MANUFACTURING METHOD OF Ni-BASE SUPERALLOY AND HOT-PROCESSED MATERIAL

Abstract

To provide a manufacturing method of a Ni-base superalloy capable of accumulating a plastic strain while capable of maintaining a crystal grain to be small, and also to provide a hot-processed material.SOLUTION: A manufacturing method of a Ni-base superalloy includes: a raw material preparation step of preparing a hot-processing raw material of a Ni-base superalloy; a hot-processing step of heating the hot-processing raw material to a hot-processing preparation temperature T1 of 970-1005°C to perform hot-processing and obtaining a hot-processed material; and a cooling step of cooling the hot-processed material until a temperature thereof reaches 200°C or lower, after the temperature of the hot-processed material has lowered from a hot-processing finishing temperature T2 included within a range of 980-965°C to an intermediate temperature T3 included within a range of T2-500°C, under a condition of a cooling rate of the hot-processed material being 300°C / min. or faster, in a process of cooling the hot-processed material after the hot-processing step. A hot-processed material can be obtained through the manufacturing method.SELECTED DRAWING: Figure 1

Description

Technical Field

[0001] The present invention relates to a method for manufacturing a Ni-based superalloy and a hot-worked material.

Background Art

[0002] In aircraft jet engines and power generation gas turbines, the operating temperature tends to rise in order to improve fuel efficiency, and many components made of superalloys with excellent mechanical properties at high temperatures are used. And, for example, in rotating components such as turbine disks, 0.2% proof stress, tensile strength, fatigue characteristics, creep characteristics, etc. at high temperatures are required. These components are formed into near-net shapes of products by performing hot working such as forging on a material made of a superalloy.

[0003] As a superalloy used for the materials of these components, there is a Ni-based superalloy, and typically, Alloy 718, which is a Ni-based superalloy containing 50.0 to 55.0% by mass of Ni, can be mentioned. Further, as a method for improving the 0.2% proof stress of these components, it is effective to keep the crystal grains in the component structure fine. And it is effective to leave the plastic strain introduced by the above hot working even in the product state, and a Ni-based superalloy having a Grain Orientation Spread (GOS), which is the intragranular orientation difference parameter of the accumulated strain, of 0.7° or more has been proposed (Patent Document 1).

Prior Art Documents

Patent Documents

[0004]

Patent Document 1

Summary of the Invention

Problems to be Solved by the Invention

[0005] Accumulating strain inside a component made of a Ni-based superheat-resistant alloy (i.e., increasing GOS) is an effective method for maintaining a high 0.2% proof stress at high temperatures. Therefore, for a Ni-based superheat-resistant alloy after hot working (i.e., a hot-worked material), usually, a heat treatment consisting of a combination of solution treatment and aging treatment is performed to finish it into a component. Since solution treatment can promote the reduction of plastic strain accumulated in the hot-worked material due to microstructural changes such as recrystallization, Patent Document 1 recommends performing "direct aging treatment" that omits this solution treatment. By direct aging treatment, the plastic strain introduced and accumulated inside the hot-worked material by the above hot working can be maintained even after heat treatment. And since the crystal grains refined by the above hot working can also be maintained after heat treatment, a component with a high 0.2% proof stress at high temperatures can be obtained. However, in actual operation, there were cases where, despite performing hot working, plastic strain due to the above hot working was not accumulated inside the hot-worked material before direct aging treatment. Also, there were cases where, despite performing hot working, the crystal grains of the hot-worked material before direct aging treatment were large.

[0006] An object of the present invention is to provide a method for manufacturing a Ni-based superheat-resistant alloy capable of accumulating plastic strain and maintaining small crystal grains in a hot-worked material before heat treatment. And to provide a hot-worked material in which plastic strain is accumulated and crystal grains are small.

Means for Solving the Problems

[0007] That is, the present invention A material preparation step of preparing a material for hot working of a Ni-based superalloy having, by mass, C: 0.08% or less, Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, Cr: 17.0 - 21.0%, Ni: 50.0 - 55.0%, Co: 1.0% or less, Mo: 2.8 - 3.3%, Al: 0.20 - 0.80%, Ti: 0.65 - 1.15%, Nb: 4.75 - 5.50%, Ta: 2.0% or less, V: 1.0% or less, B: 0.006% or less, Zr: 0.10% or less, Mg: 0.005% or less, and the balance being Fe and inevitable impurities. A hot working step of heating the above-mentioned material for hot working to a hot working preparation temperature T1 of 970 - 1005 °C and performing hot working to obtain a hot worked material. In the process of cooling the hot worked material after the above-mentioned hot working step, after the temperature of this hot worked material drops from the hot working end temperature T2 included in the range of 980 - 965 °C to an intermediate temperature T3 included in the range of T2 - 500 °C, the hot worked material is cooled until the temperature of the hot worked material becomes 200 °C or less under the condition that the cooling rate of the hot worked material is 300 °C / min or more. It is a manufacturing method of a Ni-based superalloy having the above.

[0008] In addition, in the above-mentioned cooling step, it is preferable to cool under the condition that the cooling rate until the temperature of the hot worked material drops from T2 to T3 is 120 °C / min or less. Moreover, it is preferable to have a direct aging treatment step of reheating the hot worked material after the above-mentioned cooling step to 600 - 750 °C and performing aging treatment.

[0009] And, the present invention is A hot worked material to be subjected to the direct aging treatment step, The above hot-worked material is a Ni-based superalloy having, by mass%, C: 0.08% or less, Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, Cr: 17.0 to 21.0%, Ni: 50.0 to 55.0%, Co: 1.0% or less, Mo: 2.8 to 3.3%, Al: 0.20 to 0.80%, Ti: 0.65 to 1.15%, Nb: 4.75 to 5.50%, Ta: 2.0% or less, V: 1.0% or less, B: 0.006% or less, Zr: 0.10% or less, Mg: 0.005% or less, with the balance being Fe and inevitable impurities, It is a hot-worked material in which the in-grain orientation difference parameter Grain Orientation Spread (GOS) measured by the SEM-EBSD method is 0.70° or more and the average crystal grain size is 12 μm or less.

Advantages of the Invention

[0010] According to the present invention, it is possible to provide a hot-worked material in which plastic strain is accumulated and the crystal grains are small.

Brief Description of the Drawings

[0011]

Figure 1

Figure 2

Figure 3

Embodiments for Carrying Out the Invention

[0012] The inventor of the present invention has reached the present invention by finding that, as a method for maintaining a high 0.2% proof stress at high temperatures of Ni-based superalloys, specifically Alloy 718, it is effective to accumulate a sufficient amount of strain inside at the time before heat treatment (i.e., at the time of hot-worked material) and to keep the crystal grains fine.

[0013] <Composition> The reasons for limiting the chemical composition (mass%) of the Ni-based superalloy defined in the present invention are as follows. Note that the lower limit of each element indicated by "below" includes 0%.

[0014] ·C C forms MC carbides and M 23 C6 carbides in the alloy. The former has a pinning effect that suppresses the growth of crystal grains, and the latter improves the grain boundary strength by precipitating at the grain boundary. However, when the addition amount increases, coarse MC carbides are formed, which become fracture initiation points and reduce the fatigue properties. Therefore, the content of C is set to 0.08% or less. When C is contained to surely obtain the effects of the above C, it is preferable to set the lower limit of C to 0.01%. If the effects of the above carbides are not necessary, it may be not added.

[0015] ·Si, Mn, P, S Si, Mn, P, and S preferably have a low content because they reduce the grain boundary strength, and each may be 0%. However, when used for members of aircraft jet engines or power generation gas turbines, since sufficient strength can be obtained even when a certain amount is contained, Si can be tolerated in the range of 0.35% or less, Mn in the range of 0.35% or less, P in the range of 0.015% or less, and S in the range of 0.015% or less.

[0016] ·Cr Cr is an element effective for improving oxidation resistance and corrosion resistance in the use environment. Also, M 23Forming C6 carbides has the effect of enhancing the grain boundary strength. To exhibit these effects, 17.0% or more is required. On the other hand, when Cr is contained in excess, a brittle phase such as σ (sigma) phase is formed, which deteriorates the mechanical properties and hot workability. Therefore, the upper limit is set at 21.0%.

[0017] ·Ni Ni is a major element constituting the Ni-based superalloy of the present invention and is an element effective for obtaining excellent high-temperature strength. For example, due to the synergistic effect of a certain amount or more of Ni and other elements such as Al, Ti, and Nb like in the 718 alloy, excellent high-temperature strength can be obtained. To surely obtain the effect of Ni, it is preferable to contain 50.0% or more of Ni. And, since expensive Ni can be partly replaced by Fe described later, the upper limit can be set at 55.0%.

[0018] ·Co Co can improve the stability of the structure at high temperatures and obtain a high 0.2% proof stress. However, Co is an expensive element among the contained elements, and the content is set at 1.0% or less to reduce the alloy cost. Note that when the same effect as Co addition can be obtained by other elements other than Co, it can be not added.

[0019] ·Mo Mo contributes to the solid solution strengthening of the matrix and has the effect of improving the 0.2% proof stress at high temperatures. To exhibit this effect, the content of Mo should be 2.8% or more. However, when Mo becomes excessive, an intermetallic compound phase is formed and the strength is rather impaired, so the upper limit is set at 3.3%.

[0020] ·Al Al forms a γ' phase (gamma prime phase) which is a precipitation strengthening phase and is an element that improves the 0.2% proof stress. To obtain its effect, a minimum content of 0.20% is required, but excessive addition causes a large amount of γ' phase to precipitate and deteriorates the hot workability. Therefore, the upper limit is set at 0.80%.

[0021] ·Ti Ti, like Al, forms the γ' phase and is an element that improves the 0.2% proof stress. The effect can be obtained when the content is 0.65% or more. On the other hand, when added in excess, the η phase (eta phase), which is a brittle phase, precipitates, significantly deteriorating the hot workability and mechanical properties. Therefore, the upper limit is set at 1.15%.

[0022] ·Nb Nb, like Al or Ti, forms the γ' phase and is an element that solid-solution strengthens the γ' phase to enhance the high-temperature strength. Also, for example, in the 718 alloy, it forms the γ'' phase (gamma double prime phase), which is a precipitation strengthening phase, to increase the strength, and is used to control the grain size by forming the δ (delta) phase as pinning particles. Therefore, the lower limit of Nb is set at 4.75% to ensure these effects. However, excessive addition significantly impairs the hot workability, so the upper limit is set at 5.50%.

[0023] ·Ta Ta, like Al or Ti, forms the γ' phase and is an element that solid-solution strengthens the γ' phase to enhance the high-temperature strength. It also has a pinning effect of forming MC carbides to suppress the growth of grains. However, it is a very expensive element, so it is set at 2.0% or less to suppress the alloy cost. When Ta is contained to ensure the effects of Ta mentioned above, it is advisable to set the lower limit of Ta at 0.5%. Note that if the same effects as Ta addition can be obtained by other elements other than Ta, no addition is acceptable.

[0024] ·V V, like Ta, solid-solution strengthens the γ' phase to enhance the high-temperature strength. In addition, it forms MC carbides and is used for grain size control as pinning particles. However, excessive addition causes coarsening of the MC carbides, deteriorating the fatigue properties and hot workability, so it is set at 1.0% or less. When V is contained to ensure the effects of V mentioned above, it is advisable to set the lower limit of V at 0.5%. Note that if the same effects as V addition can be obtained by other elements other than V, no addition is acceptable.

[0025] ·B B is an element that improves the grain boundary strength and mainly improves the creep strength and ductility. On the other hand, B has a great effect of lowering the melting point, and excessive addition will conversely lower the grain boundary strength. Also, when coarse borides are formed, the hot workability decreases, so the upper limit is set at 0.006%. When B is contained to surely obtain the effect of the above-mentioned B, it is advisable to set the lower limit of B at 0.002%. Note that if the same effect as the addition of B can be obtained by other elements other than B, it can be not added.

[0026] ·Zr Zr, like B, improves the grain boundary strength, but excessive addition causes a decrease in the melting point and hot workability, so the upper limit is set at 0.10%. When Zr is contained to surely obtain the effect of the above-mentioned Zr, it is advisable to set the lower limit of Zr at 0.01%. Note that if the same effect as the addition of Zr can be obtained by other elements other than Zr, it can be not added.

[0027] ·Mg Mg has the effect of fixing S as a sulfide and improving the hot workability. However, excessive addition reduces the ductility, so it is set at 0.005% or less. When Mg is contained to surely obtain the effect of the above-mentioned Mg, it is advisable to set the lower limit of Mg at 0.0005%. Note that if the same effect as the addition of Mg can be obtained by other elements other than Mg, it can be not added.

[0028] ·Remainder The remainder is Fe and inevitable impurities. At this time, Cu may be considered as an inevitable impurity, but it may be contained as long as it is 0.30% or less. Fe is used as a substitute for expensive Ni and is effective in reducing the alloy cost. However, when Fe is contained in excess, embrittlement phases such as σ phase are formed, deteriorating the mechanical properties and hot workability. In this regard, in addition to containing 50.0% or more of Ni, the super heat-resistant alloy of the present invention can also contain a certain amount of other element species to address the above problems caused by excessive Fe content. The super heat-resistant alloy of the present invention can be referred to as a "Ni-based super heat-resistant alloy" in that it contains 50.0% or more of Ni. A Ni-based super heat-resistant alloy is a Ni-based alloy used in a high-temperature region, also referred to as a superalloy, a heat-resistant superalloy, or a superalloy, and is an alloy strengthened by a precipitation phase such as γ'. A typical alloy within the range of the alloying elements described above is Alloy 718.

[0029] <Stock Preparation Process> In obtaining the Ni-based heat-resistant alloy according to the present invention, first, a "stock" to be subjected to hot working is prepared. This stock has the component composition described above. And this stock may be an ingot obtained by casting molten metal, or one subjected to heat treatment such as soaking (sizing) on this. Also, it may be a billet or the like obtained by performing block processing on these ingots. And in the case of the present invention, considering that these stocks are formed into a near-net shape of the product by hot working described later, the shape of the stock is preferably one that has been processed into a rough shape by die forging or the like. By doing so, if the hot working described later is finish die forging, a member formed into a near-net shape can be efficiently obtained.

[0030] <Hot Working Process> As a method for improving the 0.2% proof stress of the Ni-based super heat-resistant alloy at high temperature, in the present invention, it is important to accumulate a sufficient amount of strain inside the Ni-based super heat-resistant alloy and to make the crystal grains fine at the time before heat treatment of this. For this purpose, in the present invention, the hot working performed on the above stock is utilized to introduce plastic strain into the inside of the stock during hot working, and further, to make the crystal grains in the structure fine.

[0031] That is, in the hot working process according to the present invention, first, a material having the above-described component composition is heated to a hot working preparation temperature "T1" of 970 to 1005 °C. At this time, in order to homogenize the entire material, the holding time at temperature T1 can be set to 1 to 6 hours. By setting temperature T1 to 970 °C or higher, sufficient hot workability can be ensured even when forming the material into a near-net shape by, for example, die forging during the subsequent hot working process. However, if temperature T1 is too high, the plastic strain introduced during hot working is likely to be consumed by recrystallization. And even if the crystal grains are refined by hot working, new crystal grains are likely to grow due to the above recrystallization. Therefore, temperature T1 is set to 1005 °C or lower. The preferable lower limit of temperature T1 is 980 °C, and the preferable upper limit is 1000 °C.

[0032] Note that the above temperature T1 does not necessarily have to be the temperature of the material when starting hot working. That is, considering the temperature drop that occurs during transportation after taking out the heated material from the heating furnace until it is placed in the die provided in the hot working apparatus, or the temperature drop due to heat absorption from the portion in contact with the die, the above temperature T1 can also be set "higher" compared to the temperature of the material during hot working.

[0033] Then, hot working is continuously performed on the material heated to the above temperature T1 to introduce plastic strain into the interior of the material during this hot working, and further refine the crystal grains in the structure. No special conditions are required for the hot working conditions at this time. However, it is desirable that plastic strain with a strain amount of 0.1 or more is introduced. Also, although the temperature of the material during hot working is considered to rise from its starting temperature due to heat generation during processing, even in this case, it is preferable to maintain it at 980 °C or lower. And the hot working end temperature "T2" of the material (that is, the hot worked material) when hot working is completed is within the range of 980 to 965 °C. By setting this temperature range, it is possible to maintain the hot workability of the material until the end of hot working while maintaining the plastic strain and refined crystal grains introduced during hot working. Note that it may be difficult to measure the exact temperature of the material during hot working at the portion in contact with the mold. In the present invention, the temperature of the material during hot working can be set as the maximum temperature among the portions where the temperature can be confirmed on the surface of the material.

[0034] <Cooling process> According to the hot working process of the present invention described above, plastic strain is accumulated inside the hot worked material immediately after hot working, and the crystal grains in the structure are also refined. According to the method of Patent Document 1, the hot worked material after this hot working is cooled from the above-mentioned hot working end temperature T2 to 900°C at a cooling rate of 20°C / min or more, and then air-cooled to room temperature, so that the plastic strain accumulated in the hot worked material can be reduced from being consumed by recrystallization or abnormal grain growth. In actual operation, it is realistic that some slow cooling proceeds during the period from when the hot worked material is transported from the hot working apparatus until cooling starts. If this slow cooling is, for example, by air cooling, a cooling rate of about 110°C / min can be assumed at the thin-walled part such as the end of the hot worked material, and a cooling rate of about 50°C / min can also be assumed at the thick-walled part such as the center of the hot worked material. Therefore, it is fully possible to achieve the above-mentioned cooling rate of 20°C / min or more. However, even in the cooled hot worked material obtained in this way, there may be cases where plastic strain due to hot working is not accumulated or the crystal grains are large.

[0035] That is, even when it is desired to quickly cool the hot worked material after hot working, the above-mentioned unintended slow cooling is difficult to avoid. Assuming that this hot worked material is a large one used in aircraft jet engines or power generation gas turbines, it takes time to transport this hot worked material, and when cooling of the hot worked material starts, it is considered that the temperature of this hot worked material has dropped significantly from the hot working end temperature T2. Then, at this point, the hot worked material has been in a high temperature state where recrystallization easily proceeds for a long time, and the recrystallized grains may have grown or the plastic strain may have decreased.

[0036] Therefore, the factors causing the crystal grains of the hot-worked material to grow during cooling were investigated. As a result, in the cooling immediately after the hot-working finish temperature T2, even if it is inevitable that the cooling rate decreases due to the unintended slow cooling that occurs during transportation or the like as described above, if the cooling rate remains slow thereafter, it was found that the time for the recrystallized grains to grow is given and the crystal grains grow larger. And it was found that the growth of these recrystallized grains also occurs in the temperature range of 900 °C or lower where air cooling was considered acceptable by the method of Patent Document 1. Therefore, in the present invention, in order not to give the time for the recrystallized grains to grow as much as possible in the above series of cooling processes, when the temperature at which cooling starts after finishing the transportation of the hot-worked material after hot working (that is, the temperature at which unintended slow cooling should end) is the intermediate temperature "T3", this temperature T3 is set to be within the range of T2 to 500 °C. And after the hot-worked material has dropped to the intermediate temperature T3, intentional cooling is started so that the cooling rate of the hot-worked material becomes 300 °C / min or more, and the hot-worked material is cooled under this condition until its temperature becomes 200 °C or lower.

[0037] First, considering that the factor for the growth of recrystallized grains depends not only on the high temperature but also greatly on the length of time, the reason when the plastic strain after cooling completion was not accumulated or the crystal grains were large in the method of Patent Document 1 is considered to be that the cooling from when it reached 900 °C was air-cooled to room temperature. Therefore, in the present invention, even after the temperature of the hot-worked material has dropped to the intermediate temperature T3 due to unintended slow cooling, it is effective to shorten the subsequent cooling time. And after the hot-worked material has dropped to the intermediate temperature T3, by quickly passing through the temperature range from this temperature to 200 °C or lower at a cooling rate of 300 °C / min or more, the time required for the growth of recrystallized grains can be shortened. Regarding the above temperature range, it preferably includes normal temperature or room temperature of 100 °C or lower, more preferably 50 °C or lower. And regarding the above cooling rate, it is preferably 400 °C / min or more, more preferably 500 °C / min or more, and still more preferably 600 °C / min or more. As a cooling means capable of achieving such a cooling rate, for example, gas cooling or water cooling can be used.

[0038] And by setting the cooling rate after the temperature of the hot-worked material has dropped to the intermediate temperature T3 to 300 °C / min or more, in the present invention, it is possible to lower the lower limit of this intermediate temperature T3 to 500 °C. However, preferably, the lower limit of the temperature T3 is 600 °C. More preferably, the lower limit of the temperature T3 is 700 °C, and even more preferably, the lower limit of the temperature T3 is 800 °C. Thereby, even when it takes time to transport the hot-worked material after the hot working is finished and the above-mentioned unintended slow cooling cannot be avoided (for example, even when the cooling rate during the period when the temperature of the hot-worked material drops from T2 to T3 is 120 °C / min or less), plastic strain can be left inside the cooled hot-worked material, and the crystal grains in the structure can also be kept small.

[0039] In addition, the cooling rate according to the present invention can be taken as the value obtained by dividing the difference (°C) between the temperature at the start of cooling and the temperature at the end of cooling of the hot-worked material by the time (min) required for the cooling. And the temperature of the hot-worked material used at this time can be, for example, the highest temperature among the temperatures that can be confirmed on its surface. At this time, the cooling rate obtained based on the internal temperature of the hot-worked material may be slower than the cooling rate obtained based on the above surface temperature. However, the value of the cooling rate of 300 °C / min or more in the present invention is a sufficiently large value that can suppress the growth of recrystallized grains even inside the hot-worked material by achieving this on the surface of the hot-worked material. And by surely cooling the hot-worked material to a sufficiently low temperature of 200 °C or less at which recrystallization cannot proceed, even if it is the surface temperature, the above-described effects of the present invention can be achieved.

[0040] <Direct aging treatment process> In the case of the Ni-based superheat-resistant alloy obtained by the above process, plastic strain is accumulated inside thereof, and the crystal grains in the structure are also maintained small. By performing heat treatment on this Ni-based superheat-resistant alloy, various members can be manufactured. And at this time, it is preferable to perform direct aging treatment omitting solution treatment. By direct aging treatment, the plastic strain accumulated inside the above Ni-based superheat-resistant alloy can be left even after heat treatment. And since the above small crystal grains can also be maintained after heat treatment, parts with high 0.2% proof stress at high temperature can be obtained. Regarding performing direct aging treatment, the conditions of this aging treatment itself can follow conventional methods. For example, aging treatment of holding the above Ni-based superheat-resistant alloy at 600 to 750 °C can be performed. And, for example, after performing the first-stage aging treatment of holding the above Ni-based superheat-resistant alloy at 700 to 750 °C for 2 to 20 hours, the second-stage aging treatment of holding it at 600 to 650 °C for 2 to 20 hours can be performed.

[0041] <Hot-worked material> According to the manufacturing method of the present invention described above, a hot-worked material with plastic strain accumulated inside and the crystal grains in the structure maintained small can be obtained. And, for example, it is a hot-worked material with a grain orientation spread (GOS) which is a parameter of intragranular orientation difference measured by the SEM-EBSD method of 0.70° or more and an average crystal grain size of 12 μm or less.

[0042] ·GOS As an important component of the hot-worked material of the present invention, there is GOS in the cross-sectional structure. GOS is generally measured by the SEM-EBSD method, calculates the orientation difference between points (pixels) constituting crystal grains, and obtains the value by averaging it. That is, it indirectly represents the magnitude of strain within the crystal grains. And when the value of this GOS is 0.70° or more, the plastic strain of the hot-worked material can be left even in the member after direct aging treatment, and the 0.2% proof stress at high temperature of the product can be maintained high. Regarding the value of this GOS, it is preferably 1.00° or more. Note that the upper limit of GOS is not particularly limited. However, for large and complex-shaped members such as those used in aircraft jet engines and power generation gas turbines, 8.00° is preferable as the upper limit of GOS for stably leaving plastic strain remaining. Considering ease of manufacturing, 5.00° is more preferable, and 3.00° is even more preferable.

[0043] · Average crystal grain size Furthermore, as an important component of the hot-worked material of the present invention, there is the size of crystal grains in the cross-sectional structure. The size of crystal grains can also be measured using the above SEM-EBSD method. That is, from the crystal grain boundary map obtained by EBSD, the cross-sectional area of each crystal grain is obtained using image analysis software. At this time, a large-angle grain boundary with an azimuth difference of 15° or more can be regarded as a crystal grain boundary. Then, by obtaining the equivalent circle diameter of each crystal grain from this cross-sectional area, the average crystal grain size of the entire crystal grains can be obtained. By setting the average crystal grain size of the crystal grains to 12 μm or less in terms of the equivalent circle diameter, after direct aging treatment, the yield strength is improved due to grain refinement strengthening. Preferably it is 11 μm or less, more preferably 10 μm or less, and even more preferably 9 μm or less. Note that the lower limit is not particularly limited. However, for large and complex-shaped members such as those used in aircraft jet engines and power generation gas turbines, the lower limit of the average crystal grain size for stably maintaining a small size is about 2 μm in terms of the equivalent circle diameter. Considering ease of manufacturing, the lower limit is more preferably 3 μm, and even more preferably 4 μm. An even more preferable lower limit is 5 μm.

Examples

[0044] <Material preparation process> Using a billet equivalent to Alloy 718, a material for hot working was produced. The chemical composition of the billet is shown in Table 1. The Fe content is approximately 17.5%. Si is 0.06%, Mn is 0.06%, P is 0.012%, S is 0.0002%, which are not shown in Table 1, and Co, Cu, Ta, V, Zr, and Mg are not added. Using this billet, upset forging and ring rolling were performed in the heating temperature range of 920 - 1010°C. After this ring rolling, a cylindrical material was cut out such that the radial direction of the ring mill became the length direction of the cylinder. Then, after heating this material to 980°C, it was machined into a so-called double-cone-shaped compression test piece (height 9 mm) for hot working, where the bottoms of two frustums of a cone (upper base diameter 4 mm, lower base diameter 10 mm, height 3.5 mm) were joined together with a disk (diameter 10 mm, height 2 mm) sandwiched in between.

[0045]

Table 1

[0046] <Hot working process> The above hot working material was heated to the hot working preparation temperature T1 of 980°C, and hot working was performed by compression from the height direction with a strain amount of approximately 0.4. Then, two types of hot worked materials were obtained, where the surface temperature of the hot worked material (hot working end temperature T2) at the end of this hot working was 980°C and 965°C.

[0047] <Cooling process> The two types of hot worked materials that had completed the above hot working process were slowly cooled by air cooling with a cooling rate of approximately 50°C / min until their surface temperatures dropped from the above hot working end temperatures T2 of 980°C and 965°C to their respective intermediate temperatures T3 (after approximately 10 minutes from the start of slow cooling, the surface temperature of the hot worked material would drop by 500°C). After the surface temperature of each of these hot-worked materials dropped to the respective intermediate temperature T3, the hot-worked materials were cooled with nitrogen gas under the condition that the cooling rate was about 600 °C / min, so that the surface temperature of these hot-worked materials was rapidly cooled until it reached room temperature, and samples were prepared. At this time, samples were also prepared in which the above-mentioned air cooling was continued even after the surface temperature of the hot-worked material dropped to the above-mentioned intermediate temperature T3, and the hot-worked material was slowly cooled until the surface temperature reached room temperature. A schematic flow of the series of processes from the above hot working process to the cooling process is shown in Fig. 1.

[0048] <Hot-worked material> Evaluation tests were conducted on the metallographic structures of the respective samples obtained according to the above-mentioned processes. The evaluation position of the sample was set to the position inside the test piece (strain amount 0.5). The metal cross-sectional structure was measured by the SEM-EBSD method, and the grain orientation spread (GOS), which is a parameter of the average azimuth difference between each measurement point in the grain and all points in the grain for each crystal grain, was analyzed. The measurement was performed in a field of view of 100 μm × 100 μm, and the GOS corresponding to each crystal grain in the measurement field of view was weighted by the area of the crystal grain, and the weighted value was taken as the representative value of the measurement field of view. Regarding the average grain size of the crystal grains, the equivalent circle diameters of all the crystal grains in the above-mentioned measurement field of view were obtained, and their average value was taken. Figs. 2 and 3 respectively show the results of the GOS and the average crystal grain size of each sample with respect to the intermediate temperature T3 when shifting to rapid cooling. And these specific values are as shown in Table 2 (temperature T2: 980 °C) and Table 3 (temperature T2: 965 °C) (the control of the intermediate temperature T3 was actually carried out by the slow cooling time until the hot-worked material reached the intermediate temperature T3).

[0049]

Table 2

[0050]

Table 3

[0051] As can be seen from Fig. 2, even if the hot-worked material immediately after the completion of hot working is slowly cooled, instead of ending the cooling while maintaining this slow cooling, it is shifted from this slow cooling to rapid cooling, and by increasing the intermediate temperature T3 at the time of this shift, it can be seen that the GOS after cooling is high and plastic strain can be maintained inside the sample. Similarly, as can be seen from Fig. 3, by increasing the above intermediate temperature T3, it can be seen that the average crystal grain size after cooling can also be maintained small. Therefore, by directly performing the aging treatment on such a sample, the plastic strain accumulated inside the sample can be left even after the heat treatment, and the crystal grains can also be maintained small, so that a component with a high 0.2% proof stress at high temperature can be obtained.

Claims

1. A material preparation step of preparing a material for hot working of a Ni-based superalloy having a component composition of, by mass%, C: 0.08% or less, Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, Cr: 17.0 to 21.0%, Ni: 50.0 to 55.0%, Co: 1.0% or less, Mo: 2.8 to 3.3%, Al: 0.20 to 0.80%, Ti: 0.65 to 1.15%, Nb: 4.75 to 5.50%, Ta: 2.0% or less, V: 1.0% or less, B: 0.006% or less, Zr: 0.10% or less, Mg: 0.005% or less, and the balance being Fe and inevitable impurities; A hot working step of heating the material for hot working to a hot working preparation temperature T1 of 970 to 1005 °C to perform hot working to obtain a hot worked material; A cooling step of cooling the hot worked material after the hot working step. After the temperature of the hot worked material drops from a hot working end temperature T2 included in the range of 980 to 965 °C to an intermediate temperature T3 included in the range of T2 to 500 °C, the hot worked material is cooled until the temperature of the hot worked material becomes 200 °C or less under the condition that the cooling rate is 300 °C / min or more; A method for manufacturing a Ni-based superalloy, characterized by comprising the above steps.

2. The method for manufacturing a Ni-based superalloy according to Claim 1, wherein the cooling step cools the hot worked material under the condition that the cooling rate from the temperature T2 to the temperature T3 of the hot worked material is 120 °C / min or less.

3. A direct aging treatment step of reheating the hot worked material after the cooling step to 600 to 750 °C to perform aging treatment; A method for manufacturing a Ni-based superalloy according to Claim 1 or 2, characterized by comprising the above step.

4. The hot worked material to be subjected to the direct aging treatment step, wherein the hot worked material is a Ni-based superalloy having a component composition of, by mass%, C: 0.08% or less, Si: 0.35% or less, Mn: 0.35% or less, P: 0.015% or less, S: 0.015% or less, Cr: 17.0 to 21.0%, Ni: 50.0 to 55.0%, Co: 1.0% or less, Mo: 2.8 to 3.3%, Al: 0.20 to 0.80%, Ti: 0.65 to 1.15%, Nb: 4.75 to 5.50%, Ta: 2.0% or less, V: 1.0% or less, B: 0.006% or less, Zr: 0.10% or less, Mg: 0.005% or less, and the balance being Fe and inevitable impurities; A hot-worked material, characterized in that the Grain Orientation Spread (GOS) of the intragranular orientation difference parameter measured by the SEM-EBSD method is 0.70° or more, and the average crystal grain size is 12 μm or less.

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