JP2007162041A - Ni-BASE SUPERALLOY WITH HIGH STRENGTH AND HIGH DUCTILITY, MEMBER USING THE SAME, AND MANUFACTURING METHOD OF THE MEMBER - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide an Ni-base superalloy which can provide high high-temperature strength and excellent ductility in the form of both a directionally solidified material and an ordinary casting and is suitable for application to industrial gas turbines and centrifugal wheels for turbochargers or microturbines. <P>SOLUTION: The Ni-base superalloy has a composition containing, by weight, 0.06 to 0.3% C, 0.01 to 0.05% B, 0.5 to 3.0% Hf, 10.2 to 25% Co, 1 to 12% Ta, 1.5 to 16% Cr, 2 to 15% W, 3.5 to 6.5% Al, 0.5 to 9% Re and 0.2 to 2% Nb. <P>COPYRIGHT: (C)2007,JPO&INPIT

Description

  The present invention relates to a Ni-base superalloy excellent in high-temperature strength and ductility, and more particularly, to a Ni-base superalloy having high strength and excellent ductility, regardless of whether it is a normal cast material or a unidirectionally solidified material. The present invention also relates to a centrifugal wheel for a turbocharger or a microturbine manufactured by a casting of a Ni-base superalloy, or a moving blade or a stationary blade of an axial flow gas turbine.

  In an aircraft engine gas turbine or an industrial gas turbine, a Ni-base superalloy is used for a part heated to a high temperature, for example, a moving blade or a stationary blade. Ni-based superalloys are also used in centrifugal wheels for turbochargers or microturbines.

  Combustion gas temperatures of gas turbines tend to increase year by year from the viewpoint of improving thermal efficiency, and in accordance with this, unidirectionally solidified materials are used for moving blades or stationary blades instead of conventional cast materials that have been used conventionally. Came to be used. Among gas turbines, single-crystal wings with higher high-temperature strength have been put to practical use in relatively small aircraft engine gas turbines.

  Many Ni-base superalloys developed for unidirectionally solidified materials contain Ti as a strengthening element (see, for example, Patent Document 1).

  Since many of the alloys developed for single crystal materials do not contain grain boundary strengthening elements, it is difficult to apply them to large-sized products that easily generate grain boundaries during casting, such as industrial gas turbines. In order to make it possible to apply the single crystal material to a large product, an alloy in which a grain boundary strengthening element such as C, B, Hf, or Zr is added to the single crystal material has been developed (for example, refer to Patent Document 2). . However, industrial gas turbine blades are large and the internal cooling structure is complex, so that grain boundaries are likely to occur during casting, and the yield of single crystal blades is significantly higher than that of aircraft engine gas turbine blades. Low. Therefore, it is desired to develop a Ni-base superalloy for unidirectionally solidified material having high strength.

  On the other hand, a normal cast material is used for a centrifugal charger of a turbocharger or a micro turbine in which unidirectional solidification in a direction parallel to the centrifugal stress direction is difficult (see, for example, Patent Document 3). This is because ordinary cast products cast with high-strength Ni-base superalloys developed on the premise of directional solidification have extremely low ductility due to low crystal grain boundary strength and cannot be applied to centrifugal wheels. is there.

USP 5069873 Specification USP 6051083 Specification USP 3720509 specification

  Alloys developed for single crystal materials, even those containing grain boundary strengthening elements, have lower strength at the time of grain boundary generation compared to very high intra-grain strength, and crystallized during casting. It is difficult to apply to large complex wings where grain boundaries are likely to occur.

  Alloys developed for unidirectionally solidified materials can improve the strength in the solidification direction by increasing the volume fraction of the γ 'phase, which is the precipitation strengthening phase, and increasing the amount of refractory metal elements such as W, Re, and Ta. On the other hand, the strength of the grain boundary is relatively lowered. An alloy having an increased strength in the solidification direction has a problem that the strength in the direction perpendicular to the solidification direction, that is, the strength of the crystal grain boundary is significantly reduced.

  The alloy shown in Patent Document 2 has sufficient grain boundary strength to improve the casting yield of a single crystal blade, but is applied to a unidirectionally solidified blade for an industrial gas turbine having a large complex shape. Therefore, the grain boundary strength is slightly insufficient. The alloy for unidirectionally solidified material disclosed in Patent Document 1 has almost sufficient grain boundary strength, but has low strength in the solidification direction.

  The alloy for ordinary casting material shown in Patent Document 3 has moderate ductility but low strength at high temperature, and for application to a turbocharger or a centrifugal wheel of a micro turbine, strength improvement corresponding to higher temperatures is required. Required.

  The object of the present invention is to obtain high high-temperature strength and excellent ductility in both unidirectionally solidified materials and ordinary cast materials, and to be applied to industrial gas turbines, centrifugal chargers for turbochargers or microturbines. It is to provide a suitable Ni-base superalloy.

  The first of the present invention is by weight, C: 0.06 to 0.3%, B: 0.01 to 0.05%, Hf: 0.2 to 3.0%, Co: 10.2 to 25%, Ta: 1 to 12%, Cr: 1.5 to 16%, Mo: 0 to 0.95%, W: 2 to 15%, Al: 3.5 to 6.5%, Re: 0.00. 5-9%, Nb: 0.2-2%, V: 0-1%, Zr: 0-0.02%, at least one selected from platinum group elements: 0-2%, selected from rare earth elements At least one selected from 0 to 2%, at least one selected from alkaline earth metals and Si: 0 to 0.1%, at least one selected from Fe, Ga and Ge: 0 to 5%, The balance is Ni-based superalloy made of Ni and inevitable impurities. This alloy has high high-temperature strength and excellent ductility regardless of whether it is a unidirectionally solidified material or a normal cast material. Note that the inevitable impurities refer to impurities mixed with the raw material, such as impurities contained in the Cr raw material, when adding Cr, for example. Examples of such impurities include Si, S, O, N, P, Mn, and Cu.

  The Ni-base superalloy having the above-described component composition has high high-temperature strength even if it is a single crystal material. Moreover, what contains 1.1 to 3.0% of Hf by weight% has very excellent ductility.

  In the second aspect of the present invention, by weight, C: 0.16 to 0.3%, B: 0.016 to 0.05%, Hf: 1.4 to 3.0%, Co: 10.2 to 25%, Ta: 1 to 4.9%, Cr: 1.5 to 8%, Mo: 0 to 0.95%, W: 7.2 to 15%, Al: 3.5 to 6.5%, Re: 1.1 to 9%, Nb: 0.2 to 2%, V: 0 to 1%, Zr: 0 to 0.02%, at least one platinum group element: 0 to 2%, rare earth element At least one kind: 0 to 2%, at least one kind of alkaline earth metal and Si: 0 to 0.1%, at least one kind selected from Fe, Ga and Ge: 0 to 5%, the balance being inevitable with Ni It is a Ni-base superalloy made of impurities. This alloy is particularly suitable for use with ordinary castings.

  In the Ni-base superalloy according to the second invention, when Cr is 1.5 to 7% by weight and W is 9 to 15%, the high temperature strength can be improved while maintaining excellent ductility. it can. In order to further improve the high-temperature strength, it is desirable to contain 1.5 to 7% Cr and 11.2 to 15% W by weight%. In the Ni-based superalloy for ordinary castings according to the second invention, C is 0.18 to 0.3%, Hf is 1.8 to 3.0%, and Cr is 1.5 to 7% by weight. , W is 11.2-15%, both strength and ductility are extremely excellent.

  The third of the present invention is weight percent, C: 0.06 to 0.3%, B: 0.01 to 0.05%, Hf: 1.4 to 3.0%, Co: 10.2 to 25%, Ta: 1 to 12%, Cr: 1.5 to 16%, Mo: 0 to 0.95%, W: 7.2 to 15%, Al: 3.5 to 6.5%, Re: 1.1-9%, Nb: 0.2-2%, V: 0-1%, Zr: 0-0.02%, at least one platinum group element: 0-2%, at least one rare earth element Species: 0 to 2%, alkaline earth metal and at least one of Si: 0 to 0.1%, at least one selected from Fe, Ga and Ge: 0 to 5%, the balance being Ni and inevitable impurities This is a Ni-base superalloy. This alloy is particularly suitable for use with unidirectionally solidified materials.

  In the Ni-base superalloy according to the third aspect of the present invention, those having Ta in the range of 1 to 6.5% and W in the range of 9 to 15% are excellent in high temperature strength and oxidation resistance. Those having Ta of 1 to 6.5% and W of 10.5 to 5% are particularly excellent in high temperature strength. In a unidirectionally solidified material, in order to further improve high-temperature strength while maintaining oxidation resistance, Ta is included in an amount of 1 to 4.9% and W is included in an amount of 11.2 to 15%. It is desirable. Further, it is extremely desirable that the total amount of Ta and W is 15 to 17% and the ratio of W / W + Ta is 0.6 to 0.8.

  The Ni-base superalloy of the present invention can have high high-temperature strength and ductility by performing solution heat treatment and aging heat treatment. The solvus temperature of the Ni-base superalloy of the present invention is 1240 ° C. or lower, and the partial melting temperature is 1260 ° C. or higher. Here, the solvus temperature is defined as a temperature at which the γ ′ phase, which is a precipitation strengthening phase, is dissolved in the γ phase in the dendrite core portion. The partial melting temperature is the temperature at which the eutectic part where elements with low melting points segregate during casting begins, and the solution heat treatment that maximizes the strength of the alloy is performed at a temperature above the solvus temperature and below the partial melting temperature. Is called. The interval between the solvus temperature and the partial melting temperature is preferably 20 ° C. or more, and the alloy of the present invention is also preferable from this point.

  The castings of the Ni-base superalloy of the present invention are suitable for use in axial-flow industrial gas turbine blades or centrifugal chargers or micro-turbine centrifugal wheels. Unidirectionally solidified castings cast by the unidirectional solidification method or single crystal castings cast by the unidirectional solidification method are particularly suitable for blades of large size and complicated shape among the blades of axial flow gas turbines. . Wings refer to moving or stationary blades. Using the Ni-base superalloy of the present invention, casting is performed by the selector method or the seed crystal method, and only the important part of the casting is made a single crystal, and there are grain boundaries in other relatively less important parts. It can also be done. Since the alloy of the present invention has both excellent solidification direction strength (strength in crystal grains) and grain boundary strength, the important part may be a single crystal and grain boundaries may exist in other parts. It is suitable for such applications.

  The alloy of the present invention is subjected to component adjustment as a master ingot, then divided into appropriate sizes and subjected to casting.

  When a centrifugal wheel for a turbocharger or a micro turbine is manufactured using the Ni-base superalloy of the present invention, the surface of the wing part is fine crystal, and the part from the wing part to the hub part is directed from the wing part to the hub part. It is desirable to cast so that the columnar crystals in the solidification direction and the hub portion are coarse crystals having a crystal grain size of 5 mm or more. Thereby, it is possible to manufacture a casting with low costs and few defects. In casting, it is preferable that the molten metal that is in contact with the solidification front face is always solidified to the pouring gate in all parts of the product portion. In addition, it is desirable to perform HIP treatment for 2 hours or more under conditions of a temperature of 1185 to 1285 ° C. and a pressure of 120 to 185 MPa after casting and before solution heat treatment, which can reduce defects. . The fine crystal means a crystal grain size of 1 mm or less, and includes chill crystals.

  According to the present invention, it was possible to obtain a Ni-base superalloy having high high-temperature strength and ductility regardless of whether it was a unidirectionally solidified material or a normal cast material. A gas turbine made of the Ni-base superalloy of the present invention and cast by the unidirectional solidification method is expected to greatly improve thermal efficiency at low cost. In addition, unidirectionally solidified material cannot be applied for reasons of shape, and a turbocharger or a micro-turbine centrifugal wheel manufactured from a normal cast material can achieve high strength and high temperature.

  The effects and appropriate contents of the individual elements in the Ni-base superalloy of the present invention will be described.

C forms Mf carbides with Hf, Ta, Nb, etc., M ₂₃ C ₆ and M ₆ C carbides with Cr, W, Mo, etc., and prevents crystal grain boundaries from moving at high temperatures. It is an element that has an effect of strengthening grain boundaries and plays a particularly important role in the present invention. In order to exert this effect, it is necessary to add at least 0.06% or more. In the case of a normal cast material, when higher grain boundary strength is required, 0.09% or more, more preferably 0.16% or more is added. When it is desired to increase both strength and ductility in a normal cast material, it is preferable to add 0.18% or more. When 0.18% or more of carbon is added, carbides are crystallized or precipitated at all crystal grain boundaries, preventing the grain boundaries from moving during HIP treatment or solution heat treatment and deformation, and high temperature strength In addition, there is a significant effect on ductility improvement. However, if the amount of C is excessively increased, elements effective for solid solution strengthening of the γ phase and the γ ′ phase are taken into the carbide, so that the high-temperature strength is lowered. Excess carbides also reduce fatigue strength. Therefore, the upper limit of C needs to be regulated to 0.3%.

  B has an effect of filling non-matching portions of the crystal grain boundaries and increasing the bond strength of the crystal grain boundaries. In the alloy of the present invention, at least 0.01% of B must be added. When a higher grain boundary strength is required as a normal cast material, it is desirable to add 0.016% or more. However, since B significantly lowers the melting point of the Ni-base superalloy, it is necessary to make it 0.05% at the maximum.

  Hf segregates at the grain boundaries and improves the ductility of the grain boundaries. Improvement in the solidification direction strength of the alloy is achieved by improving the strength in the crystal grains of the alloy. However, when the strength in the crystal grains of the alloy is improved and significantly exceeds the strength of the grain boundaries, the strength of the grain boundaries is relatively lowered, and the solidification perpendicular direction is perpendicular to the grain boundaries. The ductility is significantly reduced, and as a result, the strength in the direction perpendicular to solidification is reduced. In the case of a normal cast material, if the strength of the crystal grain boundary is low and the ductility is low, the strength as a casting is not improved no matter how much the crystal grains are strengthened. Hf is an essential element for preventing such a phenomenon, and it is preferably added at least 0.2% or more, particularly 0.5% or more. When considering ductility, it is recommended to add 1.1% or more in either a normal cast material or a unidirectional solidified material, and if high grain boundary strength is required, it should be 1.4% or more. Is good. When Hf is added in an amount of 1.8% or more, the area ratio of the eutectic structure is increased, and the effect of preventing the movement of the crystal grain boundary is increased like the carbide, and the strength of the crystal grain boundary is improved. This is particularly effective for ordinary cast materials. However, excessive addition lowers the melting point of the alloy in the same manner as B, so it is necessary to suppress the addition amount to 3.0% or less.

  Co has the effect of lowering the solid solution temperature of the γ 'phase and facilitating solution heat treatment. Especially when used in partial solution formation, as in the case of the alloy of the present invention, the solution rate can be reduced even at a low heat treatment temperature. It becomes possible to enlarge. In order to obtain the effect, addition of at least 10.2% is necessary. However, excessive addition of Co destabilizes the γ 'phase and rather leads to a decrease in strength. Therefore, Co must be at most 25%.

  Ta is a very effective element because it is superior to Ti and Nb as a solid solution strengthening element of the γ 'phase. In order to improve the strength in the direction of solidification, the larger the amount added, the better, and at least 1% addition is necessary. However, excessive addition deteriorates the phase stability of the alloy and leads to a decrease in strength. Therefore, it is necessary to regulate the maximum to 12%. Further, as a point of the present invention, it has been clarified that the ductility of the casting is improved by increasing the amount of W and relatively decreasing the amount of Ta. In order to improve the strength of the alloy, the larger the total amount of W and Ta, the better. On the other hand, when the total amount exceeds a certain level, the TCP phase precipitates and the strength is lowered. Therefore, in order to achieve high strength and high ductility, it is preferable to increase the W content and lower the Ta amount accordingly, and Ta should be 6.5% or less, especially 4.9% or less. Is preferred.

  W is an element that mainly strengthens the γ phase in solid solution, contrary to Ta. Therefore, in order to improve the strength in the direction of solidification, the larger the amount added, the better, and it is necessary to add at least 2%. Furthermore, as described above, in order to achieve both high strength and high ductility, a composition in which W is higher than Ta is appropriate, and W is 7.2% or more, more preferably 9% or more, and even more preferably 11. 5% or more is appropriate. However, like Ta, excessive addition of W deteriorates the phase stability of the alloy, leads to the precipitation of harmful phases such as the TCP phase, and remarkably lowers the corrosion resistance. Therefore, it is necessary to regulate the maximum to 15%.

  Since W and Ta have substantially the same mass number, the atomic% ratio and weight% ratio, which are important for expressing the characteristics of the alloy, are almost the same. When the ratio represented by W / W + Ta is in the range of 0.6 to 0.8, an alloy having particularly excellent strength and ductility was obtained. Further, when the total amount of Ta and W was 15 to 17%, very high strength was obtained.

  Mo has the same genus as W, and the effect on various properties of the Ni-base superalloy is almost the same as W. However, the present inventors have found that Mo significantly deteriorates the corrosion resistance in the combustion environment as compared with W. Therefore, it is 0 to 0.95% in the alloy of the present invention.

  Re, like W and Mo, is an element that mainly strengthens the γ phase by solid solution strengthening. Compared to Mo and W, it does not deteriorate the corrosion resistance in the combustion environment, so it is a very effective element to achieve both corrosion resistance and high-temperature strength. Can be strengthened. In order to obtain this effect, it is necessary to add at least 0.5%, preferably 1.1% or more. However, since Re has a remarkably low distribution rate to the γ ′ phase, it affects the phase stability. Therefore, it is necessary to keep the addition amount at 9% at the maximum.

Cr is an essential element for forming a protective film of Cr ₂ O ₃ and maintaining the corrosion resistance of the Ni-base superalloy. Therefore, at least 1.5% addition is necessary. However, excessive addition deteriorates the phase stability of the alloy as in the case of W and leads to precipitation of harmful phases such as the TCP phase, so it is necessary to regulate it to 16% or less. Further, when the high temperature strength is emphasized and the addition amount of W or Re needs to be increased, the addition amount of Cr is preferably 8% or less, and when the high temperature strength is more important, it is preferably 7% or less.

Al is an essential element for forming the γ ′ phase (Ni ₃ Al), and it is necessary to add at least 3.5% or more. When the volume fraction of the γ ′ phase is increased and importance is given to the strength in the solidification direction, it is preferably 5% or more. Moreover, Al improves oxidation resistance and corrosion resistance by forming an Al ₂ O ₃ protective film. However, if added excessively, the solid solution strengthening degree of the γ ′ phase is lowered and the high-temperature strength is lowered. Therefore, the addition amount needs to be 6.5% at the maximum.

  Nb is less effective than Ti, but has the effect of preventing the formation of a complex oxide of Cr and Al and improving the corrosion resistance of the alloy. On the other hand, the effect is smaller than that of Ta, but the effect of solid solution strengthening of the γ 'phase is higher than that of Ti. Therefore, Nb is an effective element that can improve the corrosion resistance without reducing the high-temperature strength, and it is necessary to add 0.2% or more. However, in order to maintain the phase stability of the γ ′ phase, the amount of Nb added needs to be 2% or less. When the corrosion resistance is particularly important, addition of 0.5% or more is preferable.

  Ti has the effect of preventing the formation of a complex oxide of Cr and Al and improving the corrosion resistance of the alloy. However, when both elements of Ta and Nb are added as in the present alloy system, if Ti is further added thereto, the stability of the γ ′ phase is inhibited. Also, when considering the balance between strength, corrosion resistance and oxidation resistance, it is better to add Nb or Ta than Ti. Furthermore, for alloy systems that contain grain boundary strengthening elements, such as the present invention alloy, and tend to lower the partial melting temperature of the alloy, the balance of the alloy elements is optimized and the partial melting temperature of the alloy is made as high as possible. Is effective for improving the solution heat treatment property, and as a result, the strength is improved. When Ta, Nb, and Ti are compared, the effect per atomic percent on the partial melting temperature reduction of the alloy is Ta <Nb <Ti, and addition of Ti is preferable in order to improve the partial melting temperature of the alloy. Absent. Therefore, Ti was not added in the alloy of the present invention. As described above, in the alloy of the present invention, it is possible to obtain excellent characteristics by controlling the balance between W entering the γ phase side and Ta entering the γ ′ phase side. Excellent characteristics were obtained where the ratio was small. Therefore, when Ti entering the γ ′ phase side is added, among the elements entering the γ ′ phase side, which is originally kept low, the amount of Ta added, which is most effective for improving the high temperature strength, must be further reduced. It will disappear. From this fact, it is important that Ti is not added in the alloy of the present invention.

  Conventionally, Zr has been considered to have an effect of improving the strength of the crystal grain boundary as in the case of Hf. However, according to the study by the present inventors, the effect on the improvement of the crystal grain boundary strength is remarkably higher than that of Hf. I found it small. Further, like the relationship between Ta, Nb and Ti described above, the effect of Zr on lowering the partial melting temperature is greater than Hf. Therefore, Zr, which has an effect on improving the grain boundary strength smaller than Hf and has an effect on lowering the partial melting temperature, which is larger than Hf, cannot be added to the alloy. When Zr is added, the amount of Hf, which is one of important elements in the alloy of the present invention, must be greatly reduced. Therefore, in the alloy of the present invention, Zr is restricted to 0 to 0.02% of the impurity level. As a result, the amount of Hf added can be increased, and the crystal grain boundary strength can be improved.

  When V is added, the solid solubility limit of Ta and Nb decreases, leading to a decrease in high temperature strength. Further, since the corrosion resistance is remarkably lowered, the alloy of the present invention is made 0 to 1% of the impurity level.

Rare earth elements such as Y improve the adhesion of the Al ₂ O ₃ protective film and greatly improve the oxidation resistance. However, since the melting point of the Ni-base superalloy is remarkably lowered, the addition amount is preferably 0 to 2%. The rare earth element is an element belonging to Group 3A of the periodic table, and includes Y, lanthanoids such as Sc and La, Ce, and actinoids such as Ac. The effects of these elements are almost the same, and even if one or a mixture of two or more kinds is added, the effects are almost the same. Therefore, the total amount of these elements is set to 0 to 2%.

  Alkaline earth metals and Si have the effect of improving the adhesion of the oxide film, but excessive addition reduces the ductility of the grain boundaries. Therefore, the total amount is preferably 0 to 0.1%.

  Platinum group elements such as Pt and Ru have the effect of extending the solid solution limit of W, Re, etc., which are effective elements for improving the high-temperature strength, but are very expensive elements, so the addition amount is 0 to 2%. And

  Fe, Ga and Ge have the effect of improving the adhesion of the oxide film. Ga and Ge have an effect of improving the high temperature strength by forming an intermetallic compound with Ni. However, excessive addition of both decreases the ductility of the grain boundaries. Therefore, the total amount is 0 to 5%.

  Hereinafter, the results of producing unidirectionally solidified castings and ordinary castings, cutting out test pieces, and measuring the strength and ductility will be described. Table 1 shows the component composition of the Ni-base superalloy used in the experiment. alloy1061 is an alloy included in the composition range of USP 6051083. Examples of the present invention are alloys 1064 to 1066, 1071 to 1073, 1077, 1079 to 1081, 1086 to 1104.

  The casting used for the evaluation was a flat plate of 100 mm × 15 mm × 230 mm, and one or both castings of normal casting (CC) and unidirectional solidification (DS) were produced. A master ingot adjusted in advance to the composition shown in Table 1 was used, the CC flat plate was cast by a general vacuum casting method, and the DS flat plate was cast by a mold drawing type unidirectional solidification method. After casting, solution heat treatment and aging heat treatment were performed, and then each test specimen for evaluation was collected by machining. The CC material was subjected to HIP treatment prior to solution heat treatment. The HIP treatment was performed in Ar for 4 hours under conditions of a temperature of 1200 ° C. and a pressure of 150 MPa. The standard solution heat treatment (ST) condition was that the temperature was above the solvus and below the partial melting temperature. After solution heat treatment, gas spray was rapidly cooled. Some of the alloys having a large temperature difference between the solvus temperature and the partial melting temperature have prepared test pieces having several types of ST temperatures. The aging heat treatment was a two-stage heat treatment in which all alloys were heated at 1080 ° C. for 4 hours, then rapidly cooled to room temperature, then heated at 871 ° C. for 20 hours, and then rapidly cooled to room temperature.

FIG. 1 shows the metal structure of the DS flat plate. In FIG. 1, the solidification direction, the grain boundary, and the direction perpendicular to solidification are shown. For the DS material, the creep rupture strength in the solidification (DS-L) direction and the solidification right angle (DS-T) direction was measured. The creep rupture strength in the DS-L direction was evaluated by measuring the creep rupture time under the conditions of 850 ° C.-40 kgf / mm ² or 1040 ° C.-14 kgf / mm ² . The creep rupture strength in the DS-T direction was evaluated by measuring the creep rupture time under the condition of 982 ° C.-14 kgf / mm ² . Since the CC flat plate is equiaxed, the sampling direction of the test piece is not particularly defined, and the test piece was collected from an arbitrary place. The creep rupture strength was evaluated by measuring the creep rupture time under the condition of 982 ° C.-14 kgf / mm ² . In addition, the test piece shape and conditions of the creep test and the tensile test were set to conform to ASTM or JIS standards. Corrosion resistance was evaluated by a burner rig test at 900 ° C. for both DS and CC materials. The superiority or inferiority of the corrosion resistance was determined by the length of time for the change in corrosion weight to reach 20 mg / cm ² . The test was performed for 10 hours per cycle, and the weight change was measured every cycle. A heavy oil containing 0.06 mass% of sulfur was used as the fuel, and a 0.1 mass% NaCl solution was sprayed into the combustion gas at 30 cc / min in order to accelerate corrosion. For evaluation of oxidation resistance, a flat plate having a length of 10 mm, a width of 15 mm, and a thickness of 3 mm was used for both the DS material and the CC material. These were heated in the atmosphere at 1100 ° C. per cycle for 100 hours, the amount of weight change was measured for each cycle, and the superiority or inferiority of oxidation resistance was judged by the magnitude of the absolute value. The test was performed for a maximum of 10 cycles for a total of 1000 hours.

  First, the evaluation results of the DS material will be described.

  Table 2 shows the creep rupture time in the DS-L direction of alloys 1061 to 1063. FIG. 2 shows the relationship between the creep rupture time in the DS-L direction and the Co content for these alloys. The longer the creep rupture time, the higher the creep rupture strength. Alloy 1061 was developed as a single crystal (SC) material, and when it is made of DS material, the creep rupture strength in the DS-L direction is low. As compared with alloy 1061, the one with the increased amount of Co has a high creep rupture strength in the DS-L direction. It was confirmed that the alloy 1063 containing about 10% Co has a creep rupture time four times or longer than the alloy 1061.

  Table 3 shows the creep rupture time in the DS-T direction of alloy 1063 and alloys 1064 to 1066 in which the amounts of Co and Hf were increased as compared to alloy 1063. FIG. 3 shows the relationship between the creep life in the DS-T direction and the Hf amount for these alloys. Although alloy 1063 has a high creep rupture strength in the DS-L direction, it has a low strength in the DS-T direction, that is, a grain boundary. It was found that the creep rupture strength in the DS-T direction was significantly improved by slightly increasing the Co amount and further increasing the Hf amount relative to the alloy 1063.

  As described above, the amount of Co and the amount of Hf are both increased with respect to the alloy 1061, and Co is set to 10.2% or more, and Hf is set to 0.5% or more, preferably 1.1% or more. It was confirmed that the creep rupture strength in the DS-T direction was high.

  Since it was found that the DS material of the Ni-base superalloy containing a large amount of Co and Hf has high strength in the DS-L direction and the DS-T direction, next, the amount of W and the amount of Ta were examined.

  W is an element that mainly enters the γ phase side, and conversely, Ta is an element that mainly enters the γ ′ phase side, which is a precipitated phase. An alloy having a large amount of W has a larger lattice constant on the γ phase side, and generally a smaller lattice constant mismatch defined by (lattice constant of γ ′ phase−lattice constant of γ phase) / (average of lattice constants of both phases). . Lattice mismatch is an important factor that greatly affects the deformation mechanism of Ni-base superalloys.

Table 4 shows that the alloy 1065 and the alloys 1071 to 1073 confirmed to have a high creep rupture strength in the DS-T direction from the results shown in Table 3 are 850 ° C.-40 kgf / mm ² creep rupture time in the DS-L direction. showed that. FIG. 4 shows the relationship between the ratio of W / W + Ta in% by weight and the 850 ° C.-40 kgf / mm ² creep rupture time in the DS-L direction for these alloys. Further, Table 5 shows 982 ° C.-14 kgf / mm ² creep rupture time in the DS-T direction for these alloys, and FIG. 5 shows the ratio of W / W + Ta and 982 ° C.-14 kgf / mm ^{2 in the} DS-T direction. The relationship of creep rupture time was shown.

As a result, it was found that the creep rupture strength in the DS-L direction and the creep rupture strength in the DS-T direction were improved as the ratio of W / W + Ta was increased. The creep rupture time 2654 hours at 850 ° C.-40 kgf / mm ² in the DS-L direction of alloy 1073 exceeds the creep rupture time 2470 hours of the single crystal (SC) material of alloy 1061 under the same conditions.

  From the above, the alloy in which the amount of Co and Hf is increased with respect to the alloy 1061 and the weight ratio of W / W + Ta is in the range of 0.6 to 0.8 is DS-equivalent to the SC material of the alloy 1061. It has been confirmed that it has excellent crystal grain boundary strength (DS-T direction strength) while having creep rupture strength in the L direction, and is suitable for a large wing having a complex shape. In particular, as is apparent from the creep characteristics of alloy 1073, when the W amount exceeds 11% and the Ta amount falls below 4%, the strength in both the DS-L direction and the DS-T direction is extremely high. It was found to have the following characteristics.

  Table 6 shows room temperature tensile elongations of alloy 1065 and alloy 1071 to 1073 in the DS-T direction. FIG. 6 shows the relationship between the room temperature tensile elongation in the DS-T direction and the ratio of W / W + Ta for these alloys. All the alloys had a room temperature tensile elongation of 3% or more, and were proved to have high ductility. It was found that the ductility at room temperature was most excellent when the ratio of W / W + Ta was around 0.65, and the weight ratio of W / W + Ta was preferably 0.6 to 0.8 from the viewpoint of improving ductility.

Next, the creep rupture strength in the DS-L direction and the DS-T direction was evaluated for several Ni-base superalloys having a weight ratio of W / W + Ta in the range of 0.6 to 0.8. The creep rupture strength in the DS-L direction was measured under two conditions of 850 ° C.-40 kgf / mm ² and 1040 ° C.-14 kgf / mm ² . Table 7 shows the evaluation results for alloys 1072, 1073, 1086 to 1098. FIG. 7 shows the relationship between the DS / L direction 1040 ° C.-14 kgf / mm ² creep rupture time and the weight ratio of W / W + Ta for these alloys. From the results shown in Table 7 and FIG. 7, an increase in strength was observed in all alloys except for alloys 1091 and 1097 in which the amount of Co was increased to 14% or more based on alloys 1072 and 1073. From this result, it can be seen that it is better to suppress the Co content to 14% or less. For alloys with an amount of Re in the 1.4% range, alloy 1093 with increased amounts of W and Ta compared to alloy 1073 has the highest creep rupture strength. However, considering long-term tissue stability, alloy 1092 in which W is increased from alloy 1073 is considered appropriate. Alloy 1090, 1096, 1098 with increased amount of Re, which is a remarkably expensive element compared with W or Ta, has extremely high creep rupture strength. Increasing the amount proved to be effective. In addition, with regard to Hf, since the creep rupture time of alloy 1088 in the DS-T direction is significantly shorter than that of other alloys, it is considered preferable to add up to 1.5% when emphasizing the strength in the DS-T direction. It is done.

  FIG. 16 shows the relationship between the Ta amount and W amount and the creep rupture time in the DS-L direction for alloys 1072, 1073, and 1086 to 1098. When the weight percentage ratio of W / W + Ta is 0.6 to 0.8, W amount is 10.5 to 15%, and Ta amount is 1 to 6.5%, the creep rupture strength is high. It can be seen that extremely high strength is obtained when the total amount is 15 to 17%.

  FIG. 8 shows the oxidation resistance test results of the alloy 1092 DS material and the comparative material. The oxidation resistance is improved as compared with either the single crystal (SC) material of alloy 1061 or the 14% Cr alloy (CC material) that has been proven in industrial gas turbines. FIG. 9 shows the corrosion resistance test results of the alloy 1092 DS material and the comparative material by the burner rig. Corrosion resistance is improved compared to any of SC material of alloy 1061, DS material shown in USP 5069873, and 14% Cr alloy (CC material) that has been proven in industrial gas turbines.

  Next, the evaluation result of CC material is demonstrated.

  CC materials are used for centrifugal wheels of turbochargers and microturbines. Table 8 shows the creep rupture test results for the alloy materials 1061, 1077, and 1079. FIG. 10 shows the relationship between the creep rupture time and the Hf amount for the CC materials of these alloys. It was confirmed that the creep rupture strength was high when the Hf content was increased based on the alloy 1061 CC material.

  Therefore, the influence of the C amount on the creep rupture strength was examined based on the alloy 1079 CC material in which the Hf amount was increased to improve the creep rupture strength. Table 9 shows the creep rupture time of CC materials of alloys 1079 to 1081 and 1201, and FIG. 11 shows the relationship between the creep rupture time and the C content. It was confirmed that the creep rupture strength was further improved by increasing the amount of C based on the alloy 1079.

Among the alloys shown in Table 9, the alloy 1081 with the highest C content of 0.2% was evaluated by increasing the ST temperature by 1260 ° C. and 20 ° C. As a result, the creep rupture time at 982 ° C.-14 kgf / mm ² was 2164 h. This corresponds to fix the 898 ° C. to ¹⁰⁵ hours tolerable temperature in stress 14 kgf / mm ^2. This service temperature is comparable to that of a so-called second generation DS alloy (group containing 3% Re) such as USP 5069873. Although it is a CC material, it has a strength comparable to that of the second-generation DS alloy. Therefore, a CC material made of an alloy comparable to the alloy 1081 can be used as a centrifugal wheel for a turbocharger or a microturbine for which a DS alloy is difficult to apply. It turned out to be very suitable. In addition, it has become possible to replace the blades of an axial turbine that previously used a DS material with a CC material, and it has been found that a significant effect can be obtained in terms of cost. As a result of evaluating the room temperature tensile ductility of the alloy 1081 at an ST temperature of 1260 ° C., the elongation at break at room temperature was 5.3%, and it was confirmed that the ductility was sufficient while having high high-temperature strength. The alloy 1081 has almost the same amount of elements except C and Hf as the alloy 1092 of the DS material. From this, high-temperature strength is achieved by optimizing the Co amount and W / Ta ratio as in the case of the alloy 1092, and further by optimizing the Hf amount and the C amount, the grain boundary strength is also improved. It is thought that the optimal alloy was obtained for the centrifugal wheel of the microturbine.

  FIG. 12 shows the oxidation resistance test results of the alloy 1081 CC material and the comparative material. Compared to any of the SC material of alloy 1061 and the 14% Cr alloy (CC material) that has been proven in industrial gas turbines, the oxidation resistance is improved. FIG. 13 shows the corrosion resistance test results of the alloy 1081 CC material and the comparative material burner rig. Compared to any of SC material of alloy 1061, DS material shown in US Pat. No. 5,069,873 and 14% Cr alloy (CC material) proven in industrial gas turbines, CC material of alloy 1081 has improved corrosion resistance.

Table 10 shows the creep rupture test results of 982 ° C.-14 kgf / mm ² for alloys having a composition close to alloy 1081, specifically, alloys 1099 to 1104. Moreover, the creep rupture test result of the alloy corresponding to USP 3720509 is shown for comparison. Among these alloys, alloy 1101 has the longest creep rupture time. It was found that the composition in which both the C content and the Hf content were increased was excellent in terms of creep rupture strength in the range of the C content of 0.15 to 0.2% and the Hf content of 1.50 to 2.04%. Because alloy 1104 has a lower amount of W than alloy 1099 to 1103 and ST temperature is lower, the creep rupture strength is lower than that of alloy 1099 to 1103, but as a result of measuring the tensile elongation at break at room temperature, it is as high as 6.4%. It turned out to be an excellent alloy when emphasizing the strength of the grain boundaries. The creep rupture strength of alloy 1104 is significantly higher than the alloy shown in USP 3720509, which is widely used in centrifugal chargers of turbochargers and microturbines.

FIG. 17 shows the relationship between Hf content and C content and 982 ° C.-14 kgf / mm ² creep rupture time for CC materials of alloys 1077, 1079, 1099 to 1103. When the Hf content is in the range of 1.4 to 3%, and the Hf content is greater than (-10C + 3.4) and less than (-10C + 4.4), extremely high creep rupture strength is obtained. It was.

As a result of casting a 150 kg master ingot with the composition of alloy 1092 and evaluating the castability of the large master ingot, it was confirmed that there was no problem. Further, using the master ingot, the industrial axial flow gas turbine rotor blade 1 having a total product length of 230 mm shown in FIG. 14 was cast by a mold drawing type unidirectional solidification method. As a result of inspecting the macrostructure and microstructure of this unidirectionally solidified blade, and further conducting fluorescence penetrant inspection and X-ray inspection, it was confirmed that there was no problem in the castability of this alloy. Further, another unidirectionally solidified blade was heated in vacuum at 1260 ° C. for 4 hours, and then subjected to solution heat treatment by quenching by blowing Ar gas, and then heated in the same vacuum at 1080 ° C. for 4 hours. Then, after quenching to room temperature, heating at 871 ° C. for 20 hours, and then quenching to room temperature, a two-stage aging treatment was performed. As a result, the creep rupture time in the DS-L direction under the condition of 850 ° C.-40 kgf / mm ² is 2400 hours, and a good result comparable to the creep rupture time 2470 hours of the single crystal (SC) specimen of alloy 1061 is obtained. Had.

Further, a 150 kg master ingot was cast with the composition of alloy 1101, and the castability of the large master ingot was evaluated, and it was confirmed that there was no problem. Further, using the master ingot, a centrifugal turbine for micro turbine having a blade maximum diameter of φ230 mm shown in FIG. 15 was cast by a normal casting method in a vacuum. As a result of inspecting the macro structure and micro structure of the wheel and further conducting the fluorescent penetrant inspection and X-ray inspection, it was confirmed that there was no problem in the castability of the alloy. In addition, another wheel was cast and subjected to HIP treatment in Ar at a temperature of 1200 ° C. and a pressure of 150 MPa for 4 hours, then heated in vacuum at 1240 ° C. for 4 hours, and then rapidly cooled by blowing Ar gas. Then, a solution heat treatment was performed, followed by heating in 1080 ° C. for 4 hours, followed by rapid cooling to room temperature, further heating at 871 ° C. for 20 hours, and then rapid cooling to room temperature. Test specimens were collected in the direction of centrifugal stress from the maximum diameter portion at the center of the blade, and the creep rupture strength was evaluated under the condition of 982 ° C.-14 kgf / mm ² .

For the alloy shown in US Pat. No. 3,720,509, a centrifugal wheel for a microturbine having the same shape was cast by a normal casting method in a vacuum. After casting, HIP treatment, solution heat treatment, and two-stage aging treatment were performed under the same conditions as in the examples of the present invention. Also from this wheel, a test piece was collected in the same manner as the wheel made of alloy 1101, and the creep rupture strength was evaluated under the condition of 982 ° C.-14 kgf / mm ² .

  As a result, the rupture time of the test piece collected from the wheel made of the alloy shown in USP 3720509 was 450 hours, whereas the rupture time of the test piece taken from the wheel made of alloy 1101 was 1800 hours, which is shown in USP 3720509. It has been confirmed that it is about 4 times longer than the alloy to be manufactured.

Represents the metal structure of the DS flat plate. The figure which showed the relationship between the creep rupture time of the solidification direction of DS material, and Co amount. The figure which showed the relationship between the creep rupture time of the solidification perpendicular direction of DS material, and Hf amount. The figure which showed the relationship between the creep rupture time of the solidification direction of DS material, and W / W + Ta ratio. The figure which showed the relationship between the creep rupture time of the solidification perpendicular direction of DS material, and W / W + Ta ratio. The figure which showed the relationship between room temperature tensile elongation of the solidification direction of DS material, and W / W + Ta ratio. The figure which showed the influence of W / W + Ta ratio which acts on the creep rupture time of the solidification direction of DS material. The figure which showed the oxidation test result of DS material and alloy material of alloy1092. The figure which showed the corrosion test result of DS material and alloy material of alloy1092. The figure which showed the relationship between the creep rupture time of CC material, and Hf amount. The figure which showed the relationship between the creep rupture time of CC material, and C amount. The figure which showed the oxidation test result of CC material and the comparative material of alloy1081. The figure which showed the corrosion test result of CC material and the comparative material of alloy1081. The external view of the industrial axial flow gas turbine rotor blade. The figure showing the cross-sectional metal structure of the centrifugal wheel for micro turbines. The figure which showed the relationship of the amount of Ta and W, and the creep rupture time of DS-L direction about DS material. The figure which showed the relationship between the amount of Hf and C, and the creep rupture time about CC material.

Explanation of symbols

  1 ... Industrial axial flow turbine blades.

Claims (24)

  By weight, C: 0.06-0.3%, B: 0.01-0.05%, Hf: 0.2-3.0%, Co: 10.2-25%, Ta: 1 12%, Cr: 1.5 to 16%, Mo: 0 to 0.95%, W: 2 to 15%, Al: 3.5 to 6.5%, Re: 0.5 to 9%, Nb: 0.2 to 2%, V: 0 to 1%, Zr: 0 to 0.02%, at least one platinum group element: 0 to 2%, at least one rare earth element: 0 to 2%, alkaline earth At least one selected from the group of metals and Si: 0 to 0.1%, at least one selected from Fe, Ga and Ge: 0 to 5%, the balance being made of Ni and inevitable impurities, high strength and high Ductile Ni-base superalloy.   The high-strength and high-ductility Ni-base superalloy according to claim 1, wherein Hf is included in an amount of 1.1 to 3.0% by weight.   A Ni-based superalloy having an equiaxed crystal structure manufactured by a normal casting method, and in terms of% by weight, C: 0.16-0.3%, B: 0.016-0.05%, Hf: 1 .4 to 3.0%, Co: 10.2 to 25%, Ta: 1 to 4.9%, Cr: 1.5 to 8%, Mo: 0 to 0.95%, W: 7.2 15%, Al: 3.5 to 6.5%, Re: 1.1 to 9%, Nb: 0.2 to 2%, V: 0 to 1%, Zr: 0 to 0.02%, platinum group At least one element: 0-2%, at least one rare earth element: 0-2%, at least one alkaline earth metal and Si: 0-0.1%, selected from Fe, Ga and Ge A high-strength, high-ductility Ni-base superalloy characterized by comprising at least one type: 0 to 5%, and the balance being Ni and inevitable impurities.   The high-strength, high-ductility Ni-base superalloy according to claim 3, which contains Cr: 1.5-7% and W: 9-15% by weight.   The high-strength, high-ductility Ni-base superalloy according to claim 3, which contains Cr: 1.5-7% and W: 11.2-15% by weight.   C: 0.18-0.3%, Hf: 1.8-3.0%, Cr: 1.5-7%, W: 11.2-15% The high strength and high ductility Ni-base superalloy according to claim 3.   The high strength and high ductility Ni-base superalloy according to claim 3, wherein the ratio of W / W + Ta is 0.6 to 0.8.   A Ni-base superalloy cast by a unidirectional solidification method, wherein C: 0.06-0.3%, B: 0.01-0.05%, Hf: 1.4-3. 0%, Co: 10.2 to 25%, Ta: 1 to 12%, Cr: 1.5 to 16%, Mo: 0 to 0.95%, W: 7.2 to 15%, Al: 3. 5 to 6.5%, Re: 1.1 to 9%, Nb: 0.2 to 2%, V: 0 to 1%, Zr: 0 to 0.02%, at least one platinum group element: 0 ~ 2%, at least one rare earth element: 0-2%, at least one alkaline earth metal and Si: 0-0.1%, at least one selected from Fe, Ga and Ge: 0-5 %, A high strength and high ductility Ni-base superalloy characterized in that the balance consists of Ni and inevitable impurities.   The high-strength and highly ductile Ni-base superalloy according to claim 8, characterized by containing Ta: 1 to 6.5% and W: 9 to 15% by weight.   The high-strength, high-ductility Ni-base superalloy according to claim 8, characterized by containing Ta: 1-6.5% and W: 10.5-15% by weight.   The high-strength and highly ductile Ni-based superalloy according to claim 8, characterized by containing Ta: 1 to 4.9% and W: 11.2 to 15% by weight.   The ratio of W / W + Ta is 0.6 to 0.8, and the high strength and high ductility Ni-base superalloy according to claim 8.   In weight%, Ta: 1 to 6.5%, W: 10.5 to 15%, the total amount of Ta and W is 15 to 17%, and the ratio of W / W + Ta is 0.6 to 0.8. The high-strength and high-ductility Ni-base superalloy according to claim 8.   In weight percent, Ta: 1 to 4.9%, W: 11.2 to 15%, the total amount of Ta and W is 15 to 17%, and the ratio of W / W + Ta is 0.6 to 0.8. The high-strength and high-ductility Ni-base superalloy according to claim 8.   A Ni-base superalloy casting having the composition according to claim 1.   The Ni-base superalloy casting according to claim 15, which is cast by a unidirectional solidification method or a normal casting method and has a single crystal structure, a columnar crystal structure, or an equiaxed crystal structure.   A master ingot for casting a Ni-base superalloy casting having the composition according to claim 1.   A centrifugal wheel for a turbocharger or a microturbine formed by a Ni-base superalloy casting having the composition according to claim 1.   The wing surface of the wheel is composed of fine crystals, the portion from the wing portion toward the hub portion is a columnar crystal in the solidification direction from the wing portion to the hub portion, and the hub portion is composed of coarse crystals having a crystal grain size of 5 mm or more. The centrifugal wheel according to claim 18.   20. The solidified form in which the molten metal in contact with the solidification front surface continues to the gate at the blade surface of the wheel, the portion from the blade portion toward the hub portion, and all portions of the hub portion. A centrifugal wheel as described in 1.   A Ni-base superalloy having the composition according to claim 1 is cast to produce a centrifugal wheel for a turbocharger or a microturbine, wherein the blade surface is a fine crystal and the portion from the blade portion toward the hub portion Is cast in a solidified columnar crystal from the wing portion toward the hub portion, and the hub portion is a coarse crystal having a crystal grain size of 5 mm or more. After casting, the temperature is 1185 to 1285 ° C. and the pressure is 100 to 185 MPa. A method for producing a centrifugal wheel, characterized by performing a HIP treatment for more than an hour and then performing a solution heat treatment.   A Ni-base superalloy having the composition according to claim 1 is cast to produce a centrifugal wheel for a turbocharger or a microturbine, wherein the blade surface is a fine crystal and the portion from the blade portion toward the hub portion Is a solidified columnar crystal in the solidification direction from the wing part to the hub part, the hub part is a coarse crystal with a crystal grain size of 5 mm or more, and the molten metal in contact with the solidification front surface in all parts of the product part continues to the pouring gate. A method for producing a centrifugal wheel, characterized in that after casting, HIP treatment is performed for 2 hours or more under conditions of a temperature of 1185 to 1285 ° C. and a pressure of 100 to 185 MPa, followed by solution heat treatment. .   A blade for an axial-flow gas turbine formed by a Ni-base superalloy casting having the composition according to claim 1.   The blade for an axial flow type gas turbine according to claim 23, which is cast by a unidirectional solidification method and has a single crystal structure and a columnar crystal structure.

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