JP6620475B2 - Manufacturing method of Ni-base heat-resistant alloy tube - Google Patents

Description

  The present invention relates to a method for manufacturing a Ni-base heat-resistant alloy tube. Specifically, the present invention relates to a method for producing a Ni-based heat-resistant alloy tube used as a pipe material in a power generation boiler, a chemical industry plant, or the like.

Currently, a next-generation advanced ultra super critical pressure power plant is being developed for high efficiency and CO ₂ emission reduction in a coal-fired power plant. One such plant material is Alloy 617 alloy. This material exhibits very high temperature strength due to precipitation strengthening of carbides such as M ₂₃ C ₆ and intermetallic compound phases such as γ′-Ni ₃ (Al, Ti) phase.

  In Patent Document 1, the applicant discloses an Ni-based alloy having an α-Cr phase and a carbide as a strengthening phase. This alloy has excellent creep strength at high temperatures and excellent workability. In the patent document 2, the applicant discloses an austenitic stainless steel excellent in high temperature strength and corrosion resistance. Further, the applicant has disclosed a Ni-based heat-resistant alloy that can avoid SR cracking, which is a problem in repair welding, in Patent Document 3, and a patent document 4, in which the creep strength and crack resistance during welding when multilayer welding is performed. Patent Document 5 discloses Ni-base heat-resistant alloy members that can prevent surface defects during hot working even if the thickness is 20 mm or more. .

International Publication No. 2009/154161 JP 2004-3000 A JP 2013-36086 A JP 2014-148702 A JP 2014-156628 A

  When used in a thermal power plant, there are cases where bending, expanding, flattening, straightening, and the like are performed cold or hot. In particular, a small diameter pipe used for a boiler superheater pipe is subjected to hot or cold processing, such as bending and expanding a pipe, at a part of the pipe. At this time, a partially processed portion is formed in the tube. Patent Documents 1 to 5 do not mention processing at the time of such construction. When the amount of processing in this partially processed portion exceeds a certain value, fine precipitates are formed on the dislocations in the crystal grains, and these precipitates have a small amount of deformation in the grains during stress relaxation and break at the grain boundaries. Sometimes. In addition, due to the high dislocation density, recovery and recrystallization during creep may occur early and the creep strength may decrease. Therefore, it is necessary to perform a post heat treatment on the partial processed portion.

  The post heat treatment is often performed at a temperature lower than the product heat treatment temperature due to limitations of the heat treatment apparatus. For this reason, a carbide | carbonized_material or an intermetallic compound may precipitate from a base material. The carbide and intermetallic compound precipitated in such a temperature range are coarser than the carbide and intermetallic compound precipitated at the actual use temperature (for example, 700 ° C.), and the precipitation density is low. For this reason, it does not contribute to precipitation strengthening at high temperatures. In addition, the amount of solid solution of the strengthening element in the parent phase is reduced, and the amount of carbide and intermetallic compound precipitated at 700 ° C. is also reduced, so that the precipitation strengthening ability is lowered.

  Therefore, unlike the non-processed part, the partial processed part is introduced with strain. When post-heat treatment is carried out in the heat treatment temperature range and time range similar to those at the time of solution treatment, strain remains, or a large amount of carbides and intermetallic compounds precipitate in the grain boundaries / grains. In some cases, the strength of the actual machine is reduced. The post heat treatment temperature also varies depending on the content of the alloy element. This is presumably because the diffusion rate and the precipitation rate of precipitates change due to the inclusion of alloy elements. For this reason, it is necessary to perform post-heat treatment at an appropriate heat treatment temperature and time according to the alloy element content, and to suppress precipitation, while sufficiently removing strain.

  The present invention has been made to solve the above-described problems. When a Ni-based heat-resistant alloy pipe subjected to product heat treatment is constructed, it is bent or partially subjected to hot or cold processing such as pipe expansion. An object of the present invention is to provide a method for manufacturing a Ni-based heat-resistant alloy tube having excellent creep strength in a processed portion when a post-heat treatment is performed after forming a partially processed portion in the tube. .

  The inventors of the present invention performed post-heat treatment on the alloy pipe having the partially processed portion and observed the microstructure and the creep test. As a result, the following knowledge was obtained.

  (a) When the post-heat treatment temperature is the same as the product heat treatment temperature, coarse carbides and intermetallic compounds are slightly observed in the microstructure as in the case of the solution material. In addition, the creep strength is almost the same as that of the heat-treated product.

  (b) Even if the post-heat treatment temperature is equal to or lower than the product heat treatment temperature, if the heat treatment temperature is relatively high, the microstructure and creep strength are almost the same as the material as a solution, but if the post-heat treatment temperature is relatively low, A large amount of carbides and intermetallic compounds are precipitated in the crystal grains or on the crystal grain boundaries. Since these precipitates are coarser than carbides and intermetallic compounds that precipitate at the temperature (approximately 700 ° C. or less) when the alloy tube is used, the contribution to the creep strength is small. For this reason, creep strength falls.

  The inventors have further studied the case where the post-heat treatment temperature is relatively low, and obtained the following knowledge.

  (c) Dislocations introduced by hot working or cold working after product heat treatment remain in the structure, and coarse precipitates arranged in a line on the slip line are observed.

(d) Although precipitates are precipitated at the grain boundaries, they are very coarse and do not contribute to strengthening. The grain boundary coverage ρ is defined by the following equation, where l ₁ , l ₂ , l ₃ ... And the grain boundary length are L as observed by SEM. .
ρ = (l ₁ + l ₂ + l ₃ +...) / L

  (e) Precipitation strengthening elements constituting precipitates are contained in coarse carbides or intermetallic compounds precipitated during post-heat treatment, so that the amount of solid solution in the matrix phase decreases, and carbides and metals that precipitate at the use temperature The volume fraction of the intermetallic compound decreases.

  (f) Since the volume fraction of fine precipitates precipitated at the operating temperature decreases, and the precipitation amount of the precipitates at the grain boundaries decreases, and the grain boundary coverage does not increase. Local deformation is accelerated and the creep strength decreases.

  (g) If the post-heat treatment temperature is sufficiently high, it is possible to suppress the precipitation of carbides and intermetallic compounds that occur during the post-heat treatment, and to suppress the decrease in creep strength at the use temperature. The above phenomenon does not occur in the non-machined part, but is unique to the partially machined part in which strain is introduced. Therefore, it is necessary to control the conditions for the post heat treatment more strictly than the conditions for the solution treatment temperature and time. Since the volume fraction of the precipitated carbide and intermetallic compound is determined by the alloy composition, the temperature and time conditions necessary to suppress the strength decrease are determined by C and the intermetallic compound forming element contained in the alloy. Is done.

  This invention is made | formed based on such knowledge, and the summary of the manufacturing method of the following Ni-based heat-resistant alloy pipe is summarized.

[1] The chemical composition is mass%,
C: 0.02 to 0.15%,
Si: 1.0% or less,
Mn: 2.0% or less,
P: 0.03% or less,
S: 0.01% or less,
Cr: 17.0% or more and less than 25.0%,
Ni: 45.0% to 60.0%,
Mo: 6.0% to 13.0%,
Co: 4.0 to 15.0%,
Al: more than 0.3% and 1.8% or less,
Ti: 0.02 to 0.8%,
N: 0.05% or less,
O: 0.01% or less W: 0 to 10.0%,
Zr: 0 to 0.3%,
Nb: 0 to 1.5%,
Ca: 0 to 0.05%,
Mg: 0 to 0.05%,
Rare earth elements: 0-0.2%,
Hf: 0 to 0.2%,
B: 0 to 0.02%,
V: 0 to 1.5%
Ta: 0 to 2.0%,
Re: 0 to 3.0%,
Remaining: After performing product heat treatment on the alloy tube which is Fe and impurities,
After forming a processed portion by performing hot or cold processing on a part of the alloy tube,
The coverage of the grain boundaries in the metal structure of the processed part with carbides and intermetallic compounds is 40.0% or less, and the area ratio of the carbides and the intermetallic compounds in the entire metal structure is 10.0% or less. A method for producing a Ni-based heat-resistant alloy tube, wherein post-heat treatment is performed under the conditions

[2] The chemical composition is mass%,
The method for producing a Ni-based heat-resistant alloy tube according to the above [1], which contains one or more elements selected from the elements described in (A) to (D).
(A) W: 0.5 to 10%,
(B) Ca: 0.0005-0.05%, Mg: 0.0005-0.05% and rare earth elements: 0.0005-0.2%,
(C) Hf: 0.0005-0.2%, Zr: 0.001-0.3% and B: 0.0005-0.02%,
(D) Nb: 0.01-1.5%, V: 0.02-1.5%, Ta: 0.1-2.0% and Re: 0.1-3.0%.

[3] The above-mentioned [1] or [2] Ni-based heat-resistant alloy tube manufacturing method, wherein the post-heat treatment is performed under a condition satisfying the following expression (1).
830 + 130 × log (80 × C) + 25 × (1 / 2Mo + 2.3Al + Ti) ≦ T ≦ 1260 (1)
However, T in the formula (1) means the post-heat treatment temperature (° C.), and each element symbol means the content in mass% of each element.

[4] The method for producing a Ni-based heat-resistant alloy tube according to any one of [1] to [3], wherein the post-heat treatment is performed under a condition satisfying the following expression (2).
42−0.03 × T ≦ t ≦ 360−0.2 × T (2)
However, T in the formula (2) means a post-heat treatment temperature (° C.), and t means a post-heat treatment time (minute).

  [5] The above-described processing performed on a part of the alloy pipe is performed in a range of 5.0 to 40.0% with respect to the total length of the alloy pipe at a workability of 10.0 to 50.0%. The manufacturing method of the Ni base alloy pipe in any one of 1]-[4].

  According to the present invention, when a Ni-based heat-resistant alloy tube subjected to product heat treatment is constructed, a part of the tube is bent and subjected to hot or cold processing such as expansion, thereby forming a partially processed portion in the tube. After that, even if post-heat treatment is performed on the alloy tube, it is possible to have excellent creep strength in the processed portion. For this reason, it is most suitable for manufacture of the Ni-base heat-resistant alloy pipe | tube which has a process part used as a pipe material in a power generation boiler, a chemical industry plant, etc.

Hereinafter, each requirement of the present invention will be described in detail.
1. Chemical composition of Ni-base heat-resistant alloy tube In the following description, “%” of the content of each element means “mass%”.

C: 0.02-0.15%
C is an effective and important element for forming a carbide and ensuring creep strength in a Ni-base heat-resistant alloy used at high temperatures. In order to obtain such an effect, a C content of 0.02% or more is necessary. However, if the C content becomes excessive, the carbide becomes coarse and precipitates in a large amount, leading to a decrease in creep strength. Therefore, the upper limit of the C content is 0.15%. A preferable lower limit of the C content is 0.03%, and a more preferable lower limit is 0.04%. Moreover, the upper limit with preferable C content is 0.13%, and a more preferable upper limit is 0.10%.

Si: 1.0% or less Si is an element necessary for deoxidation during steelmaking and for improving the oxidation resistance and steam oxidation resistance of the alloy. However, when the content is excessive, the hot workability of the alloy is lowered. Therefore, an upper limit is set and the Si content is set to 1.0% or less. The Si content is preferably 0.8% or less. When the deoxidation action is sufficiently ensured with other elements, it is not necessary to provide a lower limit particularly for the Si content. In order to stably obtain effects such as deoxidation, oxidation resistance and steam oxidation resistance, the Si content is preferably 0.03% or more, and more preferably 0.05% or more. preferable.

Mn: 2.0% or less Mn combines with impurity impurity S contained in the alloy to form MnS, thereby improving hot workability. On the other hand, when the content is excessive, the alloy becomes hard and brittle, and on the contrary, hot workability and weldability are impaired. Therefore, an upper limit is set and the Mn content is set to 2.0% or less. The Mn content is preferably 1.2% or less. In order to stably obtain the effect of improving hot workability, the Mn content is preferably 0.02% or more, and more preferably 0.05% or more.

P: 0.03% or less P is inevitably mixed in the alloy as an impurity, but excessive P impairs weldability and hot workability. Therefore, an upper limit is set and the P content is set to 0.03% or less. The P content is preferably 0.025% or less, and more preferably 0.02% or less. In addition, although it is desirable to reduce the content of P as much as possible, extreme reduction causes an increase in steelmaking cost. Therefore, the minimum with preferable P content is 0.0005% or more, More preferably, it is 0.0007% or more.

S: 0.01% or less S is also inevitably mixed into the alloy as an impurity in the same manner as P described above, but excessive S impairs weldability and hot workability. Therefore, an upper limit is set and the S content is set to 0.01% or less. The S content is preferably 0.008% or less, and more preferably 0.007% or less. Although the S content is desirably reduced as much as possible, extreme reduction leads to an increase in steelmaking costs. Therefore, the minimum with preferable P content is 0.0001% or more, More preferably, it is 0.0002% or more.

Cr: 17.0 to 25.0% Non-Cr is an important element for ensuring oxidation resistance, steam oxidation resistance and corrosion resistance. In order to obtain effective oxidation resistance, steam oxidation resistance and high temperature corrosion resistance in a high temperature environment of 700 ° C. or higher, it is necessary to contain 17.0% or more. The corrosion resistance increases as the Cr content increases, but if it exceeds 25.0%, the structural stability is lowered and the creep strength is impaired. Moreover, not only is it necessary to increase the content of expensive Ni in order to stabilize the austenite structure, but also the weldability is reduced. The minimum with preferable Cr content is 18.0%, and a more preferable minimum is 19.0%. The upper limit with preferable Cr content is 24.0%, and a more preferable upper limit is 23.0%.

Ni: 45.0-60.0%
Ni is an element that stabilizes the austenite structure, and is also an important element for ensuring corrosion resistance. In order to obtain these effects, the Ni content is 45.0% or more. On the other hand, excessive Ni causes not only an increase in cost but also a decrease in creep strength. Therefore, the upper limit of the Ni content is 60.0%. The minimum with preferable Ni content is 46.0%, and a more preferable minimum is 48.0%. The upper limit with preferable Ni content is 58.0% or less, and a more preferable upper limit is 56.0%.

Mo: 6.0 to 13.0%
Mo is an element that dissolves in the matrix austenite, contributes to the improvement of creep strength by solid solution strengthening, and promotes precipitation of the intermetallic compound Laves phase. In order to fully exhibit the effect, Mo needs to be contained by 6.0% or more. However, even if Mo is excessively contained, coarse intermetallic compounds may be excessively precipitated, which may lower the creep strength. Furthermore, since Mo is an expensive element, excessive Mo content causes an increase in cost. Therefore, the Mo content is 13.0% or less. The minimum with preferable Mo content is 7.0%, and a more preferable minimum is 8.0%. The upper limit with preferable Mo content is 12.0%, and a more preferable upper limit is 10.0%.

Co: 4.0 to 15.0%
Co, like Ni, is an element that stabilizes the austenite structure and contributes to the improvement of creep strength by increasing the stability of the austenite structure. In order to acquire such an effect, it is necessary to contain 4.0% or more. However, if the Co content exceeds 15.0%, the cost increases. In order to sufficiently obtain the above Co effect, the Co content is preferably 5.0% or more, and more preferably 7.0% or more. The upper limit of the content is preferably 14.0% or less, more preferably 13.0% or less.

Al: more than 0.3% and 1.8% or less Al is one of the main constituent elements of the γ′-Ni ₃ (Al, Ti) phase that has a deoxidizing action and improves the creep strength of the alloy. is there. In order to utilize the effect of precipitation strengthening by the γ′-Ni ₃ (Al, Ti) phase, it is necessary to contain more than 0.3%. On the other hand, when it is contained in a large amount, the γ′-Ni ₃ (Al, Ti) phase is excessively precipitated and the precipitate becomes coarse, so that the creep strength is lowered. Therefore, an upper limit is set so that the Al content is 1.8% or less. Note that the Al content is preferably 0.4% or more, and more preferably 0.6% or more. Further, the Al content is preferably 1.6% or less, and more preferably 1.4% or less.

Ti: 0.02 to 0.8%
Ti is an element that promotes the formation of the intermetallic compound γ′-Ni ₃ (Al, Ti) phase, contributes to precipitation strengthening in grain boundaries and grains, and is useful for improving the creep strength. In order to sufficiently exhibit these effects, a content of 0.02% or more is required. On the other hand, when the Ti content is large, the volume ratio of the intermetallic compound becomes excessive, the precipitate becomes coarse, and the creep strength is reduced. For this reason, Ti content shall be 0.8% or less. The minimum with preferable Ti content is 0.05%, and a more preferable minimum is 0.1%. The upper limit with preferable Ti content is 0.7%, and a more preferable upper limit is 0.6%.

N: 0.05% or less N has an effect of stabilizing the austenite structure, and is an element that is inevitably included in a normal dissolution method. However, a large amount of N forms a carbonitride that remains undissolved with Ti and the like together with C, thereby impairing toughness. Therefore, an upper limit is set and the N content is set to 0.05% or less. Note that the N content is preferably 0.03% or less, and more preferably 0.02 or less. There is no particular need to set a lower limit for the N content. However, since an extreme reduction significantly increases the production cost, the preferable lower limit of N is 0.0005%, and the more preferable lower limit is 0.0008%. is there.

O: 0.01% or less O (oxygen) is contained in the alloy as an impurity, and when its content is excessive, hot workability is lowered, and further, toughness and ductility are deteriorated. For this reason, an upper limit is set for the O content to 0.01% or less. The O content is preferably 0.008% or less, and more preferably 0.005% or less. Regarding the content of O, there is no particular need to set a lower limit, but since an extreme reduction leads to an increase in production cost, the preferable lower limit of the O content is 0.0005%, and the more preferable lower limit is 0.0008%. It is.

W: 0-10%
W is an element that dissolves in the matrix austenite, contributes to the improvement of the creep strength by solid solution strengthening, and promotes the precipitation Laves phase of the intermetallic compound, and has properties similar to Mo. For this reason, it may replace with a part of above-mentioned Mo and may contain W. However, when the content of W becomes large, a large amount of intermetallic compounds precipitate and the creep strength is lowered. In particular, if it exceeds 10.0%, the creep strength is significantly reduced. For this reason, when it contains W, the content shall be 10.0% or less. The W content is preferably 8% or less, and more preferably 6% or less. In order to sufficiently obtain the above effect, the content is preferably 0.5% or more, and more preferably 1.0% or more.

Zr: 0 to 0.3%
Zr mainly contributes to grain boundary strengthening and has an action of improving creep strength. For this reason, you may contain Zr as needed. In order to fully exhibit this effect, it is necessary to contain 0.001% or more. However, if the Zr content exceeds 0.3%, weldability and hot workability are impaired. Therefore, the Zr content is set to 0.3% or less. A preferable lower limit of Zr is 0.005%, and a more preferable lower limit is 0.01%. A preferable upper limit of Zr is 0.1%, and a more preferable upper limit is 0.08%.

Nb: 0 to 1.5%
Nb is an element that promotes the formation of intermetallic compounds, contributes to precipitation strengthening in grain boundaries and grains, and is useful for improving creep strength. On the other hand, when the Nb content is large, the volume fraction of the intermetallic compound becomes excessive, the precipitate becomes coarse, and the creep strength is reduced. Moreover, the weldability is lowered due to the lowering of the melting point. For this reason, when it contains Nb, the content shall be 1.5% or less. The Nb content is preferably 1.0% or less, and more preferably 0.5% or less. In order to sufficiently obtain the above effects, the Nb content is preferably 0.01% or more, and more preferably 0.05% or more.

Ca: 0 to 0.05%
Ca fixes S as a sulfide and has an effect of improving hot workability. For this reason, you may contain Ca as needed. However, if the Ca content exceeds 0.05%, the toughness, ductility and cleanliness are impaired. Therefore, when Ca is contained, its content is set to 0.05% or less. The Ca content is preferably 0.01% or less. In order to sufficiently obtain the above effects, the Ca content is preferably 0.0005% or more.

Mg: 0 to 0.05%
Mg fixes S as a sulfide and has an effect of improving hot workability. For this reason, you may contain Mg as needed. However, if the Mg content exceeds 0.05%, the toughness, ductility and cleanliness are impaired. Therefore, when it contains Mg, the content shall be 0.05% or less. The Mg content is preferably 0.01% or less. In order to sufficiently obtain the above effects, the Mg content is preferably 0.0005% or more.

Rare earth elements: 0-0.2%
The rare earth element has an effect of fixing S as a sulfide and improving hot workability. In addition, rare earth elements also have an action of forming harmless and stable oxides, reducing the undesirable influence of O (oxygen), and improving corrosion resistance, creep strength and creep ductility. For this reason, you may contain rare earth elements as needed. However, if the content of rare earth elements exceeds 0.2%, inclusions such as oxides increase, which not only impairs hot workability and weldability, but also increases costs. Therefore, when the rare earth element is contained, the content is made 0.2% or less. The rare earth element content is preferably 0.1% or less. In order to sufficiently obtain the above effects, the rare earth element content is preferably 0.0005% or more.

  The “rare earth element” is a generic name for a total of 17 elements of Sc, Y and lanthanoid, and the content of the rare earth element refers to the total content of one or more elements of the rare earth elements. Since rare earth elements are generally contained in misch metal, for example, they may be added in the form of misch metal so that the amount of rare earth elements falls within the above range.

Hf: 0 to 0.2%
Hf mainly contributes to grain boundary strengthening and has an effect of improving creep strength. For this reason, you may contain Hf as needed. However, if the Hf content exceeds 0.2%, weldability and hot workability are impaired. Therefore, when it contains Hf, the content is made 0.2% or less. The Hf content is preferably 0.1% or less, and more preferably 0.06% or less. In order to sufficiently obtain the above effect, the Hf content is preferably 0.0005% or more, and more preferably 0.001% or more.

B: 0 to 0.02%
B has an effect of improving the creep strength as a grain boundary strengthening element. For this reason, you may contain B as needed. However, if the B content exceeds 0.02%, the weldability is impaired. Therefore, when B is contained, its content is set to 0.02% or less. The B content is preferably 0.008% or less, and more preferably 0.006% or less. In order to sufficiently obtain the above effect, the B content is preferably 0.0005% or more, more preferably 0.0008% or more.

V: 0 to 1.5%
V forms carbonitride and improves high temperature strength and creep strength. For this reason, you may contain V as needed. However, when the V content exceeds 1.5%, the high temperature corrosion resistance is lowered, and further the precipitation of the σ phase which is an embrittlement phase is promoted. Therefore, when V is contained, the content is set to 1.5% or less. The V content is preferably 1.0% or less, and more preferably 0.8% or less. In order to sufficiently obtain the above effects, the V content is preferably 0.02% or more, and more preferably 0.04% or more.

Ta: 0 to 2.0%
Ta, like Ti, is an element that forms an intermetallic compound, and improves high-temperature strength and creep strength. For this reason, you may contain Ta as needed. However, if the content of Ta exceeds 2.0%, the amount of intermetallic compound precipitated becomes excessive, and the creep strength decreases. Therefore, when Ta is contained, the content is set to 2.0% or less. The Ta content is preferably 1.5% or less, and more preferably 1.3% or less. To sufficiently obtain the above effects, the Ta content should be 0.1% or more. Is preferable, and more preferably 0.3% or more.

Re: 0 to 3.0%
Re mainly improves the high temperature strength and creep strength as a solid solution strengthening element. For this reason, you may contain Re as needed. However, when the Re content exceeds 3.0%, hot workability and toughness are impaired. Therefore, when it contains Re, the content shall be 3.0% or less. The Re content is preferably 2.0% or less, and more preferably 1.5% or less. In order to sufficiently obtain the above effect, the Re content is preferably 0.1% or more, and more preferably 0.3% or more.

  The chemical composition of the Ni-base heat-resistant alloy tube, which is the object of the production method according to the present invention, contains each of the above elements within the range specified above, and the balance consists of Fe and impurities. An impurity means the component mixed by raw materials and other factors, such as an ore and a scrap, when manufacturing steel materials industrially.

2. Post-processing heat treatment (post-heat treatment)
After performing heat treatment (1150-1230 ° C.) on a Ni-based heat-resistant alloy tube having the above chemical composition, bending it to a part of the alloy tube, and performing hot or cold processing such as tube expansion, a partial A processed part is formed. In this processed portion, dislocations are introduced into the crystal grains, and the precipitation of carbides and intermetallic compounds in the crystal grains is promoted compared to the portion that has not been processed. As a result, in the actual usage environment, carbides and intermetallic compounds are coarsely and abundantly precipitated in the processed part, cracking during stress relaxation, and a relative decrease in creep strength when compared with the non-processed part. Occurs. In order to suppress these, it is effective to remove strain by post-heat treatment to suppress precipitation of carbides and intermetallic compounds. Therefore, in the present invention, post-heat treatment is always performed.

  The part of the alloy pipe is, for example, in the range of 5.0 to 40.0% with respect to the total length of the alloy pipe. And the processing degree of the process given to a part of alloy pipe is 10.0-50.0%, for example. The processing applied to a part of the alloy tube does not include processing performed over the entire length of the alloy tube such as cold rolling, cold drawing, hot rolling, hot extrusion, and hot drawing.

  Here, carbides and intermetallic compounds are likely to precipitate at grain boundaries. The carbides and intermetallic compounds precipitated at the grain boundaries of the partially processed portion during the post heat treatment are coarser than the carbides and intermetallic compounds precipitated in the actual use environment. Further, if the amount of C, W, Mo, Ti, and Nb dissolved in the austenite matrix is excessively reduced by the post heat treatment, the amount of carbides and intermetallic compounds deposited in the actual use environment decreases. As a result, the creep strength is reduced.

  In general, the coverage by the carbides and intermetallic compounds at the grain boundaries in the metal structure of the material as it is solution heat treated (hereinafter also simply referred to as “grain boundary coverage”) is 5.0% or less. When the grain boundary coverage in the partially processed part of the heat-resistant alloy tube is excessively increased by the post heat treatment, the creep strength is lowered. Therefore, the post heat treatment needs to be performed under the condition that the grain boundary coverage of the processed part is 40.0% or less. Thereby, since sufficient carbides and intermetallic compounds can be precipitated at the actual use temperature also in the processed part, it is possible to suppress a relative strength decrease compared to the non-processed part. The post heat treatment is preferably performed under the condition that the grain boundary coverage is 35.0% or less, more preferably 33.0% or less.

  In the partially processed portion, dislocations are introduced within the crystal grains, and the area ratio of the carbide and intermetallic compound precipitated in the metal structure of the processed portion to the entire metal structure is excessive. When heat treatment is performed, carbides and intermetallic compounds may be coarsely and abundantly precipitated in an actual use environment, which may cause cracking during stress relaxation and a local decrease in creep strength. Therefore, the post heat treatment is preferably performed under the condition that the area ratio of the carbide and intermetallic compound precipitated in the metal structure of the processed portion to the entire metal structure is 10.0% or less. The post heat treatment is preferably performed under the condition that the area ratio is 8.0% or less, and more preferably under the condition of 6.0% or less.

The appropriate temperature range for the post-heat treatment varies depending on the chemical composition of the alloy tube, but is preferably performed under the conditions satisfying the following formula (1).
830 + 130 × log (80 × C) + 25 × (1 / 2Mo + 2.3Al + Ti) ≦ T ≦ 1260 (1)
However, T in the formula (1) means the post-heat treatment temperature (° C.), and each element symbol means the content in mass% of each element.

  In an actual use environment, in order to suppress cracking at the time of stress relaxation and a decrease in creep strength in the partially processed portion, it is preferable that the temperature of the post heat treatment is higher. However, if the post-heat treatment temperature is too high, melt cracks are likely to occur at the grain boundaries. Therefore, the upper limit of the post heat treatment is preferably 1260 ° C. On the other hand, the lower limit of the post-heat treatment temperature needs to be determined by the relationship between the C content contained in the alloy tube and the contents of W, Mo, Nb and Ti. That is, as the C content is relatively high, the amount of precipitation of carbides increases, and as the content of W, Mo, Nb, and Ti is relatively high, the precipitation amount of intermetallic compounds increases. Therefore, the post-heat treatment temperature is preferably “830 + 130 × log (80 × C) + 25 × (1 / 2Mo + 2.3Al + Ti)” (° C.) or more.

The post-heat treatment is preferably performed under conditions that satisfy the following formula (2).
42−0.03 × T ≦ t ≦ 360−0.2 × T
However, T in the formula (2) means a post-heat treatment temperature (° C.), and t means a post-heat treatment time (minute).

  In order to remove strain by post heat treatment and suppress precipitation of carbides and intermetallic compounds, it is effective to lengthen the post heat treatment time. For this reason, in order to obtain the effect of the post heat treatment sufficiently, it is preferable to set the post heat treatment time to “42−0.03 × T” (minute) or more. On the other hand, if the post-heat treatment time is too long, the crystal grain size becomes very coarse and the creep ductility may be lowered. For this reason, it is preferable to set the post-heat treatment time to “360−0.2 × T” (minutes) or less.

  The Ni-base heat-resistant alloy tube supplied to the manufacturing method according to the present invention may be melted and cast in the same manner as a normal Ni-base heat-resistant alloy. Hereinafter, a preferable manufacturing method for obtaining the Ni-base heat-resistant alloy tube will be described.

  Prior to the final plastic working by hot or cold, heat treatment is performed in order to sufficiently dissolve precipitates in the alloy precipitated during the working. When the heating temperature in this heat treatment is lower than 1050 ° C., undissolved carbonitrides and oxides containing stable Ti and B are present in the heated alloy. On the other hand, heating to a temperature exceeding 1250 ° C. may cause high-temperature intergranular cracking and ductility reduction. Therefore, it is preferable to heat to 1050 to 1250 ° C. at least once before the final plastic working by hot or cold. A preferred lower limit is 1150 ° C and a preferred upper limit is 1230 ° C.

  In the final plastic working by hot or cold, if sufficient strain is applied, recrystallization can be promoted in the final heat treatment. In order to sufficiently impart the strain necessary for recrystallization, the final plastic working is performed at a cross-section reduction rate of 10% or more. The cross-sectional reduction rate is preferably 20% or more. The larger the cross-sectional reduction rate, the better. Therefore, the upper limit is not specified, but the maximum value in normal processing is 90%. This processing step is also a step of determining the dimensions of the product.

  When the final plastic working is performed hot, the end temperature is preferably 1000 ° C. or higher in order to avoid uneven deformation in the carbide precipitation temperature range. Moreover, although there is no special restriction | limiting in the cooling conditions after plastic processing by hot, It is preferable to cool by the fastest cooling rate. In particular, in order to suppress precipitation of coarse carbonitrides, it is preferable to cool the temperature range from the hot working end temperature to 500 ° C. at an average cooling rate of 0.25 ° C./second or more.

  When the final plastic working is performed cold, the cold working may be performed once or may be performed a plurality of times with a heat treatment step interposed therebetween. When performing cold working a plurality of times, it is sufficient that at least the final cold working cross-sectional reduction rate satisfies the above-described condition, and at least the heating temperature of the heat treatment step immediately before the final cold working may satisfy the above-described condition.

  After the final plastic working, it is preferable to heat and hold the product in a temperature range of 1050 to 1250 ° C. and then perform a product heat treatment to cool. When the heating temperature is less than 1050 ° C, sufficient recrystallization does not occur, the crystal grains become a flat processed structure, the creep strength decreases, and heating to a temperature exceeding 1250 ° C causes high-temperature grain boundary cracking or ductility reduction. Because there is.

  EXAMPLES Hereinafter, although an Example demonstrates this invention more concretely, this invention is not limited to these Examples.

  Ni-base alloys 1 to 8 and A to G having chemical compositions shown in Table 1 were melted by a high-frequency vacuum melting furnace to obtain a 17 kg ingot having an outer diameter of 100 mm. Alloys 1 to 8 and E to F in Table 1 are alloys whose chemical compositions are within the range defined by the present invention. On the other hand, comparative steel alloys A to D are alloys of comparative examples whose chemical compositions deviate from the conditions specified in the present invention.

  After heating the said ingot to 1180 degreeC, it hot-forged so that finishing temperature might be 1050 degreeC, and it was set as the hot forging board material of thickness 15mm. In addition, it was air-cooled after completion | finish of hot forging. The hot forged plate was subjected to a softening heat treatment at 1150 ° C., then cold-rolled to a thickness of 10 mm, further heated and held at 1210 ° C. for 30 minutes, and water-cooled to 25 ° C.

  Each plate material having a thickness of 10 mm is divided, and some plate materials are subjected to cold rolling with a reduction rate of the plate thickness of 20% to be a test material assuming a processed portion. It was set as the test material which assumed the process part. These test materials were post-heat treated under the conditions shown in Table 2.

  A round bar tensile test piece having a diameter of 6 mm parallel to the longitudinal direction and a gauge distance of 30 mm was produced by machining from the thickness direction center of the post-heat treated plate material, and a creep rupture test was performed. The creep rupture test was carried out using the above-mentioned test piece in the atmosphere of 700 ° C., 750 ° C. and 800 ° C., and the obtained rupture strength was regressed by the Larson-Miller parameter method. The breaking strength at 20,000 hours was determined. Table 2 also shows the creep test results.

  As shown in Table 2, in Test Nos. 1 to 8, the chemical composition and post-heat treatment conditions are within the range defined by the present invention, the creep rupture strength is good, the creep strength in the processed part and the creep of the non-processed part Deviation from strength is small. On the other hand, Test No. deviating from the conditions specified in the present invention. Nos. 9 to 15 have inferior creep rupture strength characteristics in the processed part, and the difference between the creep strength in the processed part and the creep strength in the non-processed part is large.

  According to the present invention, when a Ni-based heat-resistant alloy tube subjected to product heat treatment is constructed, a part of the tube is bent and subjected to hot or cold processing such as expansion, thereby forming a partially processed portion in the tube. After that, even if post-heat treatment is performed on the alloy tube, it is possible to have excellent creep strength in the processed portion. For this reason, it is most suitable for manufacture of the Ni-base heat-resistant alloy pipe | tube which has a process part used as a pipe material in a power generation boiler, a chemical industry plant, etc.

Claims (3)

Chemical composition is mass%,
C: 0.02 to 0.15%,
Si: 1.0% or less,
Mn: 2.0% or less,
P: 0.03% or less,
S: 0.01% or less,
Cr: 17.0% or more and less than 25.0%,
Ni: 45.0% to 60.0%,
Mo: 6.0% to 13.0%,
Co: 4.0 to 15.0%,
Al: more than 0.3% and 1.8% or less,
Ti: 0.02 to 0.8%,
N: 0.05% or less,
O: 0.01% or less,
W: 0 to 10.0%
Zr: 0 to 0.3%,
Nb: 0 to 1.5%,
Ca: 0 to 0.05%,
Mg: 0 to 0.05%,
Rare earth elements: 0-0.2%,
Hf: 0 to 0.2%,
B: 0 to 0.02%,
V: 0 to 1.5%
Ta: 0 to 2.0%,
Re: 0 to 3.0%,
The rest: After heat treatment of the alloy pipe which is Fe and impurities,
After forming a processed portion by performing hot or cold processing on a part of the alloy tube,
The coverage of the grain boundary in the metal structure of the processed part with carbide and intermetallic compound is 40.0% or less, and the area ratio of the carbide and intermetallic compound in the entire metal structure is 10.0% or less. Thus, the manufacturing method of the Ni-base heat-resistant alloy tube which performs post-heat treatment on the conditions which satisfy | fill the following (1) Formula and (2) Formula .
830 + 130 × log (80 × C) + 25 × (1 / 2Mo + 2.3Al + Ti) ≦ T ≦ 1260 (1)
42−0.03 × T ≦ t ≦ 360−0.2 × T (2)
However, T in the formula (1) represents the post heat treatment temperature (° C.), each element symbol represents the content in mass% of each element, and t represents the post heat treatment time (minutes). The chemical composition is mass%,
The manufacturing method of the Ni-base heat-resistant alloy pipe | tube of Claim 1 containing 1 or more types selected from the element described from (A) to (D).
(A) W: 0.5 to 10%,
(B) Ca: 0.0005-0.05%, Mg: 0.0005-0.05% and rare earth elements: 0.0005-0.2%,
(C) Hf: 0.0005-0.2%, Zr: 0.001-0.3% and B: 0.0005-0.02%,
(D) Nb: 0.01-1.5%, V: 0.02-1.5%, Ta: 0.1-2.0% and Re: 0.1-3.0%. The processing performed on a part of the alloy tube, the range of 5.0 to 40.0% relative to the total length of the alloy tube, carried out at 10.0 to 50.0% of working ratio, according to claim 1 or 2 The manufacturing method of the Ni-base heat-resistant alloy pipe as described in 2.