JP4230970B2 - Ni-base superalloys for unidirectional solidification with excellent solidification direction strength and grain boundary strength, castings and high-temperature parts for gas turbines - Google Patents

Description

  The present invention relates to a Ni-base superalloy, a Ni-base superalloy casting, and a gas turbine component. The Ni-base superalloy of the present invention is suitable for casting by a unidirectional solidification method, and is suitable for use in a moving blade or a stationary blade of a gas turbine.

  The combustion gas temperature of a gas turbine tends to increase year by year for the purpose of improving thermal efficiency, and each high temperature member of the gas turbine is required to have a material having excellent high temperature strength. Therefore, the material for moving blades, which is exposed to the harshest environment among the high temperature components of gas turbines, has changed from ordinary cast materials of Ni-based superalloys to unidirectionally solidified materials, and in gas turbines for aircraft engines, Furthermore, single crystal materials having higher high-temperature strength are being put into practical use.

  There exists a thing of patent document 1 as a Ni base superalloy developed for unidirectional solidification materials.

  On the other hand, alloys developed for single crystal materials generally do not contain grain boundary strengthening elements such as C, B, Hf, and Zr, and therefore cannot allow the presence of grain boundaries in the casting. Therefore, since it is large in size, it is difficult to put it into practical use for an industrial gas turbine in which crystal grain boundaries are likely to occur during casting. In order to overcome such problems, there are those described in Patent Documents 2 to 5 as Ni-based superalloys obtained by adding a small amount of grain boundary strengthening elements to a single crystal alloy. Among these, the alloy described in Patent Document 2 is shown as a single crystal material of a eutectic alloy. However, these Ni-base superalloys have lower solidification perpendicular direction strength, that is, grain boundary strength, compared to very high solidification direction strength. There is a problem that the world cannot be tolerated.

USP 5,069,873 Specification USP 4,849,030 specification USP 4,719,080 specification USP 4,878,965 specification USP 6,051,083 specification

  As described above, Ni-base superalloys have changed from ordinary cast materials of Ni-base superalloys to unidirectionally solidified materials and single-crystal materials in aircraft engine gas turbines. On the other hand, in industrial gas turbines, as disclosed in Patent Document 5, single crystal materials can be put to practical use by adding grain boundary strengthening elements such as C, B, Hf, and Zr to increase the casting yield. It has been.

  However, because industrial gas turbine blades are large and the internal cooling structure is complex, crystal grain boundaries are likely to occur during casting, and the yield of single crystal blades is significantly lower than that of aircraft engine gas turbine blades. . Therefore, if a Ni-based superalloy for high-strength unidirectional solidified material is developed, it is preferable to apply the unidirectional solidified material to large blades such as industrial gas turbine blades from the viewpoint of cost.

  The strength in the solidification direction of a unidirectional solidified material can be improved by increasing the volume fraction of the γ 'phase, which is a precipitation strengthening phase, and optimizing the balance of the amount of refractory metal elements such as W, Re, and Ta. . However, on the other hand, the strength of the grain boundary is relatively lowered, and the alloy whose strength in the solidification direction is increased, the strength in the direction perpendicular to the solidification direction, that is, the strength of the crystal grain boundary is significantly reduced. There is. The alloy shown in Patent Document 5 has sufficient grain boundary strength to improve the casting yield of the single crystal blade, but it is a unidirectionally solidified blade for industrial gas turbines with large and complex shapes. In order to apply, the grain boundary strength is slightly insufficient. Further, the conventional alloy for unidirectionally solidified material disclosed in Patent Document 1 has a problem that the grain boundary strength is sufficient, but the strength in the solidification direction is low.

  The object of the present invention is to provide a Ni-based superalloy for unidirectionally solidified material that has a high strength in the solidification direction and is excellent in strength in the direction perpendicular to the solidification direction, that is, the strength of the grain boundary. It is also an object to provide high temperature parts for castings and industrial gas turbines.

  The inventors of the present invention have an effect on the strength in the solidification direction of a unidirectional solidified material, the elements Cr, Mo, W and Re mainly strengthening the γ phase that is the parent phase of the Ni-base superalloy, and the γ that is the precipitation strengthened phase. 'The effects of the elements Ta, Ti and Nb that mainly strengthen the phase were studied. As a result, the composition range of additive elements that maximizes the solidification direction strength was clarified. Next, with respect to the composition range, an alloy added with C, B, Hf, and Zr, which is an element that strengthens the grain boundary, and an alloy with an increased amount of Co are produced without reducing the strength in the solidification direction. The range of the added amount of the grain boundary strengthening element and Co that can increase the strength of the grain boundary has been clarified. From the above, it has been clarified that the following composition range is appropriate as a composition range for solving the above problems.

  The present invention, by weight, C: 0.05-0.095%, B: 0.01-0.05%, Hf: 0.55-1.8%, Co: 1.5-16%, Ta: 4.1-12%, Cr: 1.5-16%, Mo: 0 to 3%, W: 2 to 12%, Ti: 0 to 1.1%, Al: 3.5 to 6.5%, Nb: 0 to 2%, V: 0 to 1%, Zr: 0 to 0.02%, Re : 0.01 to 9%, one or more platinum group elements: 0 to 5%, one or more rare earth elements: 0 to 2%, the balance being made of Ni and inevitable impurities Ni-base superalloy for unidirectional solidification with excellent solidification direction strength and grain boundary strength.

  Further, in order to obtain an alloy having a good balance between the solidification direction strength and the grain boundary strength, C: 0.05 to 0.095%, B: 0.01 to 0.03%, Hf: 0.55 to 1.15%, Co: 3-16%, Ta: 4.1-12%, Cr: 1.5-16%, Mo: 0-0.95%, W: 2-12%, Ti: 0-1.1%, Al: 3.5-6.5%, Nb: 0 ~ 2%, V: 0 to 1%, Zr: 0 to 0.02%, Re: 0.01 to 9%, one or more platinum group elements: 0 to 2%, one or more rare earth elements : A Ni-based superalloy consisting of 0 to 2%, the balance being Ni and inevitable impurities is suitable.

  In addition, when emphasizing corrosion resistance, the weight percentage is C: 0.05 to 0.095%, B: 0.01 to 0.03%, Hf: 0.55 to 1.15%, Co: 3 to 8.5%, Ta: 4.1 to 12%, Cr: 1.5 to 16%, Mo: 0 to 0.95%, W: 2 to 12%, Ti: 0.55 to 4%, Al: 3.5 to 6.5%, Nb: 0 to 2%, V: 0 to 1%, Zr: 0 ~ 0.02%, Re: 0.01 ~ 9%, 1 or more types of platinum group elements: 0 ~ 2%, 1 or more types of rare earth elements: 0 ~ 2%, the balance from Ni and inevitable impurities A Ni-base superalloy is preferable.

  When placing importance on the creep strength around 800-900 ° C, the weight percentage is C: 0.05-0.095%, B: 0.01-0.03%, Hf: 0.55-1.15%, Co: 9.5-16%, Ta: 4.1- 12%, Cr: 1.5-16%, Mo: 0-0.95%, W: 2-12%, Ti: 0-1.1%, Al: 3.5-6.5%, Nb: 0-2%, V: 0-1 %, Zr: 0 to 0.02%, Re: 0.01 to 9%, one or more platinum group elements: 0 to 2%, one or more rare earth elements: 0 to 2%, the balance being Ni And a Ni-base superalloy consisting of inevitable impurities is preferred.

  Further, the present invention is a Ni-based superalloy casting cast with a unidirectional solidification alloy having an excellent balance between the solidification direction strength and the grain boundary strength, and the alloy was cast by the unidirectional solidification method. It is suitable for a Ni-base superalloy unidirectionally solidified casting, and particularly suitable as a high temperature member for a gas turbine.

  Next, effects and appropriate ranges of individual elements will be described. In addition,% display has shown weight%.

C forms MC type carbides with Hf, Ta, Nb, Ti, etc., and forms M ₂₃ C ₆ and M ₆ C type carbides with Cr, W, Mo, etc., and inhibits the movement of grain boundaries at high temperatures. Strengthen grain boundaries. For this effect, it is necessary to add at least 0.05%. When the amount of addition of C increases, an element effective for solid solution strengthening of the γ phase and the γ ′ phase is taken into the carbide, so that the high temperature strength is lowered. Therefore, it is preferable to limit the upper limit of C to 0.095%.

  B has an effect of filling non-matching portions of the grain boundaries and increasing the bonding strength of the grain boundaries. In this alloy, at least 0.01% of B must be added. However, B must be 0.05% at the maximum in order to remarkably lower the melting point of the Ni-base superalloy, and B is preferably 0.03% or less in order to stabilize the solidification direction strength.

  Hf segregates at the grain boundaries and has the effect of improving the ductility of the grain boundaries, and is an element that plays a particularly important role in the present invention. 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. Hf is an essential element for preventing such a phenomenon, and at least 0.55% or more must be added. However, excessive addition lowers the melting point of the alloy in the same manner as B, so it should be regulated to 1.8% or less, more preferably 1.15% or less.

  Co has the effect of lowering the solid solution temperature of the γ 'phase and facilitating solution heat treatment. Especially when it is used in partial solution formation, as in the case of the present invention alloy, the solution rate is reduced even at a low heat treatment temperature. It becomes possible to enlarge. In order to obtain the effect, at least 1.5% addition is necessary. Furthermore, when importance is attached to the strength in the solidification direction, addition of 3% or more is effective. In addition, Co has the effect of increasing the volume fraction of the γ 'phase in the vicinity of 800 to 900 ° C, and as a result, it becomes clear that the creep strength near 800 to 900 ° C increases as the amount of Co added increases. It was. In order to obtain this effect, the amount of Co added is desirably 9.5% or more. However, excessive addition of Co destabilizes the γ 'phase and rather leads to a decrease in strength. Therefore, Co must be 16% or less at the maximum. Furthermore, since Co reduces corrosion resistance, when corrosion resistance is required, it is preferable to add the amount of 8.5% or less and to add 0.55 to 4% of Ti effective for corrosion resistance.

  Ta is the most effective element as a solid solution strengthening element of the γ 'phase. In order to improve the strength in the solidification direction, the larger the addition amount, the better, and the addition of at least 4.1% is necessary. However, excessive addition deteriorates the phase stability of the alloy and leads to a decrease in strength, so it is necessary to regulate it to 12% or less at the maximum.

  W is an element that mainly strengthens the γ phase in solid solution, contrary to Ta. In order to improve the strength in the solidification direction, the larger the amount added, the better. It is necessary to add at least 2%. However, excessive addition of W deteriorates the phase stability of the alloy, leads to precipitation of harmful phases such as TCP phase, and remarkably lowers the corrosion resistance. Therefore, it is necessary to regulate the maximum to 12%.

  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, in the alloy of the present invention, the addition amount of Mo is preferably 3% at the maximum, and when the corrosion resistance is important, it is preferably 0.95% or less.

  Re, like W and Mo, is an element that mainly strengthens the γ phase in solid solution. Compared to Mo or W, the effect of reducing the corrosion resistance in the combustion environment is small, so it is an extremely effective element for achieving both corrosion resistance and high temperature strength. By replacing it with the added amount of W or Mo, the corrosion resistance is improved and the alloy is effectively strengthened. However, since Re has a remarkably low distribution rate to the γ ′ phase, it tends to affect the phase stability. Therefore, the addition amount needs to be 9% or less at the maximum. On the other hand, Re is a very expensive element, so as an alloy suitable for a large-scale industrial gas turbine, it is preferable to minimize the addition amount, focus on cost, and effective strengthening by W. When obtained, the amount of Re added may be about 0.01%.

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, and in order to obtain the corrosion resistance required for industrial gas turbines, addition of 6% or more is desirable. However, excessive addition deteriorates the phase stability of the alloy and leads to precipitation of harmful phases such as the TCP phase, like W, so it must be regulated to 16% or less. In order to further improve the strength in the solidification direction, when it is necessary to increase the addition amount of W or Re, the addition amount of Cr is preferably 9% or less.

Al is an indispensable element for forming the γ 'phase (Ni ₃ Al). At least 3.5% or more is required. When the volume fraction of the γ' phase is increased and solidification direction strength is important It is preferable to add 5% or more. Al is also an essential element for improving 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 added amount needs to be 6.5% at the maximum.

  Ti has the effect of preventing the formation of a complex oxide of Cr and Al and improving the corrosion resistance of the alloy. Addition of 0.55% or more is effective when emphasizing corrosion resistance. However, if it is added excessively, the stability of the γ 'phase is hindered and the high temperature oxidation resistance is lowered. Ti, like Ta, dissolves mainly on the γ ′ phase side and strengthens the γ ′ phase by solid solution. However, since the effect of solid solution strengthening of the γ 'phase is greater in Ta, when the strength in the solidification direction is particularly important, the addition amount of Ti is preferably minimized, and the strength in the solidification direction is particularly If importance is attached, the amount of Ti added must be 1.1% at the maximum.

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

  Zr has been thought to have the effect of improving the strength of the grain boundaries, as in the case of Hf, but according to the study by the present inventors, the influence on the improvement of the grain boundary strength is significantly higher than that of Hf. I found it small. As the amount of Zr added increases, the amount of Hf added that lowers the initial melting temperature of the alloy as with Zr needs to be reduced. Therefore, in the alloy of the present invention, Zr is regulated to 0.02% or less of the impurity level. As a result, the amount of Hf added can be increased, and the 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 V significantly reduces the corrosion resistance, the addition amount is set to 1% or less of the impurity level in the alloy of the present invention.

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

  The platinum group element has the effect of extending the solid solution limit of an element effective for high-temperature strength such as W or Re in the alloy. However, since it is a very expensive element, the addition amount is preferably 2% or less. This platinum group element is preferably selected from Ru, Rh, Pd, Ir, and Pt, and one or more of them can be added.

  Inevitable impurities include Mg, Ca, Si, Ge, Fe, and the like. Of these, it is desirable that both Mg and Ca be regulated to 0.1% or less. Moreover, it is desirable that Si, Ge, and Fe are each suppressed to 5% or less.

According to the present invention, a novel unidirectional solidification Ni having a solidification direction strength equivalent to that of an expensive single crystal material and a grain boundary strength equivalent to that of a unidirectional solidification material alloy having a high yield and a low cost. A base superalloy is obtained. The gas turbine using the Ni-base superalloy according to the present invention is low in cost, greatly improves thermal efficiency, saves resources, and has environmental effects associated with reduction of CO ₂ and NO _X emissions, etc. The effect is great.

  Table 1 shows a comparison of the chemical composition and composition range of the Ni-base superalloy according to the present invention and the Ni-base superalloys described in Patent Documents 1 to 5.

  Hereinafter, the results obtained by casting a Ni-base superalloy by the unidirectional solidification method and measuring the high temperature strength in the solidification direction and in the direction perpendicular to the solidification direction will be described.

  Table 2 shows the composition, heat treatment conditions, and test results of the Ni-base superalloy according to the present invention and the Ni-base superalloy subjected to the experiment in the process of the present invention.

The cast used for the evaluation was a unidirectionally solidified flat plate of 100 mm × 15 mm × 230 mm, and was cast by a mold drawing type unidirectional solidification method using a master ingot prepared in advance in the composition shown in Table 2. After casting, heat treatment was performed under the conditions shown in Table 2, and then each test specimen for evaluation was collected by machining. FIG. 1 shows an external view (photograph) of a unidirectionally solidified flat plate made of YH64 after macro etching. FIG. 1 also shows the solidification direction, grain boundary, and solidification vertical direction, which are important for the unidirectional solidification flat plate. The creep rupture strength was evaluated under the conditions of 850 ° C.-40 kgf / mm ² or 1040 ° C.-14 kgf / mm ² for the solidification direction and 982 ° C.-14 kgf / mm ^{2 for} the direction perpendicular to solidification. Corrosion resistance was evaluated by a burner rig test at 900 ° C. The superiority or inferiority of the corrosion resistance was judged 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. The fuel used was heavy oil containing 0.06% by weight of sulfur, and a 1% by weight NaCl solution was sprayed into the combustion gas at 30cc / min for the purpose of accelerating corrosion.

The composition of the test alloy was selected for the purpose of examining the effect of each element based on YH161 included in the composition range of Patent Document 5. First, YH159 and YH160 were prepared and evaluated for the purpose of examining the effect of Hf. YH159 and YH160 correspond to a composition in which Hf is added alone to YH161 while keeping the addition amount of other elements substantially constant. By adding Hf that lowers the melting point of the alloy, the initial melting temperature of the alloy decreased, and the solution heat treatment temperature had to be lowered from 1280 ° C to 1250 ° C in YH161, so 1040 ° C-14kgf / mm ² The creep rupture time of YH161 also decreased significantly from 615 hours of YH161 to 152 hours of YH159 and 113 hours of YH160. From this, the simple addition of Hf to YH161 can be expected to improve the strength in the direction perpendicular to solidification, but the strength in the solidification direction, which is important for gas turbine members, especially moving blades, is greatly reduced, and is not practically appropriate It became clear.

Next, YH62 and YH63 were produced, and the effect of adding Co to YH161 was examined. By adding Co, while the same solution treatment temperature as YH161 is 1280 ° C., the solution rate, that is, the γ ′ phase coarsely precipitated during solidification is once dissolved in the γ phase while maintaining the solution treatment temperature, The area ratio of the region that precipitated finely during cooling and contributed to the improvement of the high temperature strength increased. Generally, as the solution rate increases, the creep strength in the solidification direction improves. For YH62 and YH63, with an increase in Co content, but the solution rate was increased, the 1040 ℃ -14kgf / mm ² creep rupture time, from 615 hours of YH161, 522 hours of YH62, 428 hours of YH62 Slightly declined. This is probably because Co lowers the creep strength at high temperatures. On the other hand, the creep rupture time in the direction perpendicular to solidification was slightly improved to 26 hours for YH62 and 54 hours for YH62 compared to 20 hours for YH161. The effect of single addition was very small.

Next, in order to evaluate the influence of the combined addition of Co and Hf, YH64 to 67, 74, 75, 162 to 166, 172 and 173 were further prepared and evaluated. FIGS. 2 and 3 show the effect of Co content on the creep rupture time at 1040 ° C.-14 kgf / mm ² when the Hf content is constant at about 1% and about 1.5%. From these figures, it can be seen that the addition of about 5 to 10% Co is effective in improving the creep rupture time at 1040 ° C.-14 kgf / mm ² in both cases. At the same time, when the amount of Co is about 1%, the solution heat treatment temperature cannot be increased, so the creep rupture time is short, and even when the amount of Co exceeds about 10%, the creep rupture time of 1040 ° C-14kgf / mm ² Tended to decrease.

On the other hand, FIG. 4 shows the effect of Co amount on the steady creep rate at 850 ° C. and 40 kgf / mm ² when the Hf amount is about 1%. From this figure, in accordance with the amount of Co is increased, the steady creep rate of 850 ℃ -40kgf / mm ² is decreased tendency is observed, it is clearly in the creep strength improving 850 ° C. is effective increase in the Co amount became.

FIG. 5 shows the effect of Hf on the creep rupture time of 1040 ° C.-14 kgf / mm ² when the Co content is about 9.6%. From this figure, it can be seen that when the Co content is about 9.6%, the creep rupture time at 1040 ° C.-14 kgf / mm ² is maximum in the region where the Hf content is about 0.5 to 1.0%. The short creep rupture time of YH63 is considered to mean that the grain boundary strength must be improved by adding Hf even in the solidification direction. On the other hand, the creep rupture time of the alloy with a higher Hf addition amount than YH164 was reduced because the addition of Hf reduced the initial melting temperature of the alloy and required a lower solution heat treatment temperature. This is thought to be due to a drop in

FIG. 6 shows the effect of Hf on the creep rupture time of 982 ° C.-14 kgf / mm ^{2 in} the direction perpendicular to solidification when the Co content is about 9.6%. It can be seen from this figure that the creep rupture strength in the direction perpendicular to solidification increases as the amount of Hf increases.

  From the above results, when the amount of Co is as low as about 1.0%, it is necessary to lower the solution heat treatment temperature by increasing the amount of Hf added, the creep rupture strength in the solidification direction is reduced, but the amount of Co is about 5 to 10 In the% range, it is not necessary to lower the solution heat treatment temperature even if the Hf addition amount is increased to about 1%. By adding Co and Hf in combination, the creep strength in both the solidification direction and the direction perpendicular to the solidification direction of the unidirectional solidification material It became clear that can be improved.

  Next, the effect of Zr, an element belonging to the same group as Hf, was examined. FIG. 7 shows the effect of Zr on the creep strength in the direction perpendicular to solidification when the Co content is about 9.7%. From this figure, it was found that Zr maximizes the creep strength in the direction perpendicular to solidification, that is, the strength of the grain boundary when the addition amount is about 0.06%. However, the effect was remarkably small compared with Hf, and the addition of 0.1% or more reduced the creep strength in the direction perpendicular to solidification. In addition, the addition of Zr also has an adverse effect that it becomes impossible to increase the amount of Hf effective for improving the grain boundary strength.

  Next, the amount of C which is a grain boundary strengthening element was examined. FIG. 8 shows the effect of C on the creep strength in the solidification direction when the Co content is about 9.6% and the Hf content is about 1.0%. When the amount of C is about 0.05% or less, it is considered that the creep rupture time in the solidification direction is short because the grain boundary strength is insufficient. After that, the creep rupture time reaches a maximum at a C content of about 0.07%, and then the creep rupture time decreases because an element effective for solid solution strengthening such as Ta is consumed for carbide formation due to an increase in the C content. Possible cause. FIG. 9 shows the effect of C on the creep strength in the direction perpendicular to solidification, that is, the grain boundary strength when the Co content is about 9.6% and the Hf content is about 1.0%. When the C content was about 0.05% or less, the creep rupture time was shortened because the grain boundary strength was insufficient. Thereafter, the creep rupture time decreased after the creep rupture time reached a maximum at a C content of about 0.07%. This is presumably because the elements effective for solid solution strengthening, such as Ta, were consumed for carbide formation due to the increase in the C content, as in the solidification direction.

Next, the effect of Ti on corrosion resistance was examined. FIG. 10 shows the effect of Ti on the corrosion resistance. The vertical axis in FIG. 10 shows the time for the change in corrosion weight to reach 20 mg / cm ² in the burner rig test at 900 ° C., and the longer this time is, the better the corrosion resistance is. As a result, the corrosion resistance improved as the Ti content increased, and YH178 and YH179, in which the Co content was reduced to 8% or less, showed excellent corrosion resistance. Although these alloys were not subjected to strength evaluation, YH176 to 179 with Ti content exceeding 1% showed a decrease in solution heat treatment temperature, which is considered not preferable in terms of high-temperature strength.

YH163 is selected from the above experimental alloys, and a solid unidirectionally solidified gas turbine blade 1 (total length: 170 mm) shown in FIG. 11 is cast, followed by a solution heat treatment temperature of 1280 ° C., followed by 1080 ° C./4 hours. And creep specimens were collected after aging treatment at 871 ° C./20 hours. This solid strength evaluation test piece shows a fracture time of 521 hours, which is almost equivalent to a test piece taken from a unidirectionally solidified flat plate in a creep test under the condition of 1040 ° C-14kgf / mm ² , and there is no problem with the solid strength characteristics. Was confirmed. This is sufficiently longer than the creep rupture time of 197 hours under the same conditions of alloy YH186 included in the composition range shown in Patent Document 1.

  According to the present invention, a Ni-based superalloy for unidirectional solidification has been developed in which the solidification direction strength and the grain boundary strength are balanced. Thereby, it became possible to manufacture the moving blade and stationary blade of an industrial gas turbine with a unidirectionally solidified member.

The figure which shows the external appearance after macro-etching of the casting by one Example of this invention cast by the unidirectional solidification method. The diagram which showed the relationship between Co amount and creep rupture time about Ni base superalloy whose Hf amount is 1.0 weight%. The diagram which showed the relationship between Co amount and creep rupture time about Ni base superalloy whose Hf amount is 1.5 weight%. The diagram which showed the relationship between Co amount and steady state creep rate about Ni base superalloy whose Hf amount is 1.0 weight%. The diagram which showed the relationship between the amount of Hf and the creep rupture time of a solidification direction about the directionally solidified Ni base superalloy whose Co amount is 9.6 weight%. A diagram showing the relationship between Hf content and creep rupture time in the direction perpendicular to solidification for a unidirectionally solidified Ni-base superalloy having a Co content of 9.6% by weight. A diagram showing the relationship between the Zr content and the creep rupture time in the direction perpendicular to solidification for a unidirectionally solidified Ni-base superalloy with a Co content of 9.7 wt%. The diagram which showed the relationship between the amount of C and the creep rupture time of a solidification direction about the unidirectionally solidified Ni base superalloy whose Co amount is 9.6 weight% and Hf amount is 1.0 weight%. A diagram showing the relationship between C content and creep rupture time in the direction perpendicular to solidification for a unidirectionally solidified Ni-base superalloy having a Co content of 9.6 wt% and an Hf content of 1.0 wt%. The diagram which showed the relationship between the amount of Ti and the time for the amount of change in corrosion weight to reach 20 mg / cm ² for the unidirectionally solidified Ni-base superalloy. The perspective view of a gas turbine rotor blade.

Explanation of symbols

  1 ... Gas turbine blade.

Claims (4)

In weight%, C: 0.05~0.095%, B : 0.01~0.05%, Hf: 0.55~1.8%, Co: 9.5 ~16%, Ta: 4.1~12%, Cr: 1.5~16%, Mo: 0~ 3%, W: 2-12%, Al: 3.5-6.5%, Nb: 0.5-2 %, V: 0-1%, Zr: 0-0.02%, Re: 0.01-9%, platinum group element 1 Species or 2 or more: 0 to 5%, 1 or 2 or more of rare earth elements: 0 to 2%, the balance is composed of Ni and inevitable impurities. Excellent solidification direction strength and grain boundary strength Ni-base superalloy for unidirectional solidification. In weight%, C: 0.05~0.095%, B : 0.01~0.03%, Hf: 0.55~1.15%, Co: 9.5 ~16%, Ta: 4.1~12%, Cr: 1.5~16%, Mo: 0~ 0.95%, W: 2-12%, Al: 3.5-6.5%, Nb: 0.5-2 %, V: 0-1%, Zr: 0-0.02%, Re: 0.01-9%, one of the platinum group elements Species or two or more: 0 to 2%, one or more rare earth elements: 0 to 2%, the balance consists of Ni and inevitable impurities, excellent solidification direction strength and grain boundary strength Ni-base superalloy for unidirectional solidification. A Ni-base superalloy unidirectionally solidified casting cast by the unidirectional solidification method using the Ni-base superalloy according to claim 1 or 2. A high-temperature component for a gas turbine comprising a unidirectionally solidified cast of a Ni-base superalloy cast by the unidirectional solidification method using the Ni-base superalloy according to claim 1 or 2.

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