardness and tensile strength of the alloy are improved.
In the alloy of the present embodiment, dislocations are introduced by cold working and are aged after a number of stacking faults are introduced, so that the Suzuki effect is manifested, and the effect and phase due to the presence of the fine γ ′ phase described above. As a result, the above-described excellent mechanical properties can be obtained, and excellent sag resistance when used as a spring material can be obtained.

"Production method"
In order to manufacture the alloy according to the present embodiment having the above composition, the material blended so as to have the above composition ratio is melted by a melting method such as vacuum melting using a high frequency vacuum induction melting furnace or the like, and cooled in a furnace and cast. A lump (ingot) is produced, and the obtained ingot is subjected to heat treatment for solid solution. This ingot can be obtained by subjecting it to a desired shape by cold plastic working and then aging treatment. Or you may give a hot plastic working as needed before a cold plastic working in the said process.
The heat treatment for solid solution (homogenization treatment) indicates a treatment of uniformly heating to 1150 to 1250 ° C. for about 4 to 8 hours. The hot rolling temperature is preferably within the above-mentioned range, for example, 1200 ° C.

Then, after performing an aging treatment at 700 to 850 ° C. for 2 to 4 hours, it is molded into a desired shape and further subjected to an aging treatment at 700 to 850 ° C. for 2 to 4 hours, or the desired shape In place of the aging treatment before forming into a cold, after cold working at a processing rate of 20% or more and less than 60% and processing into the desired shape, after heating at 700 to 850 ° C. for about 1 to 4 hours, By returning to room temperature, a heat-resistant alloy having a desired shape can be obtained. When the target shape is a leaf spring, it is processed into the shape of a leaf spring, and when it is a coil spring, it is processed into the shape of a coil spring.
Alternatively, a heat treatment for homogenization may be performed at a temperature of about 1200 ° C. for the purpose of eliminating the heterogeneous precipitate after the aging treatment.
The obtained heat-resistant alloy has a structure in which the γ parent phase is 25 to 48% by volume, the γ ′ phase is 52 to 65%, and the μ phase is 10% or less. The μ phase may be a non-precipitated structure.

The alloy of the present embodiment has the above-described excellent mechanical characteristics and spring characteristics, and thus is effective as various spring materials, and exceeds 650 ° C. which is the limit of sag resistance of Inconel (registered trademark) 718. In addition, depending on the processing rate, excellent sag resistance exceeding 700 ° C., which is the limit of sag resistance of Waspalloy (registered trademark), can be obtained. For this reason, it is suitable as a spring material corresponding to a structure in which the exhaust valve is arranged closer to the engine in order to increase the efficiency of the engine. It is possible to provide a spring material that is less likely to deteriorate in elasticity.
Of course, the heat-resistant alloy of the present embodiment can be widely applied to other heat-resistant members for heat-resistant members such as engines and turbines.

FIG. 4 shows an example of a coil spring formed using the heat-resistant alloy according to the present embodiment. This coil spring 1 is formed by forming the above-mentioned heat-resistant alloy into a coil by wire drawing, It can be obtained by polishing the seating surfaces 2 at both ends of the coil and applying the aging treatment under the above-mentioned conditions after polishing.
FIG. 5 shows a second example of a coil spring formed using the heat-resistant alloy according to this embodiment, and this coil spring 3 is formed into a coil shape by forming the above-mentioned heat-resistant alloy into a wire material by wire drawing. It can be obtained by aging treatment under the above-mentioned conditions while processing and keeping the coil end portions 3a, 3b partially linear.

  The coil spring shown in FIGS. 4 and 5 formed using the heat-resistant alloy of this embodiment is merely an example of a spring shape. The shape of the spring formed using the heat-resistant alloy of this embodiment may be any shape spring such as a leaf spring, a torsion bar, a bamboo shoot spring, a disc spring, and a spring.

Examples of the heat-resistant alloy according to the present invention will be described below, but it is needless to say that the heat-resistant alloy of the present invention is not limited to the description of the following examples.
As a heat-resistant alloy sample, Ni: 35%, Cr: 17.5%, Mo: 4.0%, Nb: 3.0%, Fe: 1.6%, Al: 2%, Ti: 0.8%, Mn: 0.5%, Si: 0.2%, Zr: 0.05%, C: 0.05%, materials are blended so as to have a composition represented by the balance Co and inevitable impurities, and a high-frequency vacuum induction melting furnace After obtaining an ingot, an alloy sample A was produced by the following procedure.
As a heat-resistant alloy sample, a material was prepared so as to have a composition of Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C. After blending, an ingot was obtained by a high-frequency vacuum induction melting furnace, and then a heat-resistant alloy sample B was produced by the following procedure.
As a heat-resistant alloy sample, a material was prepared so as to have a composition of Co-45Ni-17.5Cr-4Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C. After blending, an ingot was obtained with a high-frequency vacuum induction melting furnace, and then a heat-resistant alloy sample C was produced according to the following procedure.

As a heat-resistant alloy sample, after blending the materials so as to have a composition of Co-35Ni-14Cr-4Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C After obtaining an ingot with a high-frequency vacuum induction melting furnace, a heat-resistant alloy sample D was produced by the following procedure.
As a heat-resistant alloy sample, after blending the materials so as to have a composition of Co-45Ni-14Cr-4Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C After obtaining an ingot with a high-frequency vacuum induction melting furnace, a heat-resistant alloy sample E was produced according to the following procedure.
As a heat-resistant alloy sample for comparison, the composition is Co: 13.2%, Ni: 58.8%, Cr: 19.5%, Mo: 4.2%, Ti: 3%, Al: 1.3. After the materials were mixed as described above, an ingot was obtained by a high-frequency vacuum induction melting furnace, and then a comparative sample F was manufactured by the following procedure. This comparative sample F has a composition of an alloy known as Waspalloy (registered trademark).

After obtaining a plurality of ingots having a length of 100 mm and a diameter of 28 mm from the molten alloy of each composition, each ingot was homogenized at 1200 ° C. for 3 hours, and then a plate material having a thickness of 10 mm was obtained by hot forging. Each plate was heated at 1200 ° C. for 1 hour. Each obtained plate material was divided into three equal parts, and each of the three divided plate materials was subjected to cold rolling with a processing rate of 20%, 40%, 60%, and after aging treatment at 800 ° C. for 2 hours, From the thickness center of the test piece, a tensile test piece or a sag test piece was cut out by electric discharge machining and lapped and polished for use in the test. In addition, about the sample of AF, micro Vickers hardness HV (weight 200gr) was measured in the plate | board thickness center part.
Comparative sample F was prepared by preparing a rod having a length of 70 mm, a rod having a diameter of 80 mm, a rod having a length of 80 mm and a rod having a diameter of 100 mm, and a rod having a diameter of 100 mm and a diameter of 20 mm. The plate was hot forged and subjected to each test.
As shown in FIG. 6, the tensile test piece has a total length of 10 mm, a width of 3 mm, a thickness of 0.38 mm, and a standard test piece in which a belt-like portion having a width of 1.8 mm and a length of 4.8 mm is formed at the center in the length direction. It is. The dimensions of the other parts are as shown in FIG. Tensile strength and Young's modulus were measured using this test piece.
The test piece has a plate thickness of 0.22 mm, a plate width of 3 mm, and a length of 20 mm. The sag test is performed at a temperature of 700 ° C., a height spacer of 1.3 mm, a bending stress of about 550 MPa, and a bending span of 13 mm. Sag after 48 hours and 96 hours was measured. The settling rate (%) was calculated according to the calculation formula of (H ₁ / H ₀ ) × 100%. Here, H ₀ indicates the spring height when a predetermined bending stress is applied, and H ₁ indicates the spring height after unloading. The smaller this value, the less the sag.
The results of the hardness test, tensile test, and sag test are shown in Table 1 below.

In each of Samples A to F, the higher the rolling reduction during rolling, the harder and the tensile strength tends to increase. The Young's modulus is slightly small when the rolling reduction is 20%, but tends to be saturated when the rolling reduction is 40% or more. As for sag, any sample has a tendency to sag more easily as the reduction ratio becomes higher, and the developed materials A, D, and E have a reduction rate of 20% compared to the comparative material F (Waspalloy: registered trademark). Became smaller.
When the processing rate is increased, the μ phase is likely to precipitate due to processing energy. In 40% processing, the μ phase is about 9% by volume, and in 60% processing, the μ phase is about 18% by volume. When the μ phase is precipitated in this type of alloy, Mo of the parent phase is eroded by the μ phase, and it seems that the sag resistance is lowered. For this reason, the precipitation of the μ phase is preferably 10% or less, more preferably 9% or less by volume ratio (see FIG. 34 described later).

Next, in this type of alloy, the proper range of Cr content and Al content will be verified.
FIG. 7 is a calculation state diagram using Thermo-Calc (manufactured by Thermo-Calc Software: ver. 5.0, database: TCW, N17) for the Co-35Ni-17.5Cr-8Mo-3Nb-2Al alloy. As shown in FIG. 7, it can be determined that the γ ′ phase can be precipitated in the range where the Al content is in the range of 0 to 3.5 mass%. For this reason, it can be judged that the addition amount of Al is preferably in the range of 1.5 to 2.5% in order to obtain a sufficient amount of γ ′ phase for the alloy of this composition system.

FIG. 8 is a calculation state diagram using the Thermo-Calc for a Co-35Ni- (10-25) Cr-4Mo-3Nb-2Al alloy, and FIG. 9 is a Co-35Ni- (10-25) Cr-8Mo-3Nb- It is a calculation state figure using the Thermo-Calc about 2Al alloy.
FIG. 10 is a calculation state diagram of the Co-45Ni- (10-25) Cr-4Mo-3Nb-2Al alloy using the Thermo-Calc, and FIG. 11 is a Co-45Ni- (10-25) Cr-8Mo-3Nb- It is a calculation state figure using the Thermo-Calc about 2Al alloy.
From these phase diagrams and the results shown in Table 1, in order to obtain an alloy structure having a γ matrix and a γ ′ phase and obtain the above-mentioned excellent sag resistance and excellent mechanical properties in the present alloy, Cr It can be judged that the content is preferably in the range of 13.5 to 18% by mass.

FIG. 12 shows a stress strain diagram of a sample prepared at a reduction rate of 20% for the alloy of sample A, and FIG. 13 shows a stress strain diagram of a sample prepared at a reduction rate of 60% for the alloy of sample A. Show.
As shown in FIG. 12, a strain curve presumably due to the Suzuki effect at 500 ° C. was shown. For this reason, it is considered that the alloy of this example exhibits excellent mechanical properties as a result of the stacking fault suppressing the movement of dislocations by the Suzuki effect.

FIG. 14 shows the temperature dependence of the yield stress of each sample produced by cold rolling at a working rate of 20% and 60% or by swaging at a working rate of 20% and 66% in the example alloys. The composition of the alloy is Co-35Ni-17.5Cr-8Mo-3Nb-2Al-1.6Fe-0.8Ti-0.5Mn-0.2Si-0.05Zr-0.05C.
In cold rolling, a plate material having a thickness of 10 mm is produced by hot forging, the surface scale is removed by a grinder, the plate thickness is set to 9.5 mm, and the plate material is roughly divided into three parts (length) by a cutting machine. The sample plate is divided into three equal directions, approximately 40 mm x 45 mm, and a plate thickness of 9.5 mm. A plate thickness of 7.6 mm (processing rate of 20%), a plate thickness of 5.7 mm (processing rate of 40%), and a plate thickness of 3 Cold-rolled at each processing rate up to 0.8 mm (processing rate 60%). Further, each sample plate was cut into half sizes (two equal parts in the length direction) and subjected to aging treatment at 750 ° C. × 2 hours and 800 ° C. × 2 hours.
In the swaging process, a round bar of φ13.0 mm was processed at each processing rate. The outer diameter after processing was φ11.63 mm for a sample with a processing rate of 20%, φ10.07 mm for a sample with a processing rate of 40%, and φ7.58 mm for a sample with a processing rate of 66%. In the case of cold rolling of the plate material, the processing rate (%) is calculated by a calculation formula of (1-plate thickness after processing / original plate thickness) × 100.
The sample alloy with any processing rate shown in FIG. 14 showed an excellent value as a heat-resistant alloy.

FIG. 15 shows the temperature dependence of the yield stress and the temperature dependence of the yield stress of SPRON510 (registered trademark) for each sample produced by swaging at 20%, 40%, and 66% of the processing rate in an alloy having the same composition. SPRON510® is a Co—Ni based superalloy having a composition of 35% Ni-32% Ni-20% Cr-10% Mo.
SPRON510 (registered trademark) is an alloy exhibiting excellent yield stress as a heat-resistant alloy, but the example alloys exhibited higher yield stress than SPRON510 (registered trademark) at any processing rate. In addition, it can be seen that the example alloys having the above composition are excellent in the strengthening mechanism by aging, and the yield strength can be remarkably improved after aging.

  FIG. 16 is an enlarged structure photograph showing the γ ′ phase region of the sample B rolled at a working rate of 20% and a sample rolled at a working rate of 60%. From the structure photograph shown in FIG. 16, the presence of a γ ′ phase having a particle size much smaller than 100 nm was confirmed. The thermal stability of this γ ′ phase and detailed analysis results will be described below.

FIGS. 17 to 19 show the results of confirming the growth state of the γ ′ phase by heat-treating the alloy of Sample B obtained by 20%, 40%, and 66% swaging to 700 ° C. for a predetermined time. Indicates.
17 to 19, even when the γ ′ phase having a particle size of 20 to 30 nm before heating is heated for 96 hours, 276 hours, and 380 hours, the particle size of 30 It remained at a size of ˜40 nm. From this test result, it can be assumed that the alloy of this example can obtain extremely excellent heat resistance because the γ 'phase hardly grows even when heated for a long time.

With respect to the alloy of Sample B subjected to the 20% processing rate swaging, elemental analysis was performed on the region where the γ ′ phase was precipitated using a three-dimensional atom probe.
The measurement results of Ni, Al, and Nb are shown in FIGS. As shown in each figure, it was found that Ni, Al, and Nb were concentrated at the position where the γ ′ phase was present.
FIGS. 21A to 21C show the Al-6% isoconcentration surface, the Ni-50% isoconcentration surface, and the Nb-5% isoconcentration surface.
FIG. 22 shows the result of measuring the concentration of each position of Nb, Ni, Co, Mo, Al, Fe, Ti, and Cr so as to correspond to the measurement position of the Al-6% isoconcentration surface shown in FIG. Show.
As shown in FIG. 22, since Ni, Co, Cr, Al, Nb, and Ti are present in accordance with the generation position of the γ ′ phase, the γ ′ phase of this example alloy has the composition formula (Ni, Co ) ₃ (Al, Nb, Cr, Co, Ti) It can be determined that it is a γ ′ phase represented by. Moreover, from this elemental analysis result, Ni: 57-61%, Co15-17%, Al: 5-6%, Nb: 10-13%, Ti: 3-5%, Cr: 2-3% by mass ratio It can be estimated that the γ ′ phase contains.

FIG. 23 shows another result of elemental analysis of a region where the γ ′ phase exists for the same alloy sample by an atom probe using a focused ion beam. The γ ′ phase is determined from the element distribution state of Ni, Al, and Nb. It can be seen that the particle size is about 20 nm. It was also found that the distance between particles with other adjacent γ ′ phases was about 10 nm.
The γ ′ phase represented by the composition formula (Ni, Co) ₃ (Al, Nb, Cr, Co, Ti) has a particle size of about 20 nm and is separated from the other γ ′ phase at a distance of 10 nm. Therefore, it is estimated that the excellent mechanical properties of the present alloy are exhibited. In other words, the expression of the Suzuki effect requires the presence of fine regions in which fine particles for suppressing the movement of dislocations are aligned at a predetermined interval, and the arrangement of stacking faults is related to the fine particle arrangement. Need to be.
Due to the arrangement of the γ ′ phase and the presence of stacking faults present in the metal structure, the alloy of this example has an excellent mechanical property due to the presence of a fine γ ′ phase that hardly grows at high temperatures. It can be presumed that excellent mechanical properties far exceeding the above-mentioned SPRON510 (registered trademark) are obtained due to the manifestation of the effect.

FIG. 24 shows the result of measuring the micro Vickers hardness (weight 1 kg) at room temperature for the alloy of Sample B.
The hardness of the samples with any processing rate is improved by performing an aging treatment. Moreover, the hardness improved as the processing rate increased.
FIG. 25 shows the micro Vickers hardness (weight 1000 gr) after swaging the alloy of sample B at a processing rate of 20%, aging treatment at 800 ° C. for 3 hours, and heating to 700 ° C. for 100 hours or 240 hours. ) Shows the measurement results. For comparison, the micro Vickers hardness after subjecting the alloy of Sample F to 20% rolling at a processing rate, aging treatment at 800 ° C. for 3 hours, and heating to 700 ° C. for 100 hours or 240 hours. The measurement results are shown.
The alloy of sample F is a prominent alloy as a spring material on the market, whereas the alloy of sample B showed higher hardness even after being heated to a high temperature for a long time.

FIG. 26 is a structure photograph for comparing and comparing the metal structures of the sample F alloy and the sample B alloy.
The metal structure of the alloy of Sample F shown in FIG. 26 (a) was dark-fielded by a transmission electron microscope for the alloy after heat treatment at 1080 ° C. for 4 hours, aging at 850 ° C. for 2 hours, and heating at 760 ° C. for 10 hours. This is a metallographic image of the image. The metal structure shown in FIG. 26A is a mixed phase structure of a primary particle γ ′ phase (particle size 195 nm) having a large particle size and a γ ′ phase (23.9 nm) having a fine particle size. It can be presumed that the metal structure shown in FIG. 26 (a) is heated and the γ ′ phase absorbs the surrounding γ ′ phase and grows and grows into a coarse primary crystal γ ′ phase.
On the other hand, the metallographic structure of the alloy of Sample B shown in FIG. 26 (b) is a field emission type for the alloy after uniform heat treatment at 1200 ° C., aging treatment at 800 ° C. for 3 hours, and heating to 700 ° C. for 96 hours. It is the metal structure image | photographed with the scanning electron microscope. As shown in FIG. 26A, the alloy of Sample B has a metal structure in which a uniform γ ′ phase having a particle size of 20 to 30 nm is precipitated.
Compared to the metal structure, the alloy of Sample B is an aggregate of a fine γ 'phase and a γ parent phase existing around it, and there is no primary crystal γ' phase with a coarsened grain size. It seems to have an excellent structure.

FIG. 27 shows the result of measuring the true stress-true strain relationship for each temperature for the same alloy sample. FIG. 27 (a) shows the measurement result at 500 ° C., and FIG. 27 (b) shows the measurement result at 700 ° C. FIG. 28 (c) shows the measurement result at 900 ° C. Each figure shows a sample before processing (ST), a sample swaged at a processing rate of 66%, a sample swaged at a processing rate of 20% and then aged at 800 ° C. for 3 hours, and a swaging processing at a processing rate of 66% Each sample after aging treatment at 800 ° C. for 3 hours is shown in comparison.
A flow due to dynamic strain aging labeled DSA in each figure was observed. From this, it can be assumed that the Suzuki effect is manifested by a mechanism similar to that of Splon 510 (registered trademark).

  FIG. 28 shows a photograph of the microstructure of the sample B alloy after 20% swaging and aging at 800 ° C. for 3 hours. As shown in FIG. 28 (a), the stacking faults are arranged in layers. As shown in FIG. 28 (b), it is possible to confirm the presence of a structure showing a state in which the stacking faults intersect, and as shown in FIG. 28 (c), a structure in which the stacking faults are aligned. I was able to confirm.

FIG. 29 (a) shows a transmission electron micrograph of an alloy obtained by subjecting the alloy of sample B to 66% swaging and aging treatment at 800 ° C. for 3 hours, and FIG. A transmission electron micrograph of an alloy that has been subjected to% swaging and aged at 800 ° C. for 3 hours is shown. Any of the in tissue photographs, it is understood that the gamma 'phase indicating an L1 ₂ structure is in spherical particle size 20 to 30 nm.

  A wire rod having a drawing rate of 0, 20, 40, 60% is manufactured using the developed alloy B described above, and a compression spring shown in FIG. 4 and a torsion spring shown in FIG. A sagging test was performed. Table 2 shows the tensile properties and Vickers hardness (weight 300 gr) of the wire used in the sag test. In Table 2, the alloy of sample B is abbreviated as development material 00, 20, 40, 60, and the alloy of sample F is abbreviated as comparative material F (Waspalloy: registered trademark). The developed material 00 is obtained by performing aging treatment at 800 ° C. for 3 hours after homogenizing heat treatment (Vickers hardness HV217) to precipitate a γ ′ phase (Vickers hardness HV372), and is not subjected to wire drawing.

After coiling a compression spring having the same shape as the compression spring 1 shown in FIG. 4 using the wires shown in Table 2 (surface bending strain of 10%), the seating surfaces at both ends of the spring are polished to a flat surface, an aging treatment, After setting, it was subjected to a settling test. The load loss rate after applying the load of the spring and each sample for 6 hours, applying the load for 24 hours, and applying the load for 96 hours was determined.
A compression spring having a wire diameter d of 1.4 to 1.8 mm, a center diameter D of 18.6 mm, a total number of windings of 5.8, and an effective number of windings of 3.8 was prototyped. The total length of the compression spring is 52 mm.

  Table 3 shows the relationship between the aging treatment conditions and the sag after the coiling process (surface processing rate of 10%) for the developed material 00. From this, it can be seen that by increasing the aging temperature after processing, the sag resistance is improved by the Suzuki effect. This trend was the same with other developed materials. It can also be seen that an aging treatment of 700 ° C. or higher is necessary to make the load loss rate 35% or less. Therefore, the developed material was subjected to an aging treatment of 800 ° C. × 2 h after coiling and subjected to a sag test.

FIG. 30 shows a method for measuring load loss. A load P ₁ and the height a at the time of shearing stress 392MPa load was measured by inserting a spring into the heating furnace load state, 6 hours at 700 ° C., 24 hours, unloading removed springs from the furnace after 96 hours · released was measured load P ₂ when loaded again to the height a. The load loss rate (%) was calculated according to the calculation formula of (1-P ₂ / P ₁ ) × 100%.

From the results shown in Table 4, the load loss rate at the time of tightening at 700 ° C. for 96 hours is 24% better for the alloy of Sample B than the No. 1 and No. 6 when the tightening stress is 392 MPa, and 650 Even in comparison between No. 2 and No. 7 at 96 ° C., the alloy of Sample B was 18% better.
Comparative material F (Waspalloy: registered trademark) using an alloy of sample F, which is a conventional material, is commercially available as an alloy having excellent sag resistance, but it is 20% and 40% processed with the sample according to the present invention. It can be seen that all the samples of the rate have sag resistance superior to that of the comparative material F. However, when the processing rate is 60%, the sag resistance is worse than that of the comparative material F.

  FIG. 31 is a comparison of the load loss rate of the developed material and the comparative material after a tightening temperature of 700 ° C. and a tightening stress of 392 MPa and after tightening times of 6 h, 24 h, and 96 h. From this, it can be seen that although the load loss increases in the order of the developed materials 00, 20, and 40 and the sag resistance deteriorates, the load loss is smaller than that of the comparative material F and exhibits excellent sag resistance. In addition, the developed material 60 is less sag resistant than the comparative material F. Since dislocations increase more than the dislocations suppressed by the Suzuki effect together with the processing rate, it is considered that the sag resistance deteriorates. FIG. 32 shows a comparison of the load loss rate with respect to the tightening temperature with the comparative material.

FIG. 33 shows a measurement position when a sag test is performed using the torsion spring 3 having the shape shown in FIG. A wire rod having a wire diameter of 1.8 mm made of the alloy of the sample B was produced by wire drawing with a processing rate of 20%, and using this wire rod, a center diameter D: 12.2 mm and a total number of turns of 9.5 in FIG. A torsion spring 3 having the size shown in the spring specifications of No. 5 was prototyped and subjected to aging treatment at 800 ° C. for 2 hours to prepare a torsion spring specimen.
With respect to this torsion spring test body, as shown in FIG. 33, one spring end 3a was fixed, and the other spring end 3b was set to a tightening angle of 53 to 66 °. The actual tightening stress was 500 to 620 MPa. In the torque measurement, when the tightening temperature was 700 ° C., the sag rate measured at an actual angle of 44 ° was calculated. As a measurement point in the case of tightening at 650 ° C., the sag rate of the actual tightening angle was measured. For comparison, a torsion spring specimen of the same size was made with the alloy of Sample F and subjected to a sag test under the same conditions. The alloy of sample F was subjected to an aging treatment at 800 ° C. for 4 hours after the spring forming in the same manner as the alloy of the above-described specimen.

Table 5 shows the results of the torsion spring sag test. The torque loss rate when tightened at 700 ° C. for 96 hours was 7.6% better for the alloy of sample B when compared with No. 9 and No. 11 when the tightening stress was 500 MPa, and No. 10 and No. at 620 MPa. Even in the comparison of .12, the alloy of Sample B was 15.9% better.
The torque loss rate at the time of tightening at 650 ° C. for 96 hours was 2.7% better for the alloy of sample B in comparison with No. 13 and No. 15 when the tightening stress was 500 MPa, and No. 14 at 620 MPa. In comparison with No. 16, Sample B was 2.6% better.

FIG. 34 (a) is a structural photograph showing the state of the precipitate (μ phase) of the metal structure after 40% wire drawing and aging treatment at 800 ° C. for 2 hours for the alloy of Sample B, FIG. ) Is a structure photograph showing the state of precipitates (μ phase) of the metal structure after 60% wire drawing of the alloy of Sample B and aging treatment at 800 ° C. for 2 hours.
With respect to the structure photograph shown in FIG. 34 (a), image determination by black and white binarization processing was performed, and the volume ratio of the white dot-like μ phase was calculated for the 40% wire drawing material of Sample B. %. Moreover, the same binarization process was performed with respect to the other sample of the same composition, and the volume ratio of the white dot-like μ phase could be calculated as 8.7%.
The structure photograph shown in FIG. 34 (b) was subjected to image determination by black-and-white binarization processing, and the volume ratio of the white dot-like μ phase was calculated for the 60% wire drawing material of Sample B. %. Moreover, when the same binarization process was performed with respect to the other sample of the same composition, the area ratio of the white dot-like μ phase could be calculated as 18.4%. Conversion from these binarization processes was carried out by converting the area ratio obtained from the binarization process to 3/2 to obtain a volume ratio. In the same alloy, the μ phase was not detected when the alloy was subjected to only aging treatment without drawing after homogenization heat treatment or 20% wire drawing and aging treatment at 800 ° C. for 2 hours.

  From the above-mentioned various test results, regarding the 60% wire drawn material, the volume fraction of the μ phase greatly exceeds 10% and becomes 14 to 16%, and Mo in the parent phase that improves sag resistance is the μ phase. It can be assumed that the resistance to sag at high temperatures has decreased due to the erosion. For this reason, when heat-treating the alloy according to the present invention by heat treatment, it is important to obtain a good sag resistance when the volume fraction of the μ phase is 10% or less, more preferably 9% or less. It can be judged.

  DESCRIPTION OF SYMBOLS 1 ... Compression spring, 3 ... Torsion spring, 3a, 3b ... Torsion spring edge part.

Claims (11)

  Ni: 34-46%, Cr: 13.5-18%, Mo: 2-9%, Nb 2.5-3.5%, Fe: 1-2%, Al: 1.5-2. 5%, including Ti: 0.6 to 1.0%, having a composition of the balance Co and inevitable impurities, and having a structure of a γ parent phase, a particle size of 20 to 50 nm and a volume ratio of 52 to 65% A heat-resistant alloy excellent in high-temperature strength, characterized by comprising a γ ′ phase and a μ phase having a volume ratio of 10% or less.   In the composition, Mn: 0.3 to 0.6%, Si: 0.15 to 0.25%, Zr: 0.04 to 0.06%, C: 0.04 to 0.06% The heat-resistant alloy excellent in high-temperature strength according to claim 1, comprising one or more kinds.   The γ ′ phase is represented by a composition formula (Ni, Co) 3 (Al, Nb, Cr, Co, Ti), and mass ratios of Ni: 57 to 61%, Al: 5 to 6%, Nb: 10 to 10%. The heat-resistant alloy having excellent high-temperature strength according to claim 1 or 2, comprising 13%, Ti: 3 to 5%, Cr: 2 to 3%, and Co 15 to 17%.   The plate material has a Vickers hardness of HV420 or more and less than 560, a tensile strength of 1080 MPa or more and less than 1580 MPa, and a wire material has a Vickers hardness of HV360 or more and less than 630, and a tensile strength of 1210 MPa or more and less than 2170 MPa. The heat-resistant alloy excellent in high temperature strength as described in any one of Claims 3. Ni: 34-46%, Cr: 13.5-18%, Mo: 2-9%, Nb 2.5-3.5%, Fe: 1-2%, Al: 1.5-2. An ingot containing 5%, Ti: 0.6 to 1.0% and having the balance Co and inevitable impurity composition is homogenized at 1150 to 1250 ° C. for 4 to 8 hours, and then 2 to 4 at 700 to 850 ° C. It was subjected to aging treatment time, and molded into a predetermined shape, further 700-850 Ri by the applying aging treatment 2-4 hours at ° C., and γ matrix phase, particle size 20~50nm and volume fraction 52 A method for producing a heat-resistant alloy excellent in high-temperature strength, characterized by having a structure having γ 'phase of ˜65% and a μ phase having a volume ratio of 10% or less. Ni: 34-46%, Cr: 13.5-18%, Mo: 2-9%, Nb 2.5-3.5%, Fe: 1-2%, Al: 1.5-2. An ingot containing 5%, Ti: 0.6 to 1.0% and having the balance Co and inevitable impurities composition is homogenized at 1150 to 1250 ° C. for 4 to 8 hours, and the processing rate is 20% or more and 60%. Less than a cold working, shape | mold in a predetermined | prescribed shape, and give an aging process for 1 to 4 hours at 700-850 degreeC, and (gamma) mother phase, particle size 20-50nm and volume ratio 52-65% A method for producing a heat-resistant alloy excellent in high-temperature strength, characterized by having a structure comprising a γ 'phase and a μ phase having a volume ratio of 10% or less . In the composition, Mn: 0.3 to 0.6%, Si: 0.15 to 0.25%, Zr: 0.04 to 0.06%, C: 0.04 to 0.06% The method for producing a heat-resistant alloy having excellent high-temperature strength according to claim 5 or 6 , wherein an ingot having a composition containing one kind or two or more kinds is used. The γ ′ phase is represented by a composition formula (Ni, Co) 3 (Al, Nb, Cr, Co, Ti), and mass ratios of Ni: 57 to 61%, Al: 5 to 6%, Nb: 10 to 10%. The heat-resistant alloy having excellent high-temperature strength according to any one of claims 5 to 7, comprising 13%, Ti: 3 to 5%, Cr: 2 to 3%, and Co 15 to 17%. Manufacturing method. By the manufacturing method as described in any one of Claims 5-8, Vickers hardness HV420 or more and less than 560 in a board | plate material, Tensile strength 1080MPa or more and less than 1580MPa, Vickers hardness HV360 or more and less than 630 in a wire rod, tensile strength 1210MPa or more, the production method of high temperature alloys in high temperature strength you characterized in that less than 2170MPa. By treating 2-3 hours at 800 ° C. the aging treatment performed after forming into a predetermined shape, claims, characterized in that a 700 ° C., more than 35% of sag due to 550MPa bending stress at 96 hours method for producing a high heat-resistant alloy in high temperature strength according to any one of 5 claims 9.   A heat-resistant alloy spring formed of the heat-resistant alloy according to any one of claims 1 to 4.