JP6705508B2 - NiCrFe alloy - Google Patents

Description

The present invention relates to an austenitic heat resistant alloy, and more particularly to a NiCrFe alloy.

Conventionally, facilities such as a thermal power generation boiler and a chemical plant operate in a high temperature environment (for example, 400 to 800° C.) and further come into contact with a process fluid containing sulfide and/or chloride. Therefore, the materials used for these facilities are required to have high-temperature creep strength and corrosion resistance.

Materials used for such equipment include, for example, 18-8 series stainless steel such as SUS304H, SUS316H, SUS321H, and SUS347H, and a NiCrFe alloy typified by Alloy800H defined as NCF800H in JIS standard.

The NiCrFe alloy is superior in corrosion resistance and high temperature strength as compared with 18-8 series stainless steel. Further, the NiCrFe alloy is more economical than the Ni-based alloy represented by Alloy 617. Therefore, the NiCrFe alloy is widely used in a severe environment of use.

NiCrFe alloys used in such a harsh use environment are disclosed in JP2013-227644A (Patent Document 1), JPA-6-264169 (Patent Document 2), and JP2002-256398A (Patent Document). 3) and JP-A-8-13104 (Patent Document 4).

The austenitic heat-resistant alloy disclosed in Patent Document 1 is, by mass%, C: less than 0.02%, Si: 2% or less, Mn: 2% or less, Cr: 20% or more and less than 28%, Ni: 35%. To 50% or less, W: 2.0 to 7.0%, Mo: less than 2.5% (including 0%), Nb: less than 2.5% (including 0%), Ti: 3. Less than 0% (including 0%), Al: 0.3% or less, P: 0.04% or less, S: 0.01% or less and N: 0.05% or less, and the balance Fe and impurities And f1=(1/2)W+Mo is 1.0 to 5.0, f2=(1/2)W+Mo+Nb+2Ti is 2.0 to 8.0, and f3=Nb+2Ti is 0.5 to 5.0. Is.

The heat-resistant and corrosion-resistant alloys disclosed in Patent Document 2 are, by weight %, nickel 55-65%, chromium 19-25%, aluminum 1-4.5%, yttrium 0.045-0.3%, titanium. 0.15 to 1%, carbon 0.005 to 0.5%, silicon 0.1 to 1.5%, manganese 1% or less, at least one element selected from the group consisting of magnesium, calcium and cerium. 0.005% total, less than 0.5% total of magnesium and calcium, less than 1% cerium, 0.0001-0.1% boron, 0.5% or less zirconium, 0.0001-0.2% nitrogen, cobalt Up to 10% and the balance consists of iron and associated impurities.

The austenitic alloy disclosed in Patent Document 3 contains C: 0.01 to 0.1%, Mn: 0.05 to 2%, Cr: 19 to 26%, and Ni: 10 to 35% in mass%. The content of Si satisfies the expression 0.01<Si<(Cr+0.15×Ni-18)/10.

The heat-resistant alloy disclosed in Patent Document 4 is, by weight %, C: 0.02 to 0.15%, Si: 0.70 to 3.00%, Mn: 0.50% or less, Ni: 30.0. .About.40.0%, Cr: 18.0 to 25.0%, Al: 0.50 to 2.00%, Ti: 0.10 to 1.00%, with the balance being Fe and unavoidable impurities.

JP, 2013-227644, A JP-A-6-264169 JP, 2002-256398, A JP-A-8-13104

Hans van Wortel: "Control of Relaxation Cracking in Austenitic High Temperature Components", CORROSION 2007 (2007), NACE, Paper No. 07423

The austenitic heat-resistant alloy disclosed in Patent Document 1 controls the generation of Laves phase and regulates the creep strength and toughness by defining the W, Mo, Nb, and Ti contents. The heat-resistant and corrosion-resistant alloy disclosed in Patent Document 2 improves the high temperature oxidation resistance by precipitating γ′ during creep. The austenitic alloy disclosed in Patent Document 3 has improved carburizing property by suppressing peeling of the oxide film mainly composed of Cr ₂ O ₃ formed on the material surface. The heat-resistant alloy disclosed in Patent Document 4 contains a specific amount of Cr, reduces Mn, and contains a certain amount of Si, so that good oxidation resistance can be obtained even if the Ni content is reduced. There is.

On the other hand, Non-Patent Document 1 discloses that the NiCrFe alloy has high susceptibility to stress relaxation cracking. That is, in the NiCrFe alloy, it is necessary to perform stress relief heat treatment on the bent portion and the welded portion where residual stress exists after the construction. Therefore, the NiCrFe alloy is required to have not only excellent creep strength but also excellent stress relaxation crack resistance.

An object of the present invention is to provide a NiCrFe alloy excellent in creep strength and resistance to stress relaxation cracking.

The NiCrFe alloy according to the present invention is, in mass%, C: 0.03 to 0.15%, Si: 1.00% or less, Mn: 2.00% or less, P: 0.040% or less, S: 0.0. 0050% or less, Cr: 18.0 to 25.0%, Ni: 25.0 to 40.0%, Ti: 0.10 to 1.60%, Al: 0.05 to 1.00%, N: 0.020% or less, O: 0.008% or less, rare earth element (REM): 0.001 to 0.100%, B: 0 to 0.010%, Ca: 0 to 0.010%, Mg: 0 ~0.010%, V:0~0.5%, Nb:0~1.0%, Ta:0~1.0%, Hf:0~1.0%, Mo:0~1.0%. , W: 0 to 2.0%, Co: 0 to 3.0%, and Cu: 0 to 3.0%, the balance consisting of Fe and impurities and satisfying the formulas (1) to (3). Has a chemical composition.
0.50≦Ti+48Al/27≦2.20 (1)
0.40≦Ti/(Ti+48Al/27)≦0.80 (2)
Σ[REM/(A(REM))]-S/32-2/3・O/16≧0 (3)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the above formula. The atomic weight of each rare earth element is substituted into A(REM) in the formula (3).

The NiCrFe alloy according to the present invention is excellent in creep strength and stress relaxation crack resistance.

FIG. 1 is a diagram showing a relationship between fn2 of each test number in Examples and the sum (mass %) of γ′ and η phases after aging treatment.

The present inventors have investigated in detail the creep strength and stress relaxation crack resistance of NiCrFe alloys. As a result, the present inventors have obtained the following findings.

(A) In order to obtain excellent creep strength, the amount of γ′ (intermetallic compound: Ni ₃ (Ti, Al)) precipitated during creep under a high temperature environment may be increased. If γ′ is sufficiently precipitated during creep in a high temperature environment, precipitation strengthening increases the creep strength of the alloy. However, if γ'precipitates excessively, the deformability in the austenite grains decreases and stress concentration occurs at the grain interfaces. As a result, the stress relaxation cracking resistance of the alloy decreases. Therefore, in order to achieve both excellent creep strength and excellent resistance to stress relaxation cracking, it is necessary to adjust the amount of γ'precipitated during creep under a high temperature environment. In order to make the amount of γ'precipitated to be an appropriate amount, the contents of Ti and Al constituting γ'may be adjusted.

Specifically, the chemical composition of the NiCrFe alloy satisfies the formula (1) in order to maintain the stress relaxation crack resistance while ensuring the creep strength.
0.50≦Ti+48Al/27≦2.20 (1)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (1).

It is defined as fn1=Ti+48Al/27. fn1 is an index indicating the amount of γ'precipitated during creep. fn1 is the total content of Al and Ti converted to the amount of Ti. If fn1 is lower than 0.50, a sufficient amount of γ'precipitated cannot be obtained. Therefore, NiCrFe alloy cannot obtain excellent creep strength. On the other hand, when fn1 is higher than 2.20, the stress relaxation cracking resistance of the NiCrFe alloy decreases due to the precipitation of a large amount of γ'.

(B) The morphology of γ'precipitated during creep in a high temperature environment may change over time. Specifically, fine γ′ precipitates in the initial stage of creep, but with the passage of time, γ′ may change to a coarse and acicular η phase (Ni ₃ Ti) during creep in a high temperature environment. is there. The formation of the η phase reduces the creep strength of the NiCrFe alloy.

Therefore, the present inventors have studied in detail the case where the γ′ phase changes to the η phase under a high temperature environment. As a result, it was considered that the Ti content with respect to the total content of Al and Ti converted into the Ti content may be related to the change from the γ′ phase to the η phase. Therefore, the present inventors have studied in detail the Ti content relative to the total content of Al and Ti converted to the Ti content and the structure during creep.

It is defined as fn2=Ti/(Ti+48Al/27). fn2 is the ratio of the Ti content to the total content of Al and Ti converted into the Ti content. FIG. 1 shows the relationship between fn2 and the sum of the γ′ and η phases after the aging treatment. FIG. 1 was obtained by the following method. Of the examples described below, for the NiCrFe alloys whose chemical compositions are within the scope of the present invention, and the above-described formula (1) and below-described formula (3) are within the scope of the present invention, fn2 and a method described below are used. It was prepared by using the obtained γ′ after aging treatment and the Ti, Al, and Ni contents in the η phase. Further, γ′ and η phase were discriminated by the method described later. “◯” in FIG. 1 means an example in which the number density of the η phase after the aging treatment was less than 5/100 μm ² . On the other hand, “●” in FIG. 1 means an example in which the number density of the η phase after the aging treatment was 5/100 μm ² or more.

Referring to FIG. 1, if fn2 is less than 0.40, a sufficient precipitation amount of γ'is not obtained. In this case, the NiCrFe alloy cannot obtain excellent creep strength. On the other hand, when fn2 exceeds 0.80, γ'changes to the η phase. As a result, the NiCrFe alloy cannot obtain excellent creep strength. Therefore, when fn2 is 0.40 to 0.80, the creep strength of the NiCrFe alloy can be increased.

As described above, when the chemical composition of the NiCrFe alloy of the present invention satisfies the formula (2), an appropriate amount of γ′ is precipitated, and the precipitation of the η phase is suppressed over time, and excellent creep strength is obtained. ..
0.40≦Ti/(Ti+48Al/27)≦0.80 (2)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (2).

One of the causes of (C) stress relaxation cracking is S segregated at grain boundaries. Therefore, the stress relaxation cracking resistance of the NiCrFe alloy can be enhanced by reducing S, which is an impurity that segregates at grain boundaries and causes grain boundary embrittlement. On the other hand, the rare earth element (REM) combines with a small amount of S in the alloy that cannot be removed by refining to form inclusions. That is, the REM can fix S as an inclusion.

Therefore, if the REM content is adjusted to an appropriate amount, the stress relaxation cracking resistance of the NiCrFe alloy can be enhanced. REM binds to S as well as to O easily. Therefore, in order to immobilize S by REM, the REM content should be adjusted in consideration of the amount of REM bound to O.

When the chemical composition of the NiCrFe alloy of the present invention satisfies the formula (3), S is sufficiently fixed by REM and excellent stress relaxation crack resistance is obtained.
Σ[REM/(A(REM))]-S/32-2/3・O/16≧0 (3)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (3), and the atomic weight of each rare earth element is substituted for A(REM).

In Σ[REM/(A(REM))], the addition sum of the values obtained by dividing each REM content (mass %) contained in the NiCrFe alloy by the atomic weight of the REM is substituted.

fn3=Σ[REM/(A(REM))]-S/32-2/3·O/16. REM is a general term for a total of 17 elements of Sc, Y and lanthanoid. If fn3 is 0 or more, REM can be sufficiently fixed with S as an inclusion, and the stress relaxation cracking resistance can be improved.

The NiCrFe alloy according to the present invention completed based on the above findings is, in mass %, C: 0.03 to 0.15%, Si: 1.00% or less, Mn: 2.00% or less, P: 0. 040% or less, S: 0.0050% or less, Cr: 18.0 to 25.0%, Ni: 25.0 to 40.0%, Ti: 0.10 to 1.60%, Al: 0.05. ~1.00%, N: 0.020% or less, O: 0.008% or less, rare earth element (REM): 0.001 to 0.100%, B:0 to 0.010%, Ca:0 to 0.010%, Mg:0 to 0.010%, V:0 to 0.5%, Nb:0 to 1.0%, Ta:0 to 1.0%, Hf:0 to 1.0%, Mo: 0 to 1.0%, W: 0 to 2.0%, Co: 0 to 3.0%, and Cu: 0 to 3.0%, the balance consisting of Fe and impurities. It has a chemical composition that satisfies the formulas (1) to (3).
0.50≦Ti+48Al/27≦2.20 (1)
0.40≦Ti/(Ti+48Al/27)≦0.80 (2)
Σ[REM/(A(REM))]-S/32-2/3・O/16≧0 (3)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formulas (1) to (3). The atomic weight of each rare earth element is substituted into A(REM) in the formula (3).

The above chemical composition may contain B: 0.0001 to 0.010%.

The above chemical composition may contain one or two selected from the group consisting of Ca: 0.0001 to 0.010% and Mg: 0.0001 to 0.010%.

The chemical composition is V: 0.01 to 0.5%, Nb: 0.01 to 1.0%, Ta: 0.01 to 1.0%, and Hf: 0.01 to 1.0%. You may contain 1 type(s) or 2 or more types selected from the group which consists of.

The above chemical composition is Mo: 0.01 to 1.0%, W: 0.01 to 2.0%, Co: 0.01 to 3.0%, and Cu: 0.01 to 3.0%. You may contain 1 type(s) or 2 or more types selected from the group which consists of.

The NiCrFe alloy according to the present invention has excellent creep strength and excellent stress relaxation crack resistance. More specifically, the NiCrFe alloy is subjected to cold rolling with a cross-section reduction rate of 20%, and then held at 10% tensile strain at a strain rate of 0.05 min ⁻¹ in an air atmosphere at 650° C. Does not break for more than 300 hours.

Hereinafter, the NiCrFe alloy according to the present invention will be described in detail. "%" regarding an element means mass% unless otherwise specified.

[Chemical composition]
The chemical composition of the NiCrFe alloy of the present invention contains the following elements.

C: 0.03 to 0.15%
Carbon (C) stabilizes austenite and enhances the creep strength of the alloy at high temperatures. If the C content is too low, these effects cannot be obtained. On the other hand, if the C content is too high, a large amount of coarse carbide precipitates, and the ductility of the grain boundary decreases. Moreover, the toughness and creep strength of the alloy are reduced. Therefore, the C content is 0.03 to 0.15%. The preferable lower limit of the C content is 0.04%, more preferably more than 0.04%, further preferably 0.05%, further preferably 0.06%. The preferable upper limit of the C content is 0.12%, more preferably 0.10%.

Si: 1.00% or less Silicon (Si) is inevitably contained. Si deoxidizes the alloy and enhances the corrosion resistance and oxidation resistance of the alloy at high temperatures. However, if the Si content is too high, the stability of austenite decreases, and the toughness and creep strength of the alloy decrease. Therefore, the Si content is 1.00% or less. The preferable upper limit of the Si content is 0.80%, more preferably 0.60%, and further preferably less than 0.60%. Extreme reduction of the Si content lowers the deoxidizing effect and lowers the corrosion resistance and oxidation resistance of the alloy at high temperatures. Further, it significantly increases the manufacturing cost. Therefore, the lower limit of the Si content is preferably 0.02%, more preferably 0.05%.

Mn: 2.00% or less Manganese (Mn) is unavoidably contained. Mn deoxidizes the alloy and stabilizes austenite. However, if the Mn content is too high, embrittlement occurs, and the toughness and creep ductility of the alloy decrease. Therefore, the Mn content is 2.00% or less. The preferable upper limit of the Mn content is 1.80%, more preferably 1.50%. The extreme reduction of the Mn content reduces the deoxidizing effect and the austenite stabilization. Further, it significantly increases the manufacturing cost. Therefore, the preferable lower limit of the Mn content is 0.10%, more preferably 0.30%, and further preferably more than 0.50%.

P: 0.040% or less Phosphorus (P) is an impurity. P reduces the hot workability and weldability of the alloy, and reduces the creep ductility of the alloy after long-term use. Therefore, the P content is 0.040% or less. The preferable upper limit of the P content is 0.035%, more preferably 0.030%. It is preferable that the P content is as low as possible. However, the extreme reduction of P content increases the manufacturing cost. Therefore, the preferable lower limit of the P content is 0.0005%, and more preferably 0.0008%.

S: 0.0050% or less Sulfur (S) is an impurity. S reduces the stress relaxation cracking resistance of the alloy and also reduces the hot workability, weldability and creep ductility of the alloy. Therefore, the S content is 0.0050% or less. The preferable upper limit of the S content is 0.0030%. It is preferable that the S content is as low as possible. However, the extreme reduction of the S content increases the manufacturing cost. Therefore, the preferred lower limit of the S content is 0.0002%, and more preferably 0.0003%.

Cr: 18.0 to 25.0%
Chromium (Cr) enhances the high temperature oxidation and corrosion resistance of the alloy. If the Cr content is too low, these effects cannot be obtained. On the other hand, if the Cr content is too high, the stability of austenite at high temperature is reduced and the creep strength of the alloy is reduced. Therefore, the Cr content is 18.0 to 25.0%. The preferable lower limit of the Cr content is 18.5%, more preferably 19.0%. The preferable upper limit of the Cr content is 24.5%, more preferably 24.0%.

Ni: 25.0-40.0%
Nickel (Ni) stabilizes the austenite structure. Ni further forms γ', increasing the creep strength of the alloy. If the Ni content is too low, γ'is less likely to be formed and these effects cannot be obtained. On the other hand, if the Ni content is too high, the manufacturing cost will increase. Therefore, the Ni content is 25.0 to 40.0%. The preferable lower limit of the Ni content is 26.0%, more preferably 27.0%. The preferable upper limit of the Ni content is 37.0%, and more preferably 35.0%.

Ti: 0.10 to 1.60%
Titanium (Ti) combines with Ni to form γ'. Ti also combines with C to form TiC, which enhances the creep strength and tensile strength of the alloy at high temperatures. If the Ti content is too low, these effects cannot be obtained. On the other hand, if the Ti content is too high, γ'precipitates excessively, and the stress relaxation cracking resistance of the alloy decreases. Therefore, the Ti content is 0.10 to 1.60%. The preferable lower limit of the Ti content is 0.20%, more preferably 0.30%, and further preferably more than 0.60%. Moreover, the preferable upper limit of the Ti content is 1.50%, more preferably less than 1.50%, and further preferably 1.40%.

Al: 0.05-1.00%
Aluminum (Al) deoxidizes the alloy. Al also combines with Ni to form γ', which enhances the creep strength and tensile strength of the alloy at high temperatures. If the Al content is too low, these effects cannot be obtained. On the other hand, if the Al content is too high, a large amount of γ'is precipitated, and the stress relaxation crack resistance, creep ductility and toughness of the alloy are reduced. Therefore, the Al content is 0.05 to 1.00%. The preferable lower limit of the Al content is 0.08%, more preferably 0.10%. The preferable upper limit of the Al content is 0.90%, more preferably 0.80%.

N: 0.020% or less Nitrogen (N) is an impurity. N precipitates as coarse TiN to reduce the amount of solid solution Ti and reduce the creep strength of the alloy. N further reduces the toughness and hot workability of the alloy. Therefore, the N content is 0.020% or less. The preferable upper limit of the N content is 0.017%, and more preferably 0.015%. The N content is preferably as low as possible. However, extreme reductions increase manufacturing costs. Therefore, the lower limit of the N content is preferably 0.002%, and more preferably 0.004%.

O: 0.008% or less O (oxygen) is an impurity. O reduces the hot workability of the alloy and reduces the toughness and ductility of the alloy. Therefore, the O content is 0.008% or less. The preferable upper limit of the O content is 0.006%, and more preferably 0.005%. The O content is preferably as low as possible. However, extreme reductions increase manufacturing costs. Therefore, the preferable lower limit of the O content is 0.0005%, and more preferably 0.0008%.

REM: 0.001 to 0.100%
The rare earth element (REM) forms a compound with S to reduce the content of S dissolved in the matrix, and enhances the stress relaxation cracking resistance of the alloy. REM further improves the hot workability and oxidation resistance of the alloy. If the REM content is too low, these effects cannot be obtained. On the other hand, if the REM content is too high, the hot workability and weldability of the alloy deteriorate. Therefore, the REM content is 0.001 to 0.100%. The preferable lower limit of the REM content is 0.003%, more preferably 0.005%. The preferable upper limit of the REM content is 0.090%, more preferably 0.080%.

REM is a general term for a total of 17 elements of Sc, Y and lanthanoid, and the REM content refers to the total content of one or more elements of REM. Further, REM is generally contained in misch metal. Therefore, for example, it may be added as a misch metal to the molten metal and adjusted so that the amount of REM falls within the above range.

The balance of the chemical composition of the NiCrFe alloy according to the invention consists of Fe and impurities. Here, the impurities are those that are mixed from ore as a raw material, scrap, or a manufacturing environment when industrially producing the NiCrFe alloy, and do not adversely affect the NiCrFe alloy of the present embodiment. Means acceptable.

[About arbitrary elements]
The NiCrFe alloy according to the present invention may further contain B in place of part of Fe.

B: 0 to 0.010%
Boron (B) is an optional element and may not be contained. When contained, B improves the creep strength of the alloy by finely dispersing grain boundary carbides. B further segregates at the grain boundaries to assist the effect of REM. If B is contained even in a small amount, the above effect can be obtained to some extent. However, if the B content is too high, the weldability and hot workability of the alloy deteriorate. Therefore, the B content is 0 to 0.010%. The preferable upper limit of the B content is 0.008%. A preferable lower limit of the B content for effectively obtaining the above effect is 0.0001%, and more preferably 0.0005%.

The NiCrFe alloy according to the present invention may further contain one kind or two kinds selected from the group consisting of Ca and Mg instead of part of Fe. All of these elements form a compound with S and assist the effect of REM.

Ca: 0 to 0.010%
Calcium (Ca) is an optional element and may not be contained. When included, Ca forms a compound with S and assists the S fixing effect of REM. If Ca is contained even in a small amount, the above effect can be obtained to some extent. However, if the Ca content is too high, an oxide is formed and the hot workability of the alloy is deteriorated. Therefore, the Ca content is 0 to 0.010%. The preferable upper limit of the Ca content is 0.008%. The preferable lower limit of the Ca content for effectively obtaining the above effect is 0.0001%, more preferably 0.0002%, and further preferably 0.0003%.

Mg: 0 to 0.010%
Magnesium (Mg) is an optional element and may not be contained. When contained, it forms a compound with S and assists the S fixing effect of REM. If Mg is contained even a little, the above effect can be obtained to some extent. However, if the Mg content is too high, an oxide is formed and the hot workability of the alloy is deteriorated. Therefore, the Mg content is 0 to 0.010%. The preferable upper limit of the Mg content is 0.008%. The preferable lower limit of the Mg content for effectively obtaining the above effect is 0.0001%, more preferably 0.0002%, and further preferably 0.0003%.

The NiCrFe alloy according to the present invention may further contain one kind or two or more kinds selected from the group consisting of V, Nb, Ta and Hf in place of a part of Fe. All of these elements form carbides and carbonitrides and enhance the creep strength of the alloy.

V: 0 to 0.5%
Vanadium (V) is an optional element and may not be contained. When contained, V forms fine carbides and carbonitrides with C and N and enhances the creep strength of the alloy. If V is contained even a little, the above effect can be obtained to some extent. However, if the V content is too high, a large amount of carbides and carbonitrides are deposited, and the creep ductility of the alloy decreases. Therefore, the V content is 0 to 0.5%. The preferable upper limit of the V content is 0.4%. The lower limit of the V content for effectively obtaining the above effect is 0.01%.

Nb: 0 to 1.0%
Niobium (Nb) is an optional element and may not be contained. When contained, Nb forms fine carbides and carbonitrides with C and N, and enhances the creep strength of the alloy. If Nb is contained even a little, the above effect can be obtained to some extent. However, if the Nb content is too high, a large amount of carbides and carbonitrides are deposited, and the creep ductility and toughness of the alloy are reduced. Therefore, the Nb content is 0 to 1.0%. The preferable upper limit of the Nb content is 0.4%. The lower limit of the Nb content for effectively obtaining the above effect is 0.01%.

Ta: 0 to 1.0%
Tantalum (Ta) is an optional element and may not be contained. When contained, Ta forms fine carbides and carbonitrides with C and N and enhances the creep strength of the alloy. If Ta is contained even a little, the above effect can be obtained to some extent. However, if the Ta content is too high, a large amount of carbides and carbonitrides will precipitate, and the creep ductility and toughness of the alloy will decrease. Therefore, the Ta content is 0 to 1.0%. The preferable upper limit of the Ta content is 0.4%. The lower limit of the Ta content for effectively obtaining the above effect is 0.01%.

Hf: 0-1.0%
Hafnium (Hf) is an optional element and may not be contained. When included, Hf forms fine carbides and carbonitrides with C and N and enhances the creep strength of the alloy. If Hf is contained even a little, the above effect can be obtained to some extent. However, if the Hf content is too high, a large amount of carbides and carbonitrides will precipitate, and the creep ductility and toughness of the alloy will decrease. Therefore, the Hf content is 0 to 1.0%. The preferable upper limit of the Hf content is 0.4%. The lower limit of the Hf content for effectively obtaining the above effect is 0.01%.

The NiCrFe alloy according to the present invention may further contain one kind or two or more kinds selected from the group consisting of Mo, W, Co and Cu in place of a part of Fe.

Mo: 0 to 1.0%
Molybdenum (Mo) is an optional element and may not be contained. When included, Mo forms a solid solution in the alloy and enhances the creep strength of the alloy at high temperatures. If Mo is contained even a little, this effect can be obtained to some extent. However, if the Mo content is too high, the stability of austenite is lost and the toughness of the alloy is reduced. Therefore, the Mo content is 0 to 1.0%. The preferable upper limit of the Mo content is 0.9%. The preferable lower limit of the Mo content for effectively obtaining the above effect is 0.01%.

W: 0-2.0%
Tungsten (W) is an optional element and may not be contained. When included, W dissolves in the alloy to enhance the creep strength of the alloy at high temperatures. If W is contained in even a small amount, this effect can be obtained to some extent. However, if the W content is too high, the stability of austenite is lost and the toughness of the alloy is reduced. Therefore, the W content is 0 to 2.0%. The preferable upper limit of the W content is 1.8%. The preferable lower limit of the W content for effectively obtaining the above effect is 0.01%.

Co: 0-3.0%
Cobalt (Co) is an optional element and may not be contained. When included, Co stabilizes austenite and also forms a solid solution with the alloy to enhance the creep strength of the alloy at high temperatures. This effect can be obtained to some extent if Co is contained in a small amount. However, if the Co content is too high, the manufacturing cost will increase. Therefore, the Co content is 0 to 3.0%. The preferable upper limit of the Co content is 2.8%. The preferable lower limit of the Co content for effectively obtaining the above effect is 0.01%.

Cu: 0 to 3.0%
Copper (Cu) is an optional element and may not be contained. When contained, Cu stabilizes austenite and suppresses precipitation of an embrittlement phase such as a σ phase during use at high temperature. This effect can be obtained to some extent if Cu is contained at least. However, if the Cu content is too high, the hot workability of the alloy deteriorates. Therefore, the Cu content is 0 to 3.0%. The preferable upper limit of the Cu content is 2.5%, more preferably less than 2.0%. The preferable lower limit of the Cu content for effectively obtaining the above effect is 0.01%.

[About Formula (1)]
The NiCrFe alloy according to the present invention further satisfies equation (1).
0.50≦Ti+48Al/27≦2.20 (1)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (1).

fn1=Ti+48Al/27 is an index showing the precipitation amount of γ′. fn1 represents the total amount of Ti when Al is converted into the amount of Ti. If fn1 is lower than 0.50, a sufficient precipitation amount of γ'is not obtained, and good creep characteristics of the alloy cannot be obtained. On the other hand, when fn1 is higher than 2.20, the amount of γ'precipitated becomes too large, and the stress relaxation cracking resistance, creep ductility and toughness of the alloy deteriorate. Therefore, fn1 is 0.50 to 2.20. In this case, γ'has an appropriate amount of precipitation, and good creep characteristics can be obtained. The preferable upper limit of fn1 is 2.00. The preferable lower limit of fn1 is 0.65.

[About formula (2)]
The above chemical composition further satisfies equation (2).
0.40≦Ti/(Ti+48Al/27)≦0.80 (2)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (2).

fn2=Ti/(Ti+48Al/27) is the ratio of the Ti content to the total content of Al and Ti converted into the Ti content. When fn2 is lower than 0.40, the Ti content is too small with respect to the Al content, and the precipitation amount of γ'is reduced. As a result, NiCrFe cannot obtain excellent creep strength. On the other hand, if fn2 is higher than 0.80, the Ti content is too large relative to the Al content and precipitates as fine γ'in the initial stage of creep, but with the passage of time, coarse and needle-like η phase is formed. Change. As a result, the creep strength and toughness of the alloy decrease. Therefore, fn2 is 0.40 to 0.80. In this case, an appropriate amount of γ′ is precipitated and does not change to the η phase even after a lapse of time, so that good creep strength can be obtained. The preferable upper limit of fn2 is 0.75.

[Regarding Expression (3)]
The chemical composition further satisfies equation (3).
Σ[REM/(A(REM))]-S/32-2/3・O/16≧0 (3)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formula (3), and the atomic weight of each REM is substituted for A(REM).

fn3=Σ[REM/(A(REM))]-S/32-2/3·O/16 is an index showing the amount of S segregated at the grain boundaries. When fn3 is a negative value, S segregates at the grain boundaries, causing grain boundary embrittlement and reducing the stress relaxation cracking resistance of the alloy. On the other hand, if fn3 is 0 or more, REM fixes S as inclusions and reduces the S content in the matrix. As a result, the stress relaxation cracking resistance of the alloy can be enhanced. Therefore, fn3 is 0 or more.

[Production method]
An example of the method for manufacturing the NiCrFe alloy of this embodiment will be described. The manufacturing method of the present embodiment includes a step of manufacturing an ingot (steel making step) and a step of manufacturing a hot rolled sheet (hot working step). Hereinafter, each step will be described in detail.

[Steel making process]
First, an alloy having the above chemical composition is melted. The melting is performed using, for example, high frequency vacuum melting. Then, an ingot is manufactured by the ingot making method.

[Hot working process]
In the hot working step, hot working is usually performed once or a plurality of times. First, the ingot is heated, and then hot working is performed. The hot working is, for example, hot forging or hot rolling. The hot working may be performed by a known method.

Further, cold working may be performed on the hot worked NiCrFe alloy. Cold working is, for example, cold rolling.

Further, heat treatment may be performed on the NiCrFe alloy processed as described above. The preferable heat treatment temperature is 1050 to 1200°C. Further, it is preferable that the NiCrFe alloy after heating and holding is cooled with water.

In the example of the manufacturing method described above, the manufacturing method of the NiCrFe alloy plate has been described. However, the NiCrFe alloy may be a rod or an alloy tube. That is, the product shape is not limited. Further, in the case of an alloy tube, it is preferable to carry out hot working by hot extrusion.

The NiCrFe alloy produced by the above steps has excellent creep strength and excellent stress relaxation crack resistance.

[Microstructure]
In the NiCrFe alloy according to the present invention, γ'and η phases are precipitated in the use environment at high temperature. That is, the microstructure of the NiCrFe alloy according to the present invention after holding at 650° C. for 3000 hours contains 2 to 6 mass% of γ′ and η phases in total, and the number density of the η phase is less than 5/100 μm ^2. Is. In the present specification, γ′ and η phase are also collectively referred to as “aging precipitate”.

When the NiCrFe alloy according to the present invention is subjected to an aging treatment in which the NiCrFe alloy is held at 650° C. for 3000 hours, if the total of γ′ and η phases is less than 2% by mass, the amount of γ′ precipitated in the alloy decreases. As a result, the NiCrFe alloy cannot obtain excellent creep strength. On the other hand, when the same aging treatment is carried out, if the total of γ'and η phases exceeds 6% by mass, the precipitation amount of γ'may be too large. In this case, the alloy does not have excellent resistance to stress relaxation cracking. Therefore, the total of γ′ and η phases after the aging treatment is 2 to 6 mass %.

Specifically, the total of γ'and η phases can be measured by the following method. The NiCrFe alloy according to the present invention is subjected to an aging treatment in which it is held at 650° C. for 3000 hours. A test piece of 10 mm×5 mm×50 mm is taken from the NiCrFe alloy after the aging treatment. When the alloy is an alloy plate, a test piece is taken from the center part of the plate thickness. On the other hand, when the alloy is an alloy tube, a test piece is taken from the center part of the wall thickness of the alloy tube. The weight of the test piece is measured in advance.

The collected test specimen 1% tartaric acid _{-1% (NH 4) 2 SO} 4 - electrolysis in an aqueous solution, and collecting the residue from the electrolyte. The collected residue is dissolved in HCl(1+4)-20% tartaric acid solution at 60° C. and the solution is filtered. The filtrate is quantified by ICP emission spectroscopy to determine the Ti, Al, and Ni concentrations in the residue. The contents of Ti, Al, and Ni in the γ′ and η phases of the test piece are determined from the obtained Ti, Al, and Ni concentrations in the residue and the weight of the test piece. The sum of the Ti, Al, and Ni contents obtained by the above method is defined as the sum (mass %) of the γ'and η phases.

When the NiCrFe alloy according to the present invention is subjected to an aging treatment of holding it at 650° C. for 3000 hours, if the number density of η phase is 5/100 μm ² or more, a part of γ′ is changed to the η phase. There is. Therefore, NiCrFe alloy cannot obtain excellent creep strength. Therefore, the number density of the η phase after the aging treatment is less than 5 pieces/100 μm ² .

Specifically, the number density of the η phase can be measured by the following method. The NiCrFe alloy according to the present invention is subjected to an aging treatment in which it is held at 650° C. for 3000 hours. Microscopic observation is performed on the NiCrFe alloy after the aging treatment. Specifically, a micro test piece is taken from the NiCrFe alloy after the aging treatment. When the alloy is an alloy plate, a test piece is taken from the center part of the plate thickness. On the other hand, when the alloy is an alloy tube, a micro test piece is taken from the center part of the wall thickness of the alloy tube. The collected micro test piece is mechanically polished. The surface of the micro test piece after mechanical polishing is electrolytically corroded with 10% oxalic acid. With respect to the micro test piece after electrolytic corrosion, 5 fields of view are observed with a scanning electron microscope (SEM: Scanning Electron Microscope), and SEM images of each field of view are generated. The observation magnification is 10,000 times, and the observation visual field is, for example, 12 μm×9 μm.

The shapes of γ'and η phase are different. Specifically, γ′ is spherical and the η phase is needle-shaped. More specifically, the aspect ratio of γ'is less than 3, and the aspect ratio of the η phase is 3 or more. Here, the aspect ratio means a value obtained by dividing the major axis length by the minor axis length of each aging precipitate.

Aging precipitates (γ' and η phase) are identified from the contrast in the SEM images of the above-mentioned visual fields. Further, the aspect ratio of the identified aged precipitate is calculated by image processing. General-purpose application software may be used to calculate the aspect ratio. If the calculated aspect ratio is 3 or more, the aging precipitate is identified as the η phase.

For the SEM image of each visual field, the specified η phase is counted, and the sum of all visual fields is obtained. The number density (number/100 μm ² ) of the η phase in the observation visual field of 100 μm ² is obtained using the number of η phases in the entire visual field and the total visual field area.

Alloys having the chemical compositions 1 to 15 shown in Table 1 were melted by a high frequency vacuum melting method.

A 50 kg ingot was manufactured using each code alloy. Hot forging and hot rolling were performed on the ingot to obtain a plate material having a thickness of 15 mm. Each plate was held at 1150° C. for 30 minutes, and then the plate was rapidly cooled (water-cooled) to carry out solution treatment. A NiCrFe alloy plate material was manufactured by the above manufacturing process. The following test was implemented using the manufactured NiCrFe alloy plate material.

[Creep rupture test]
A test piece was prepared from the produced alloy plate material. The test piece was sampled parallel to the longitudinal direction (rolling direction) from the center of thickness of the alloy plate material. The test piece was a round bar test piece, and the diameter of the parallel portion was 6 mm and the gauge length was 30 mm. A creep rupture test was conducted using the test piece. The creep rupture test was carried out in an air atmosphere at 750° C. under a tensile load of 70 MPa. The one having a break time of 3000 hours or more was evaluated as "E" (Excellent) and the one having a break time of less than 3000 hours was evaluated as "NA" (Not Acceptable).

[Microstructure observation]
A test piece was produced from the produced alloy plate material by the method described above. The produced test piece was subjected to an aging treatment in which the test piece was held at 650° C. for 3000 hours, and the sum (mass %) of γ′ and η phases was obtained by the above method. Further, the number density of η phases (number/100 μm ² ) was determined by the above method. The sum of γ′ and η phases was evaluated as “L” (Less) when less than 2% by mass, “E” (Excellent) when 2 to 6% by mass, and “TM” (Too Much) when more than 6% by mass. .. Furthermore, the number density of the η phase was 5/100 μm ² or more, and was evaluated as “η”.

[Stress relaxation cracking test]
Further, cold working was performed on the manufactured alloy plate material. Specifically, the alloy plate material was cold-rolled to a thickness of 12 mm. The cross-section reduction rate of this cold rolling was 20%. A test piece was prepared from this alloy plate material. Samples were taken parallel to the longitudinal direction (rolling direction) from the center of thickness of the alloy plate material. The test piece was a round bar test piece, and the diameter of the parallel portion was 6 mm and the gauge length was 30 mm. A stress relaxation test was performed using the test piece. In the stress relaxation test, 10% tensile strain was applied at a strain rate of 0.05 min ⁻¹ in an air atmosphere at 650° C. and the state was maintained for 300 hours. The sample which was kept for 300 hours and was not broken was evaluated as “E” (Excellent), and the sample which was broken was evaluated as “NA” (Not Acceptable).

[Test results]
The test results are shown in Table 2.

With reference to Table 2, the chemical compositions of reference numbers 1 to 8 were appropriate, fn1 was 0.50 to 2.20, fn2 was 0.40 to 0.80, and fn3 was 0 or more. Therefore, in the microstructure, the γ'and η phases were 2 to 6% by mass. Furthermore, the number density of the η phase was less than 5/100 μm ² . As a result, the creep rupture time was 3000 hours or more, which showed excellent creep strength. Further, the test piece did not break in the stress relaxation cracking test and showed excellent stress relaxation cracking resistance.

On the other hand, with the code 9, the value of fn1 was too low. Therefore, the microstructure was too small, with the sum of γ'and η phases being less than 2% by mass. As a result, the creep rupture time was less than 3000 hours and did not show excellent creep strength.

At reference numeral 10, the value of fn1 was too high. Therefore, in the microstructure, the sum of γ′ and η phases exceeded 6 mass %. Furthermore, the number density of the η phase was less than 5/100 μm ² . That is, in the microstructure, γ'exceeded 6 mass% and was too much. As a result, the test piece broke in the stress relaxation cracking test and did not show excellent stress corrosion cracking resistance.

At numbers 11 and 12, the value of fn2 was too low. Therefore, the microstructure was too small, with the sum of γ'and η phases being less than 2% by mass. As a result, the creep rupture time was less than 3000 hours and did not show excellent creep strength.

At reference numeral 13, the value of fn2 was too high. Therefore, the microstructure had a number density of η phase of 5/100 μm ² or more. As a result, the creep rupture time was less than 3000 hours and did not show excellent creep strength.

At reference numeral 14, the value of fn3 was too low. As a result, the test piece broke in the stress relaxation cracking test and did not show excellent stress corrosion cracking resistance. It is considered that S in the matrix could not be fixed.

In reference numeral 15, the REM content was too low. Furthermore, the value of fn3 was too low. As a result, the test piece broke in the stress relaxation cracking test and did not show excellent stress corrosion cracking resistance. It is considered that S in the matrix could not be fixed.

The embodiments of the present invention have been described above. However, the embodiments described above are merely examples for implementing the present invention. Therefore, the present invention is not limited to the above-described embodiments, and can be implemented by appropriately modifying the above-described embodiments without departing from the spirit thereof.

INDUSTRIAL APPLICABILITY The present invention can be widely applied to applications where creep strength and stress relaxation crack resistance are required. The present invention can be suitably used particularly as a high temperature member for a thermal power generation boiler or a chemical industrial plant for petroleum refining.

Claims (6)

In mass %,
C: 0.03 to 0.15%,
Si: 1.00% or less,
Mn: 2.00% or less,
P: 0.040% or less,
S: 0.0050% or less,
Cr: 18.0 to 25.0%,
Ni: 25.0-40.0%,
Ti: 0.10 to 1.60%,
Al: 0.05 to 1.00%,
N: 0.020% or less,
O: 0.008% or less,
Rare earth element (REM): 0.001 to 0.100%,
B: 0 to 0.010%,
Ca: 0 to 0.010%,
Mg: 0 to 0.010%,
V: 0 to 0.5%,
Nb: 0 to 1.0%,
Ta: 0 to 1.0%,
Hf: 0 to 1.0%,
Mo: 0 to 1.0%,
W: 0 to 2.0%,
Co: 0 to 3.0%, and
Cu: A NiCrFe alloy containing 0 to 3.0%, the balance being Fe and impurities, and having a chemical composition satisfying the following formulas (1) to (3).
0.50≦Ti+48Al/27≦2.20 (1)
0.40≦Ti/(Ti+48Al/27)≦0.80 (2)
Σ[REM/(A(REM))]-S/32-2/3・O/16≧0 (3)
Here, the content (mass %) of the corresponding element is substituted for the element symbol in the formulas (1) to (3). The atomic weight of each rare earth element is substituted into A(REM) in the formula (3). The NiCrFe alloy according to claim 1, wherein
The chemical composition is
B: NiCrFe alloy containing 0.0001 to 0.010%. The NiCrFe alloy according to claim 1 or 2, wherein
The chemical composition is
Ca: 0.0001 to 0.010%, and
Mg: A NiCrFe alloy containing one or two selected from the group consisting of 0.0001 to 0.010%. The NiCrFe alloy according to any one of claims 1 to 3,
The chemical composition is
V: 0.01 to 0.5%,
Nb: 0.01 to 1.0%,
Ta: 0.01 to 1.0%, and
Hf: NiCrFe alloy containing one or more selected from the group consisting of 0.01 to 1.0%. The NiCrFe alloy according to any one of claims 1 to 4,
The chemical composition is
Mo: 0.01 to 1.0%,
W: 0.01 to 2.0%,
Co: 0.01 to 3.0%, and
Cu: A NiCrFe alloy containing one or more selected from the group consisting of 0.01 to 3.0%. The NiCrFe alloy according to any one of claims 1 to 5, which is cold-rolled at a cross-section reduction rate of 20%, and then at a strain rate of 0.05 min ^-1 in an air atmosphere at 650°C. A NiCrFe alloy that does not break for 300 hours or more in a stress relaxation test in which a tensile strain of 10% is maintained.

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