JP2015042770A - HIGH-STRENGTH Ni-BASED ALLOY - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high-strength Ni-based alloy having both good corrosion resistance near room temperature, high strength, and high ductility. SOLUTION By mass%, C: less than 0.01%, Si: 0.5% or less, Mn: 0.5% or less, Cr: 15-25%, Mo alone or Mo as essential is Mo + (1/2 )? W: 1.0-5.0%, Al: 0.2-0.8%, Ti: 1.0-2.0%, Nb: 3.0-3.8%, Fe: 30% Hereinafter, Mg: 0.0007 to 0.01%, a value represented by Al / (Al + Ti-4C / 3) is 0.18 to 0.28, and a value represented by Mg / S is 0.7 or more. A high-strength Ni-based alloy consisting of the balance Ni and impurities. [Selection] Figure 1

Description

  The present invention relates to a Ni-based alloy that can be strengthened by precipitation, particularly by an aging treatment, and is used for applications requiring high strength near room temperature and corrosion resistance such as intergranular corrosion resistance.

A Ni-based alloy containing Al, Ti, and Nb is subjected to an aging treatment, whereby a γ ′ (gamma prime) phase and / or a γ ″ (gamma double prime) phase, which is a precipitation strengthening phase, in an austenite (γ) matrix. It is known that high tensile strength and creep strength can be obtained at room temperature and high temperature by precipitation strengthening by precipitation strengthening, especially γ ′ and / or γ ″ precipitation strengthening type Ni bases, which emphasize creep strength at high temperatures. In super heat-resistant alloys, it is important not only to strengthen the grain boundaries by fine precipitation of the precipitation strengthening phase, but also to strengthen the grain boundaries where creep deformation occurs. The grain boundaries are strengthened by precipitating carbides and intermetallic compounds at the grain boundaries. ing. In addition, trace impurities such as P and S that are harmful to the creep strength are easily segregated at the grain boundaries. In order to prevent these elements from segregating alone, a small amount of carbide is precipitated at the grain boundaries to solidify the precipitates. A measure such as removal from the grain boundary by melting is taken. For this reason, about 0.02% to 0.05% of C is added, and a carbide containing Cr, Ti, Nb, etc. is often precipitated in a minute amount in the grain boundaries and / or grains.
Further, since Ti and Nb are also main constituent elements of the γ ′ and / or γ ″ phase that precipitate during aging treatment and contribute to strengthening, if a large amount of primary carbides containing Ti and Nb are present, γ ′ during aging precipitation is present. And / or the precipitation amount of the γ ″ phase is reduced, and the strength may not be sufficiently obtained. For this reason, in the alloy which presupposes that a carbide | carbonized_material is precipitated, it adds a little rather in consideration of the amount of Ti and Nb consumed by a carbide | carbonized_material.

On the other hand, when emphasizing corrosion resistance, γ ′ and / or γ ″ precipitation-strengthened Ni-base superalloys are not only good in strength near room temperature but also good because Cr, Mo, etc. are added. It has corrosion resistance and is used in applications where high strength and corrosion resistance are required.Here, “around room temperature” means a temperature lower than the temperature at which creep occurs, for example, up to about 200 ° C. or 300 ° C. Shall be included.
Corrosion resistance is reduced by grain boundary precipitation of carbides containing Cr, as in stainless steel, so lower C is preferable, but general γ ′ and / or γ ″ precipitation strengthening used near room temperature Type Ni-based superalloys often contain about 0.02 to 0.05% C, and the amount of C is as low as stainless steel, and therefore often exhibits good corrosion resistance.

However, in a special application, for example, γ ′ and / or γ ″ precipitation strengthened Ni-base superalloy used in a nuclear power plant, for example, as disclosed in Patent Document 1, a Laves phase (M ₂ Nb ) Disappears and MC type carbides (M is Ti, Nb, etc.) and γ ″ phase are precipitated at the base of the austenite structure, and a method for manufacturing a reactor internal member that improves stress corrosion cracking resistance, and patent literature As disclosed in FIG. 2, the SCC resistance is improved by improving the stress corrosion cracking resistance (SCC resistance) by making the grain boundaries zigzag by precipitating M ₂₃ C ₆ type carbides consistent with austenite grains. Ni-based alloy members and heat treatment methods thereof have been proposed.

Japanese Patent Publication No. 61-28746 Japanese Patent Publication No. 4-42462

The alloy members shown in Patent Document 1 and Patent Document 2 try to improve the stress corrosion cracking resistance in a high-temperature and high-pressure water environment in a nuclear reactor by precipitating carbides by the manufacturing method and heat treatment method. Therefore, although there is no lower limit of C in the amount of C in the alloy composition, in the examples, the amount of C is 0.06% in Patent Document 1 and 0.025 to 0.05% in Patent Document 2. It is obvious that an alloy containing the amount of C is disclosed, and in order to exert the effects of the manufacturing methods shown in these patent documents, it is necessary to contain C necessary for carbide formation in the alloy. is there. This is also influenced by the assumption that stainless steel and Ni-based alloys such as high-temperature high-pressure water in reactors are supposed to be used in an environment where full-scale corrosion is not likely to occur, and it is considered to be acceptable to some extent. .
In addition, MC type carbides are generally formed in a solidified segregation part and are often unevenly distributed in a stringer shape, and MC type carbides have the effect of pinning austenite crystal grains, so MC type carbides are not uniform. There has been a problem that the austenite crystal grains are not uniform and the tensile strength and ductility are likely to vary with the distribution.
An object of the present invention is to provide a high-strength Ni-based alloy having both good corrosion resistance, high strength, and high ductility near room temperature.

Based on the alloy described in Patent Document 1, the present inventors studied to improve the corrosion resistance by further reducing the C content. On the other hand, free Ti, Nb, etc. increase due to the decrease in MC type carbides due to low C, so undesirable intermetallic compounds, η (eta) phase, δ (delta) phase, etc. Precipitation is likely to occur, the phase stability of the alloy may decrease, and hot workability and intergranular corrosion resistance decrease due to segregation of S at the grain boundaries caused by a decrease in grain boundary carbides due to low C. Therefore, the inventors studied diligently for a composition that eliminates these problems.
As a result, the inventors have found that it is effective to optimize the composition balance of Al, Ti, and C, and to add an appropriate amount of Mg as an S-fixing element.

That is, in the present invention, in mass%, C: less than 0.01%, Si: 0.5% or less, Mn: 0.5% or less, Cr: 15-25%, Mo alone or Mo is essential Mo + 0.5W : 1.0-5.0%, Al: 0.2-0.8%, Ti: 1.0-2.0%, Nb: 3.0-3.8%, Fe: 30% or less, Mg : 0.0007 to 0.01%, the value represented by Al / (Al + Ti-4C / 3) is 0.18 to 0.28, the value represented by Mg / S is 0.7 or more, and the balance It is a high-strength Ni-based alloy composed of Ni and impurities.
Further, preferably, by mass C: 0.008% or less, Si: 0.1% or less, Mn: 0.1% or less, Cr: 18-24%, Mo alone or Mo is essential Mo + 0.5W: 2.0 to 4.5%, Al: 0.35 to 0.7%, Ti: 1.2 to 1.8%, Nb: 3.2 to 3.8%, Fe: 20% or less, Mg: 0.001 to 0.005%, the value represented by Al / (Al + Ti-4C / 3) is 0.18 to 0.28, the value represented by Mg / S is 0.8 or more, and the balance Ni And a high-strength Ni-based alloy consisting of impurities.

  Since the Ni-based alloy of the present invention can achieve both good corrosion resistance near room temperature and high strength, it is used in chemical plants, drilling parts such as petroleum and natural gas, and seawater environments that are more corrosive than the environment in the reactor. When used for parts, etc., it provides higher reliability.

It is a figure which shows the relationship between 0.2% yield strength obtained by the tensile test done after the aging treatment, and Al / (Al + Ti-4C / 3). It is a figure which shows the relationship between the tensile strength obtained by the tensile test done after the aging treatment, and Al / (Al + Ti-4C / 3). It is a figure which shows the relationship between the elongation obtained by the tensile test done after the aging treatment, and Al / (Al + Ti-4C / 3). It is a figure which shows the relationship of the aperture_diaphragm | restriction obtained by the tension test done after the aging treatment, and Al / (Al + Ti-4C / 3).

First, each element prescribed | regulated by this invention and its content are demonstrated. Unless otherwise specified, the content is expressed as mass%.
C: Less than 0.01% C combines with Cr to form M ₂₃ C ₆ type carbide mainly at the austenite grain boundaries, not only lowering the overall corrosion and intergranular corrosion, but also bonding with Ti and Nb. As a result, stringer-like MC type carbides are produced, making the crystal grains non-uniform, resulting in non-uniform strength and ductility. M ₂₃ C ₆ type carbide containing Cr formed at the grain boundary suppresses the grain boundary sliding at high temperature and increases the high temperature strength and ductility. However, when used near room temperature, it is caused by precipitation at the grain boundary. Since the adverse effect of the decrease in corrosion resistance is greater, C needs to be lowered. Addition of 0.01% or more increases the corrosion resistance, so it is limited to less than 0.01%. C is preferably 0.008% or less, more preferably 0.006% or less, and further preferably 0.005% or less.

Si: 0.5% or less Si is used as a deoxidizer during alloy melting. However, when it contains excessively, ductility and workability will fall, It limits to 0.5% or less. A particularly preferable upper limit of Si is 0.1% or less, and more preferably 0.05% or less.
Mn: 0.5% or less Mn is used as a deoxidizing agent or a desulfurizing agent during alloy melting. If O or S is contained as an unavoidable impurity, it segregates at the grain boundary and lowers its melting point, thereby causing hot brittleness in which the grain boundary melts locally during hot working. Therefore, deoxidation and desulfurization using Mn I do. However, since ductility will fall when it contains excessively, it limits to 0.5% or less. The upper limit of preferable Mn is 0.1% or less, more preferably 0.05% or less.

Cr: 15-25%
Cr is an important element that improves the corrosion resistance by dissolving in the austenite matrix. However, if it is less than 15%, the above effect cannot be obtained, and excessive addition causes a decrease in the manufacturability and workability of the alloy, so it is limited to 15 to 25%. Preferably it is 18 to 24%, more preferably 18 to 23%.
Mo alone or Mo as essential Mo + 0.5W: 1.0-5.0%
Mo is an important element that improves the overall corrosion resistance and pitting corrosion resistance by dissolving in the austenite matrix together with Cr, and can be defined as Mo + 0.5 W by partially replacing it with equivalent W. If Mo alone or Mo is essential, the value of Mo + 0.5W is less than 1.0%, and the effect of improving corrosion resistance is small. On the other hand, if it exceeds 5.0%, the workability decreases. Essentially, the value of Mo + 0.5W is 1.0 to 5.0%. Preferably, it is 2.0 to 4.5%, more preferably 2.5 to 3.5%.

Al: 0.2-0.8%, Ti: 1.0-2.0%
Al and Ti form an intermetallic compound Ni ₃ (Al, Ti) called a γ ′ phase together with Ni, and are added to increase the high temperature strength of the alloy. If the content of Al is less than 0.2%, the above effect cannot be obtained, and excessive addition deteriorates the manufacturability and workability of the alloy, so Al is limited to 0.2 to 0.8%. Ti is highly effective in improving the strength around room temperature, and is added in a large amount when high strength is required. When Ti is less than 1.0%, a sufficient strength improvement effect cannot be obtained, while excessive addition does not contribute to the manufacturability and strengthening of the alloy between coarse metal plates called η phase (Ni ₃ Ti). Since the compound is formed in the vicinity of the austenite grain boundary to reduce the strength and ductility, Ti is limited to 1.0 to 2.0%. A preferable Al range is 0.35 to 0.7%, and a preferable Ti range is 1.2 to 1.8%. A more preferable range of Al is 0.35 to 0.6%, and a more preferable range of Ti is 1.3 to 1.7%.

Al / (Al + Ti-4C / 3): 0.18 to 0.28
As described above, when the amount of Ti increases, it becomes easier to form a brittle intermetallic compound called a coarse plate-like η phase (Ni ₃ Ti) in a layered manner, but this has a significant effect on the balance between the amounts of Al and Ti. . As for Ti, since it partially binds with C to form MC type carbide, the amount of Ti dissolved in the remaining austenite matrix after the formation of MC type carbide contributes to the formation of γ 'phase and η phase. . Nb also combines with C to form MC type carbides, so C in the alloy will fix both Ti and Nb. Considering that Ti and Nb combine with C to generate approximately the same amount of MC type carbide, and considering the atomic weight ratio of Ti, Nb, and C, the amount of Ti fixed as MC type carbide is 4C / 3. The amount of Ti remaining in the austenite matrix after being fixed as MC type carbide is expressed as Ti-4C / 3 in mass%, so the total amount of Al and Ti constituting the γ ′ phase is Al + Ti-4C / 3. Become. In order to improve the balance of strength and ductility of this alloy, it is important to balance Al and Ti. Therefore, a mass% ratio with Al and a numerical value of Al / (Al + Ti-4C / 3) are set. As a result of studying the range of this numerical value, if this value is less than 0.18, sufficient ductility cannot be obtained, while if it exceeds 0.28, sufficient strength cannot be obtained. Therefore, Al / (Al + Ti-4C / The value of 3) is limited to 0.18 to 0.28.
In the case of the alloy of the present invention, since the amount of Ti fixed to the MC type carbide is small in order to limit C to a low level, the influence of the MC type carbide is considered to be small, but that amount is about 0.02 to 0.05%. Compared with the alloy containing C, the amount of Ti forming the γ ′ phase is increased, and therefore, a brittle plate-like η phase tends to be easily formed in a layer form. In particular, when the aging temperature is increased to around 760 ° C. in order to increase the strength, a brittle plate-like η phase tends to be formed in a layer form at the austenite grain boundary. Therefore, it is difficult to apply a high temperature aging treatment at around 760 ° C. to obtain high strength, and it is difficult to obtain high strength. Therefore, it is important to optimize the amounts of Al and Ti and their balance to obtain high strength. On the other hand, Ti has a great effect of improving the strength around room temperature, so a large amount is added to obtain high strength. For this reason, in an alloy having a small amount of C, it is necessary to evaluate the phase stability by more accurately estimating the amount of Ti that forms the γ ′ phase. Therefore, by setting the above numerical values, it is possible to limit the component range that ensures more accurate phase stability.

Nb: 3.0 to 3.8%
Nb forms an intermetallic compound Ni ₃ Nb called γ ″ phase together with Ni, or is added to increase the high temperature strength of the alloy by forming a solid solution in the γ ′ phase (Ni ₃ (Al, Ti)) If Nb is less than 3.0%, sufficient strength cannot be obtained, while if it exceeds 3.8%, a δ phase tends to be formed, and sufficient ductility may not be obtained. The range of Nb is preferably 3.2 to 3.8%, and more preferably Nb is 3.3 to 3.8%.
Fe: 30% or less Fe is preferable to be low and is limited to 30% or less in order to reduce the overall corrosivity or to easily generate a δ phase or a Laves phase which are embrittled phases. Preferably it is 20% or less, more preferably 18% or less.

Mg: 0.0007 to 0.01%
When Mg is suppressed to less than 0.01%, the amount of precipitated carbide at the grain boundary becomes too small, so that S segregated at the grain boundary cannot be fixed, and hot due to S segregation at the grain boundary. Addition to improve hot workability and intergranular corrosion resistance by bonding with S that segregated at grain boundaries and fixing S, since workability and intergranular corrosion resistance are likely to decrease. . If Mg is less than 0.001%, the effect is not sufficient. On the other hand, if it exceeds 0.005%, the amount of oxides and sulfides increases, resulting in a decrease in cleanliness as inclusions, or with low melting point Ni. Mg is limited to 0.001 to 0.005% because the compound is increased and the hot workability is lowered. Preferably it is 0.001 to 0.004%, and more preferably 0.001 to 0.003%.
Since the purpose of adding Mg is to fix S that segregates at the grain boundaries, the amount of addition is determined according to the amount of S. In order to fix S effectively, Mg needs to be added in an atomic weight ratio of 1: 1 or more, so the value of Mg / S is limited to 0.7 or more. Preferably, 0.8 or more is good.
Ni (remainder)
The remaining Ni is an austenite generating element. Since the austenite phase is densely packed with atoms, the diffusion of atoms is slow even at high temperatures, and the high-temperature strength is higher than that of the ferrite phase. Also, the austenite base has a large solid solubility limit of the alloy element, and is advantageous for precipitation of the γ ′ phase and γ ″ phase, which are the key to precipitation strengthening, and for strengthening the austenite base itself by solid solution strengthening. Since the most effective element is Ni, in the present invention, the balance is Ni. Of course, impurities are included.

Among the impurities described above, the impurities to be particularly limited are as follows.
Impurities P and S are easily segregated at the grain boundaries, leading to deterioration of corrosion resistance and hot workability. Therefore, P is limited to 0.02% or less, and S is limited to less than 0.005%. S is preferably 0.003% or less, and more preferably 0.002% or less.
O and N combine with Al, Ti and the like to form oxide-based and nitride-based inclusions to reduce cleanliness and deteriorate corrosion resistance and fatigue strength. Since the amount of Al and Ti to be formed may be reduced to hinder the strength increase due to precipitation strengthening, it is preferable to keep it as low as possible. Therefore, preferable O is 0.008% or less, N is 0.004% or less, more preferable O is 0.006% or less, and N is 0.003% or less.
When Nb is added, a small amount of Ta may be mixed as an impurity. However, if Ta is in the range of 0.2% or less, there is little influence, and there is no need to limit it to a particularly low level. .
When Ni is added, a small amount of Co may be mixed as an impurity. However, if Co is in the range of 1% or less, the influence is small, and there is no need to limit it to a particularly low level, and it may be mixed.
B and Zr segregate at the grain boundaries to improve hot workability. However, if excessively added or mixed, a brittle compound is formed and hot workability is adversely affected. Therefore, B is 0.002% or less. Zr is limited to 0.05% or less.

  In order to obtain the above low impurity level by mass production, it is preferable to produce ingots by a combination of vacuum induction melting (VIM) and vacuum arc remelting (VAR). In this case, it is more preferable to produce an ingot by melting by a combination of vacuum induction melting (VIM) and electroslag remelting (ESR). Further, since S can be efficiently reduced by using ESR melting, it is preferable to employ ESR melting in the case of the alloy of the present invention in which S is desired to be limited to a low level.

A 10 kg ingot was prepared by vacuum induction melting. Table 1 shows the alloys No. 1 of the present invention produced. 1-8 and comparative alloy no. 21 to 25 chemical components are shown. As impurities not shown in Table 1, all of O (oxygen) in the alloys of the present invention were 0.008% or less, and N (nitrogen) was 0.002% or less.
The ingot shown in Table 1 was homogenized at 1180 ° C. for 20 h, and then hot forged to finish a bar material having a cross section of 25 mm × 40 mm. Thereafter, after holding at 1000 ° C. for 1 hour, an air-cooled solution treatment was performed, and further, an aging treatment under two conditions was performed.
The aging treatment conditions are as follows: hold at 720 ° C. for 8 hours, cool to 620 ° C. at 2 hours, hold at 620 ° C. for 8 hours, air cool (this condition is referred to as H1 heat treatment), hold at 760 ° C. for 8 hours, and at 2 hours It was cooled to 650 ° C., held at 620 ° C. for 8 hours, and then air-cooled conditions (this condition is referred to as H2 heat treatment).
After the aging treatment, a test piece was collected along the longitudinal direction of the bar material and subjected to a tensile test at room temperature. The tensile test piece was a round bar test piece having a parallel portion of 6.35 mm and a distance between gauge points of 25.4 mm, and was tested at room temperature in accordance with ASTM. The results are shown in Table 2 and FIG.

  From Table 2, it can be seen that both the alloy of the present invention and the comparative alloy exhibit high proof stress and tensile strength, and the alloy of the present invention also exhibits high elongation and drawing. On the other hand, the comparative alloy tends to exhibit low elongation and drawing as a whole compared with the alloy of the present invention, and in particular, comparative alloy No. with a low value of Al / (Al + Ti-4C / 3). Comparative Alloy No. 24, 25, and Nb are high. It can be seen that No. 21 is elongated and the aperture is low. Further, the alloy of the present invention having a low C had almost no stringer-like MC type carbide, but the comparative alloy No. 1 having a high C was used. 22 and 23 had many stringer-like MC type carbides. For this reason, for the convenience of collecting the test piece, the tensile test piece could be taken only in the longitudinal direction of the bar material, and therefore it is not reflected in the tensile test results in Table 2. 22 and 23 may have variations in ductility and strength in a direction perpendicular to the longitudinal direction of the bar material.

  By using the Ni-based alloy of the present invention, it is possible to obtain good corrosion resistance due to low C content, high strength near room temperature, and ductility. When used in parts for excavation such as plants, oil and natural gas, and parts used in the seawater environment, high reliability is achieved.

Claims (2)

  C: less than 0.01% by mass, Si: 0.5% or less, Mn: 0.5% or less, Cr: 15-25%, Mo alone or Mo as essential Mo + 0.5W: 1.0-5 0.0%, Al: 0.2-0.8%, Ti: 1.0-2.0%, Nb: 3.0-3.8%, Fe: 30% or less, Mg: 0.0007-0 0.01%, the value expressed by Al / (Al + Ti-4C / 3) is 0.18 to 0.28, the value expressed by Mg / S is 0.7 or more, and the balance is Ni and impurities. A high-strength Ni-base alloy characterized by The high-strength Ni-based alloy according to claim 1 is C: 0.008% or less, Si: 0.1% or less, Mn: 0.1% or less, Cr: 18-24%, Mo alone or Mo is essential Mo + 0.5W: 2.0-4.5%, Al: 0.35-0.7%, Ti: 1.2-1.8%, Nb: 3.2-3.8%, Fe: 20% or less, Mg: 0.001 to 0.005%, a value represented by Al / (Al + Ti-4C / 3) is 0.18 to 0.28, and a value represented by Mg / S is 0. A high-strength Ni-based alloy having a balance of Ni or more and a balance of Ni or more.