JP2861720B2 - Method for producing duplex stainless welded steel pipe excellent in strength, toughness and corrosion resistance - Google Patents

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION This invention is economically preferred, strength, toughness, and superior austenitic ferritic duplex stainless welded steel pipe corrosion resistance for use in or under C0 ₂ corrosive environments environments containing chlorides And a method for manufacturing the same.

[0002]

_2. Description of the Related Art In recent years, C0 ₂
In addition, petroleum or natural gas containing corrosive substances such as chlorides is being used, and the use of high corrosion-resistant steel pipes as steel pipes for transporting such petroleum or natural gas at high pressure is increasing. . As a high corrosion resistance steel pipe, 13Cr steel has the property of being high strength and inexpensive, but is inferior in weldability, and thus has not been used as a welded steel pipe.

Further, 13Cr steel is inferior in corrosion resistance in a high-temperature environment exceeding 100 ° C., and therefore is unsuitable for use in such a high-temperature environment. Conventionally, 18Cr-8Ni-based SUS304 has been used in a high-temperature environment. Austenitic stainless steel such as SUS316 of 16Cr-11Ni-2Mo or SUS3
Austenitic ferritic duplex stainless steels such as 29J3L and SUS329J4L are used.

However, a welded steel pipe made of austenitic stainless steel such as SUS304, SUS316, etc. has a lower strength (about 250 MPa at 0.2% proof stress) than a welded steel pipe made of 13Cr steel, and has a lower chloride content. There is a problem that the resistance to stress corrosion cracking in an environment containing many substances is inferior. On the other hand, a welded steel pipe made of duplex stainless steel is excellent in all of the properties of strength, toughness, and stress corrosion cracking resistance. However, conventional duplex stainless steel pipes made of SUS329J3L steel and SUS329J4L steel, from the viewpoint of improving corrosion resistance,
Because it contains a large amount of molybdenum and nickel,
There is a problem that the production cost is increased and the material used in an environment containing almost no hydrogen sulfide is expensive.

[0005] Therefore, inexpensive duplex stainless steels having excellent strength, toughness and stress corrosion cracking resistance have been developed. For example, Japanese Patent Application Laid-Open Nos.
Japanese Patent Publication No. 1-201446 and Japanese Patent Publication No. 4-42464 disclose a steel having a low content of molybdenum and nickel (hereinafter referred to as Prior Art 1).

On the other hand, in order to improve pitting resistance, which is important as corrosion resistance in an environment containing chlorides, it is effective to increase the contents of chromium, molybdenum and nitrogen in stainless steel. In JP-A-1-165750 and JP-A-3-82740, PI = Cr (%) + 3 × Mo () is used as an index for the pitting corrosion resistance of the duplex stainless steel, similar to the index for the austenitic stainless steel. %) + 16 × N (%) (hereinafter referred to as Prior Art 2) is disclosed.

[0007]

The above prior art 1 has the following problems. That is, for example,
As disclosed in Japanese Unexamined Patent Publication No. 3 (1993), the toughness of the duplex stainless steel deteriorates with an increase in the volume fraction (α _f ) of the ferrite phase. Therefore, the volume fraction of the ferrite phase (α _f ) is
It is usually designed to be about 0.5. However, in the case of a duplex stainless steel having an extremely small nickel content, the toughness is deteriorated even _if the volume fraction (α _f ) of the ferrite phase is about 0.5.

The prior art 2 has the following problem.
That is, since the content of each component of the ferrite phase and the austenite phase is different in the duplex stainless steel, the PI using the average composition as an indicator of the pitting corrosion resistance as in Prior Art 2 overestimates the pitting corrosion resistance. May occur. Further, even if the steel contains a considerable amount of chromium, molybdenum, and nitrogen, if the nickel content is extremely small, the pitting corrosion resistance is deteriorated.

Accordingly, object of the present invention to solve the problems described above, has a 13Cr steel equivalent 0.2% proof stress (or 400 MPa) and excellent toughness and, in an environment containing chloride or C0 ₂ It is an object of the present invention to provide a method for economically producing a welded steel pipe made of a duplex stainless steel having SUS316 steel or higher overall corrosion resistance, pitting corrosion resistance and stress corrosion cracking resistance.

[0010]

The method for manufacturing a duplex stainless steel welded pipe according to the first aspect of the present invention is characterized by the following. Carbon (C): 0.05 wt.% Or less, Silicon (Si): 1.5 wt.% Or less, Manganese (Mn): 2.0 wt.% Or less, Nickel (Ni): 3.0-5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, and the rest: Consisting of Fe and unavoidable impurities, Equation (1) below: σ _0.2 = 20 × Cr (%) + 11 × Mn (%) −Ni (%) + 133 × N (%) − 38 The index (σ _0.2 ) indicating the 0.2% proof stress obtained by ───── (1) is 400
A steel ingot or a slab having the chemical composition described above is prepared, and the steel ingot or the slab is hot-rolled to prepare a steel sheet. An original plate is prepared by performing an annealing consisting of heating at a temperature within the range of T (° C.) or less and then cooling, T = 71 × Ni (%) + 6 × Mn (%) − 36 × Cr (%) − 42 × Si (%) + 1037 × C (%) + 1113 × N (%) + 1608───── (2) Then, the raw plate was formed to prepare a base tube, and the seam portion of the base tube was prepared. And the weld metal of the seam is carbon (C): 0.05 wt.% Or less, silicon (Si): 1.5 wt.% Or less, manganese (Mn): 2.0 wt.% Or less, nickel (Ni): 3.0 to 5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, Oxygen (O): 0.035 wt.% Or less, and balance: Fe and unavoidable impurities Consisting of
A welded steel pipe having a chemical composition having an index of 2% proof stress (σ _0.2 ) of 400 or more is prepared, and then the welded steel pipe is heated at a temperature in the range of 900 to 1150 ° C.
A solution heat treatment consisting of cooling at a rate of 1 ° C./sec or more in a temperature range of up to 850 ° C .;
Volume fraction alpha _f of the ferrite phase of the base material of the welded steel pipe is 0.40
0.60.60, and the base material described below in (3)
The pitting resistance index of the ferrite phase (α
_P1 ) and the pitting resistance index (γ _P1 ) of the austenite phase determined by the following equation (4) are 23.5 or more, and α _P1 = 23 × Cr (%) / (3 × α _f +20)) ────────────── (3) γ _P1 = 20 × Cr (%) / (3 × α _f +20) −16 × N (%) / (α _f -1) ─ ─── (4) The volume fraction (α _f ) of the ferrite phase of the seam weld metal of the welded steel pipe is in the range of 0.30 to 0.50, and the pitting corrosion resistance of the ferrite phase of the weld metal is Index (α _P1 ) and pitting resistance index of austenite phase (γ
_P1 ) produces a duplex stainless steel welded pipe in which the base material is larger than a lower value of either α _P1 or γ _P1 of the base material.

[0011] The method for producing a duplex stainless steel welded pipe according to the second aspect of the invention is characterized by the following. Carbon (C): 0.05 wt.% Or less, Silicon (Si): 1.5 wt.% Or less, Manganese (Mn): 2.0 wt.% Or less, Nickel (Ni): 3.0-5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, and the rest: Consisting of Fe and unavoidable impurities. Equation (1) below: σ _0.2 = 20 × Cr (%) + 11 × Mn (%) − An index (σ _0.2 ) indicating the 0.2% proof stress obtained by Ni (%) + 133 × N (%) − 38 ────── (1) is 400
A steel ingot or a steel slab having the chemical composition described above is prepared, and the steel ingot or the steel slab is heated at a temperature of T (° C.) or less obtained by the following equation (2), and then 900 ° C. An original plate is prepared by performing hot rolling including finishing at the above temperature, and T = 71 × Ni (%) + 6 × Mn (%) − 36 × Cr (%) − 42 × Si (%) + 1037 × C (%) + 1113 × N (%) + 1608 ───── (2) Then, the raw plate was formed to prepare a raw tube, and a seam portion of the raw tube was welded. , Carbon (C): 0.05 wt.% Or less, Silicon (Si): 1.5 wt.% Or less, Manganese (Mn): 2.0 wt.% Or less, Nickel (Ni): 3.0-5.0 wt.%, Chromium (Cr) : 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, Oxygen (O): 0.035 wt.% Or less, and the balance: Fe and unavoidable impurities.
A welded steel pipe having a chemical composition having an index (σ _0.2 ) of 2% proof stress of 400 or more is prepared, and then the welded steel pipe is heated at a temperature in the range of 900 to 1150 ° C.
A solution heat treatment consisting of cooling at a rate of 1 ° C./sec or more in a temperature range of up to 850 ° C .;
Volume fraction alpha _f of the ferrite phase of the base material of the welded steel pipe is 0.40
0.60.60, and the base material described below in (3)
The pitting resistance index of the ferrite phase (α
_P1 ) and the pitting resistance index (γ _P1 ) of the austenite phase determined by the following equation (4) are 23.5 or more, and α _P1 = 23 × Cr (%) / (3 × α _f +20)) ────────────── (3) γ _P1 = 20 × Cr (%) / (3 × α _f +20) −16 × N (%) / (α _f -1) ─ ─── (4) The volume fraction (α _f ) of the ferrite phase of the seam weld metal of the welded steel pipe is in the range of 0.30 to 0.50, and the pitting corrosion resistance of the ferrite phase of the weld metal is Index (α _P1 ) and pitting resistance index of austenite phase (γ
_P1 ) produces a duplex stainless steel welded pipe in which the base material is larger than a lower value of either α _P1 or γ _P1 of the base material.

[0012]

The chemical composition of the base plate and the weld metal and the volume fraction of the ferrite phase in the method of the present invention are described below.
The reason why (α _f ) is limited to the above range will be described. (1) Carbon (C): Carbon is an austenite-forming element. However, when the carbon content exceeds 0.05 wt.%, Carbides are formed in the steel, and the corrosion resistance is deteriorated. Therefore, the carbon content of the original plate and the weld metal should be limited to 0.05 wt.% Or less.

(2) Silicon (Si): Silicon is an element useful as a deoxidizing material. However, when the silicon content exceeds 1.5 wt.%, The susceptibility to cracking at the time of welding increases, and an intermetallic compound is formed in the steel to deteriorate hot workability. Therefore, the silicon content of the original plate and the weld metal should be limited to 1.5 wt.% Or less.

(3) Manganese (Mn): Manganese is an austenite-forming element and has a deoxidizing effect. However, when the manganese content exceeds 2.0 wt.%, The pitting corrosion resistance in an environment containing chloride deteriorates.
Therefore, the manganese content of the base plate and the weld metal is 2.0 w
Should be limited to t.% or less.

(4) Nickel (Ni): Nickel is a strong austenite-forming element, and is an essential element for obtaining a volume fraction (α _f ) of a ferrite phase described later.
However, when the nickel content is less than 3.0 wt.%, The ductility, toughness, and pitting corrosion resistance are significantly deteriorated even when the volume fraction (α _f ) of the ferrite phase is adjusted to an appropriate value described later. On the other hand, if the nickel content exceeds 5.0 wt.%, Not only does the cost rise, but also the content of nitrogen, which is an austenite forming element, in order to adjust the volume fraction (α _f ) of the ferrite phase to an appropriate value. As a result, it may be disadvantageous from the viewpoint of pitting corrosion resistance. Therefore, the nickel content of the original plate and the weld metal should be limited to the range of 3.0 to 5.0 wt.%.

(5) Chromium (Cr): Chromium is a strong ferrite-forming element and has an effect of improving the general corrosion resistance and the pitting corrosion resistance. However, if the chromium content is less than 21.0 wt.%, The volume fraction (α _f ) of the ferrite phase cannot be adjusted to an appropriate value described later without forming martensite. On the other hand, the chromium content is 25.0
If it exceeds wt.%, the toughness decreases and the σ phase precipitates, so that the corrosion resistance and the hot workability deteriorate. Therefore,
The chromium content of the blank and weld metal should be limited to the range of 21.0-25.0 wt.%.

(6) Nitrogen (N): Nitrogen is a strong austenite-forming element and an effective element for imparting pitting corrosion resistance to steel. However, the nitrogen content is 0.25wt.%
If the ratio exceeds, the deformation resistance during hot rolling increases, so that cracks and the like occur in the original sheet, and blow occurs during welding. Therefore, the nitrogen content of the base plate and weld metal is 0.
Should be limited to 25 wt.% Or less.

(7) Oxygen (O): As the oxygen content increases, toughness and pitting resistance deteriorate. Therefore, the oxygen content of the original stainless steel is 0.01
wt.% or less. However, the seam portion of a stainless steel welded pipe is often welded by a submerged arc using a low basicity flux from the viewpoint of improving efficiency. Therefore, an increase in the oxygen content of the weld metal is inevitable, but from the viewpoint of preventing the deterioration of the toughness and pitting resistance of the weld metal, the oxygen content of the weld metal is 0.03%.
Should be limited to 5 wt.% Or less.

(8) Volume fraction of ferrite phase (α _f ): In duplex stainless steel, volume fraction of ferrite phase (α _f )
Has a great effect on the properties of steel. That is, when the volume fraction (α _f ) of the ferrite phase is less than 0.40, hot rollability is significantly deteriorated, and when α _f exceeds 0.60, ductility and toughness of the steel are reduced. Therefore, the volume fraction (α _f ) of the ferrite phase of the original plate should be limited to the range of 0.40 to 0.60. On the other hand, since the crystal grains of the weld metal are significantly coarser than the original sheet, the volume fraction of the ferrite phase (α _f )
Exceeds 0.50, the ductility and toughness of the weld metal deteriorate. If the volume fraction (α _f ) of the ferrite phase is less than 0.30, the stress corrosion cracking resistance is reduced. Therefore, the volume fraction (α _f ) of the ferrite phase of the weld metal should be limited to the range of 0.30 to 0.50.

The present inventor has proposed that molybdenum-free 2
The effect of components on the strength of duplex stainless steel was studied in detail. As a result, it was found that an index (σ _0.2 ) indicating the 0.2% proof stress of the duplex stainless steel containing no molybdenum is represented by the following equation (1). Therefore, in the present invention, the same as 13Cr steel
In order to provide 0.2% proof stress, the contents of chromium, manganese, nickel and nitrogen in the original plate and the weld metal should be limited so that the value obtained by the following equation (1) is 400 or more. σ _0.2 = 20 × Cr (%) + 11 × Mn (%)-Ni (%) + 133 × N (%)-38 ──── (1)

On the other hand, the pitting corrosion resistance in an environment containing chlorides was repeatedly examined. As a result, in the duplex stainless steel, the pitting corrosion resistance was also reduced because the contents of the ferrite phase and the austenite phase were different. It was found that, unlike the ferrite phase and the austenitic phase, pitting occurred initially in a phase having poor pitting resistance.

In the duplex stainless steel containing no molybdenum, the pitting resistance index (α _P1 ) in the ferrite phase and the pitting resistance index (γ _P1 ) in the austenite phase are determined by the respective contents of chromium and nitrogen and ferrite. It depends on the volume fraction of the phase (α _f ). Therefore,
In order to impart pitting corrosion resistance equivalent to that of SUS316 steel to duplex stainless steel, the pitting corrosion resistance index (α _P1 ) in the ferrite phase determined by the following equation (3) and
The index of pitting corrosion resistance (γ _P1 ) in the austenite phase determined by the equation (4) needs to be 23.5 or more. α _P1 = 23 × Cr (%) / (3 × α _f +20) ──────────────── (3) γ _P1 = 20 × Cr (%) / (3 × _{α f + 20) - 16 ×} N (%) / (α f -1) ──── (4)

From the above, the chromium and nitrogen contents of the original sheet were determined so that both the index of pitting resistance (α _P1 ) in the ferrite phase of the original sheet and the index of pitting corrosion resistance (γ _P1 ) in the austenite phase were 23.5 or more. The amount and volume fraction of the ferrite phase (α _f ) should be limited. And when the pitting corrosion resistance of the weld metal is lower than the pitting corrosion resistance of the original sheet,
Since pitting occurs preferentially in the weld metal, the pitting resistance index (α _P1 ) in the ferrite phase and the pitting resistance index (γ _P1 ) in the austenite phase of the weld metal are expressed as follows:
The chromium and nitrogen contents of the weld metal and the volume fraction of the ferrite phase (α _f ) should be limited so as to be larger than the lower value of either α _P1 or γ _P1 of the original sheet.

In the method according to the first embodiment of the present invention, a steel ingot or a slab having the above-mentioned chemical composition is hot-rolled to prepare a steel sheet. 2) An original plate is prepared by annealing at a temperature within a range of T (° C.) or less obtained by the formula and then cooling, and a base plate is prepared. T = 71 × Ni (%) + 6 × Mn (%) −36 × Cr (%) −42 × Si (%) ＋ 1037 × C (%) + 1113 × N (%) + 1608 ───── (2) Next, the original plate is formed to prepare the base tube, and the seam of the base tube is prepared. A welded steel pipe is prepared by welding the parts. Next, a solution heat treatment is performed on the welded steel pipe thus prepared under predetermined conditions.

In the above-described process, during annealing,
If the heating temperature for the steel sheet is less than 900 ° C, the original sheet will not be sufficiently softened, and it will be difficult to form the tube subsequently. On the other hand, if the heating temperature during annealing is high, the volume fraction of the ferrite phase (α _f ) increases, and if the α _f exceeds 0.60,
Not only is the ductility reduced, but also significant wrinkles may occur on the original plate during the molding of the tube. Then, the present inventors examined in detail the relationship between the volume fraction of the ferrite phase (α _f ), the heating temperature, and the composition of the chemical components. As a result, the maximum heating rate at which the volume fraction of the ferrite phase (α _f ) was 0.60 or less was obtained. Temperature T (° C)
Can be expressed by the above equation (2). Therefore, the heating temperature of the steel sheet during annealing should be limited to a range of 900 ° C. or more and T (° C.) or less obtained by the above equation (2).

In the method according to the second embodiment of the present invention, a steel ingot or a billet having the above-mentioned chemical composition is heated at a temperature not higher than T (° C.) determined by the above equation (2), A raw plate is prepared by performing hot rolling consisting of finishing at a temperature of not less than ℃, and then, as in the method of the first embodiment, the raw plate is formed to prepare a raw tube, and a seam portion of the raw tube is welded. Thus, a welded steel pipe such as a UOE steel pipe is prepared, and then the solution-treated heat treatment is performed on the thus prepared welded steel pipe under predetermined conditions. According to the method of the second embodiment, the hot rolled steel sheet can be formed into a tube without performing softening annealing, and therefore, the manufacturing cost can be further reduced.

The present inventors have repeatedly studied the conditions under which a hot rolled steel sheet can be formed into a tube without performing softening annealing. As a result, the following was found.
That is, in the steel of the present invention, the volume ratio (α _f ) of the ferrite phase of the hot-rolled steel sheet is substantially the same as the volume ratio (α _f ) of the ferrite phase when the ingot or the billet is heated.

Therefore, in order to prevent the occurrence of wrinkles during the forming of the tube, the heating temperature for the steel ingot or the slab during the hot rolling is adjusted by setting the volume ratio (α _f ) of the ferrite phase to 0.60.
The temperature should be limited to T ° C. or less obtained by the above equation (2). Also, sufficient dynamic recrystallization or recovery must occur during rolling to omit soft annealing. Therefore, the finishing temperature of hot rolling should be limited to 900 ° C or higher.

By welding a seam portion of a raw tube prepared by forming an original plate into a tube, a weld metal of a duplex stainless steel and a heat-affected zone of a base metal are formed by precipitation of carbonitride and ferrite phase. The volume fraction (α _f ) of increases significantly. Therefore, in order to eliminate the precipitated carbonitride and optimize the volume ratio (α _f ) of the ferrite phase, it is necessary to perform a solution treatment on the welded steel pipe to which the seam is welded.

The solution heat treatment described above must be performed under the conditions of heating at a temperature in the range of 900 to 1150 ° C. and cooling at a rate of 1 ° C./sec or more in a temperature range of 500 to 850 ° C. There is. That is, the heating temperature is 900 ℃
If it is less than 1, the carbonitride does not form a solid solution, so that the corrosion resistance deteriorates. On the other hand, when the heating temperature exceeds 1150 ° C., the volume ratio (α _f ) of the ferrite phase increases and the crystal grains become coarse, so that ductility and toughness deteriorate.

Chromium carbonitride is generally 500-850 ° C.
Precipitates in a temperature range within the range of However, in the solution heat treatment, if the cooling rate of the welded steel pipe heated under the above-described conditions in the temperature range of 500 to 850 ° C. is less than 1 ° C./sec, the particles accompanying the precipitation of chromium carbonitride will not be obtained. Interfacial corrosion becomes significant. Therefore, in the solution heat treatment, 50
The cooling rate in the temperature range between 0 and 850 ° C should be limited to 1 ° C / sec or more.

[0032]

Next, the present invention will be described with reference to examples and comparative examples. Example 1 A steel ingot having a chemical composition within the range of the present invention shown in Table 1 and having a chemical composition within the range of the present invention was weighed 50 kg, and at least one of the test ingots had a chemical composition outside the range of the present invention. Test steel a '
A steel ingot having a weight of gl ′ and a weight of 50 kg was prepared, and a commercially available SUS316 steel plate and a commercially available 13Cr steel plate were prepared as conventional steels.

Next, the ingots of the test steels a to i and the comparative test steels a ′ to l ′ were heated at a temperature of 1150 ° C.
The steel sheet was hot-rolled at a finishing temperature of 500 ° C., and then air-cooled at a cooling rate of 0.7 ° C./sec in a temperature range of 500 to 850 ° C. to prepare a steel sheet having a thickness of 15 mm. The steel sheet thus prepared was heated at a temperature of 1050 ° C.,
A solution heat treatment comprising cooling at a rate of 30 ° C./sec in a temperature range of 850 ° C. was performed. Thus, the volume fraction (α _f ) of the ferrite phase within the range of the present invention shown in Table 1, the index (σ _0.2 ) indicating the 0.2% proof stress obtained by the equation (1), and the index (σ _0.2 ) obtained by the equation (3) Specimens Nos. 1 to 9 of the present invention as original plates having a pitting corrosion resistance index of a ferrite phase (α _P1 ) and a pitting corrosion resistance index of an austenite phase (γ _P1 ) determined by equation (4) and a comparative specimen N
o.1-12 were prepared.

[0034]

[Table 1]

The above-described specimen No. 1 of the present invention as an original plate
~ 9, Comparative specimen No. 1 '~ 12' and from each of the conventional steel specimen, microstructure observation test specimen, tensile test specimen,
A 2 mm V notched Charpy impact test specimen, a general corrosion test specimen, a stress corrosion cracking test specimen and a test specimen for pitting corrosion potential measurement were taken, and 0.2% proof stress, tensile strength, elongation, 0 ° C and −50
The absorption energy at ℃, the overall corrosion resistance, the pitting potential, the stress corrosion cracking resistance (σ _th ) and the state of the end face cracks of the steel sheet were examined. The results are shown in Table 2.

[0036]

[Table 2]

The volume fraction (α _f ) of the ferrite phase is
The micro sample was subjected to 20% NaOH electrolytic etching and measured. 3% NaCl-10 at 200 ° C
In AtmC0 ₂ conditions, it was assessed by measuring the corrosion rate when the strip-shaped test piece having a thickness of 3mm × width 20 mm × length 30mm was immersed for 720 hours. 45% MgC stress corrosion cracking resistance
l ₂ Perform constant load tensile test in boiling solution for up to 500 hours,
The maximum stress that did not break after 500 hours
(σ _th ). The pitting potential was a potential at which the current density became 100 μA / cm ^{2 according} to JIS GO577. Then, the degree of end face cracking of the steel sheet was visually determined.

As apparent from Tables 1 and 2, 0.2%
The comparative sample No. 1 having an index (σ _0.2 ) indicating the proof stress, a pitting corrosion resistance index (α _P1 ) of the ferrite phase, and a pitting corrosion resistance index (γ _P1 ) of the austenite phase which are all out of the range of the present invention. And No. 2 was inferior in 0.2% proof stress and pitting potential. The comparative samples No. 3 and No. 5 in which the pitting corrosion resistance index (α _P1 ) of the ferrite phase was out of the range of the present invention were inferior in pitting potential. Comparative specimens No. 4 and N having a small ferrite phase volume fraction (α _f ) outside the scope of the present invention.
In o.6, the end face crack of the steel sheet was large, and in particular, the comparative specimen No.4 in which the volume fraction (α _f ) of the ferrite phase was less than 0.30,
The stress corrosion cracking resistance (σ _th ) was also inferior.

Comparative samples No. 7 and No. 8 having a small nickel content outside the range of the present invention had inferior absorption energy and pitting potential. The comparative samples No. 9 and No. 12 in which the pitting resistance index (γ _P1 ) of the austenite phase was low outside the range of the present invention were inferior in pitting potential. Comparative specimens No. 10 and No. 11 in which the volume fraction (α _f ) of the ferrite phase was out of the range of the present invention were inferior in elongation and energy absorption.

The conventional SUS316 steel sheet was inferior in 0.2% proof stress and stress corrosion cracking resistance (σ _th ). The 13Cr steel sheet was inferior in elongation, absorbed energy, general corrosion resistance and pitting corrosion resistance.

On the other hand, in the test specimens Nos. 1 to 9 of the present invention, the 0.2% proof stress was 400 MPa or more, which corresponded to the σ _0.2 obtained by the equation (1), and was excellent in ductility and toughness. I was In addition, both the general corrosion resistance under CO ₂ corrosive environment and the pitting corrosion resistance under chloride-containing environment are SUS3
It was 16 steel or more, and the stress corrosion cracking resistance was remarkably superior to SUS316 steel.

FIG. 1 is a graph showing the relationship between the nickel content, the pitting potential and the absorbed energy at 0 ° C. and -50 ° C., showing the specimens Nos. 2, 3, and 4 of the present invention and the nickel content. Is less than the scope of the present invention, a comparative sample No. 7 and
Fig. 8 shows the pitting potential and the absorbed energies at 0 ° C and -50 ° C. As is clear from FIG. 1, the pitting potentials of the comparative specimens Nos. 7 and 8 having a nickel content of less than 3 wt.
Its pitting resistance index of the ferrite phase (alpha _P1) and the austenitic phase pitting resistance index (gamma _P1) is low even within the scope of the present invention, pitting corrosion resistance was poor. The absorption energies of Comparative Sample Nos. 7 and 8 at 0 ° C. and -50 ° C. were low even though the ferrite phase volume ratio was within the range of the present invention, and the toughness was poor.

FIG. 2 shows the volume fraction (α _f ) of the ferrite phase.
5 is a graph showing the relationship between absorbed energy at 0 ° C. and -50 ° C., elongation, and cracks at the end face of a steel sheet. The specimens of the present invention Nos. 1 to 9 and the comparative specimens Nos. 4, 6 and No.10,11
It shows the absorbed energy at 0 ° C and -50 ° C, elongation, and end face cracking of the steel sheet. As is clear from FIG.
When the volume fraction (α _f ) of the ferrite phase was less than 0.40, the cracks at the steel sheet end face during rolling of comparative test pieces Nos. 4 and 6 were large,
The comparative specimens Nos. 10 and 11, in which the volume fraction (α _f ) of the ferrite phase exceeded 0.60, had low absorption energies and elongation at 0 ° C and -50 ° C, and were inferior in toughness and ductility.

Example 2 A steel ingot having a chemical composition within the scope of the present invention shown in Table 1 and having a chemical composition within the range of the present invention shown in Table 3 was used. After hot rolling at the heating temperature and the finishing temperature, or after hot rolling the steel ingot, annealing at a temperature within the range of the present invention, a 15 mm thick specimen N of the present invention shown in Table 3
o.10 to 15 were prepared. On the other hand, the weight of the test steels e, f, g
A 0 Kg steel ingot is hot-rolled at a heating temperature or a finishing temperature outside the range of the present invention also shown in Table 3, or after the above-described steel ingot is hot-rolled, annealed at a temperature outside the range of the present invention. And Table 3
The test specimens Nos. 13 to 18 having a thickness of 15 mm shown in FIG. Note that cooling after hot rolling or after annealing was air cooling.

[0045]

[Table 3]

The specimen N of the present invention prepared as described above
Samples for micro-observation were collected from o.10 to 15 and comparative specimens Nos.13 to 18, and the volume fraction of the ferrite phase (α
_f ) and hardness (HV: 10 kgf) were measured, and the measurement results are shown in Table 3.

As is clear from Table 3, the comparative samples No. 13 and No. 1 were subjected to an annealing treatment after hot rolling, in which the annealing temperature was higher than the range of the present invention.
In No. 6, the volume fraction (α _f ) of the ferrite phase was out of the range of the present invention, and there was a high possibility that wrinkles would occur during tube formation. And, the comparative sample No. 18 whose annealing temperature was out of the range of the present invention was low in hardness (HV: 10 kgf), so that it was considered difficult to form a tube by UOE.

In the case where the annealing treatment after the hot rolling was not performed, and the finishing temperature during the hot rolling was low outside the range of the present invention, the comparative sample No. 14 had a hardness (HV: 10 kgf). Therefore, it was considered that tube formation by UOE was difficult. And
In Comparative Samples No. 15 and No. 17, the heating temperature during hot rolling was out of the range of the present invention and the volume ratio (α _f ) of the ferrite phase was out of the range of the present invention, and The possibility of wrinkles occurring during formation was high.

On the other hand, the annealing temperature in the case where the annealing treatment is performed after the hot rolling, and the heating temperature and the finishing temperature in the hot rolling in the case where the annealing treatment is not performed after the hot rolling are as follows: Inventive specimen No. 1 which is within the scope of the present invention
0 to 15, the volume fraction of the ferrite phase (α _f ) is within the range of the present invention, and has an appropriate hardness (HV: 10 kgf);
Therefore, it is considered that the tube formation by UOE can be favorably performed.

Example 3 A steel ingot of test steel j having a chemical composition within the range of the present invention shown in Table 4 was hot-rolled into a 20-mm-thick steel sheet in a hot-rolling mill. Butt the obtained two original plates,
The seam was welded by two-electrode submerged arc welding using the wire having the chemical composition shown in Table 4 and the flux having the chemical composition shown in Table 6 under the conditions shown in Table 5. Thus, the two types of welded steel sheets whose chemical composition of the original sheet and the weld metal of the seam portion are within the scope of the present invention,
In addition, two types of welded steel sheets were manufactured in which the oxygen content of the weld metal was out of the range of the present invention and large. The groove shape is
The inner depth was 5.5 mm, the outer depth was 7 mm, and the bevel angle was 45 °.

The above-mentioned two types of welded steel sheets in which the chemical compositions of the original plate and the weld metal are within the scope of the present invention, and two types of welded steel sheets in which the oxygen content of the weld metal is large outside the scope of the present invention Heating at a temperature of 1050 ° C.
A solution heat treatment consisting of cooling at a rate of 20 ° C./sec in a temperature range of 500-850 ° C., thus
The base plate and the weld metal have a chemical composition within the range of the present invention, a volume fraction of the ferrite phase (α _f ), an index indicating 0.2% proof stress (σ _0.2 ), a pitting corrosion resistance index (α _P1 ) of the ferrite phase and austenite. Having a phase pitting resistance index (γ _P1 ),
Test specimens No. 16 and No. 17 of the present invention shown in Table 7 and Comparative test specimens No. 19 and No. 20 having a large oxygen content outside the range of the present invention were produced.

[0052]

[Table 4]

[0053]

[Table 5]

[0054]

[Table 6]

[0055]

[Table 7]

[0056]

[Table 7]

From the specimens Nos. 16 and 17 of the present invention and the comparative specimens Nos. 19 and 20 described above, a micro observation sample, a Charpy impact test specimen, and a grain boundary corrosion test specimen (thickness: 3)
mm × width 20mm × length 30mm) and pitting test specimen (thickness 3mm ×
(Width 40 mm x length 40 mm).

In order to examine the effect of solution heat treatment conditions on the welded steel sheet, the specimen No. 16 of the present invention was heated to a temperature in the range of 850 to 1200 ° C. and then heated to 500 to 85 ° C.
A solution heat treatment comprising cooling at a rate of 0.3 to 30 ° C./sec in a temperature range of 0 ° C. was performed, and the above-described test pieces were collected.

The Charpy impact test specimen was sampled so that the center of the notch was located at the center of the weld metal, and the intergranular corrosion test specimen and the pitting corrosion test specimen were taken so that the center was located at the center of the weld metal. Was collected. The intergranular corrosion resistance is determined by JI
Performed by a 65% nitric acid corrosion test according to S G0573 (hereinafter referred to as “Huey test”), and its pitting corrosion resistance was determined by a ferric chloride test at 20 ° C according to JIS G0578 (hereinafter referred to as “pitting corrosion test”). Was.

FIG. 3 shows the relationship between the oxygen content of the weld metal, the absorbed energy of the weld metal at −20 ° C., and the corrosion rate of the weld in a pitting corrosion test. As is clear from FIG. 3, the comparative specimens No. 19 and No. 20 in which the oxygen content of the weld metal is larger than the range of the present invention, the toughness and pitting corrosion resistance of the weld metal deteriorated.

FIG. 4 shows the heating temperature during the solution heat treatment, the corrosion rate of the weld in the Huey test, the absorbed energy of the weld metal at −20 ° C., and the volume fraction of the ferrite phase (α
_f ). As is clear from FIG. 4, when the heating temperature during the solution heat treatment was out of the range of the present invention and low, the intergranular corrosion resistance and toughness of the weld metal deteriorated. Further, when the heating temperature during the solution heat treatment is higher than the range of the present invention, the volume fraction (α _f ) of the ferrite phase of the weld metal increases outside the range of the present invention, and the toughness is significantly deteriorated. .

FIG. 5 shows the relationship between the cooling rate in the temperature range of 500 to 850 ° C. and the corrosion rate in the Huey test of the weld, that is, the intergranular corrosion resistance. As is clear from FIG. 5, if the cooling rate in the temperature range of 500 to 850 ° C. during the solution heat treatment is less than 1 ° C./sec, even if the heating temperature is 1050 ° C., the intergranular corrosion resistance of the weld metal is reduced. The properties were significantly reduced.

Example 4 The steel ingots of the test steels j and k having the chemical composition within the range of the present invention shown in Table 8 were obtained at a hot rolling mill, and the ingots of the present invention are shown in Table 9 by the symbols 1 to 3 in Table 9. Hot rolling at a heating temperature and a finishing temperature within the range, or after hot rolling the above steel ingot, annealing at a temperature within the range of the present invention, the present invention specimen No. .18-20 were produced. For comparison, comparative specimens Nos. 21 to 24 made of a steel plate having a thickness of 20 mm, in which the heating temperature or the finishing temperature during hot rolling is out of the range of the present invention indicated by reference numerals in Table 9, are shown. Manufactured.

[0064]

[Table 8]

[0065]

[Table 9]

The above-described specimens Nos. 18 to 20 of the present invention and comparative specimens Nos. 21 to 24 were formed by UOE processing into the same tubes as those used for manufacturing a welded steel pipe having an outer diameter of 20 inches. Table 9 also shows the tube formability at that time. As is clear from Table 9, the comparative specimen No. 21 in which the annealing treatment was performed after the hot rolling and the annealing temperature was out of the range of the present invention, wrinkles were generated at the time of forming the tube. Occurred. And, a comparative specimen having a low annealing temperature outside the range of the present invention.
In No.24, tube forming was not possible.

In the case where the annealing treatment after the hot rolling was not performed, and the heating temperature during the hot rolling was out of the range of the present invention, the comparative sample No. 23 had wrinkles when forming the tube. did.
The comparative sample No. 22 having a low finishing temperature outside the range of the present invention was not capable of forming a tube. On the other hand, the annealing temperature in the case where the annealing treatment is performed after the hot rolling, and the heating temperature and the finishing temperature in the hot rolling in the case where the annealing treatment is not performed after the hot rolling are all the same. The test piece Nos. 18 to 20 of the present invention within the scope of the present invention had extremely good tube formability.

Example 5 The ingots of the test steels j and k having the chemical composition within the range of the present invention shown in Table 8 were prepared under the conditions within the range of the present invention.
The raw plate having a thickness of 20 mm and the obtained original plate having a thickness of 20 mm were formed into a tube by UOE processing to prepare a raw tube. The seam portion of the tube was welded by two-electrode submerged arc welding using the wire C or E shown in Table 8 and the flux a shown in Table 6. Thus, the chemical composition of the weld metal, the volume fraction of the ferrite phase (α _f ), an index indicating 0.2% proof stress (σ _0.2 ), the pitting resistance index of the ferrite phase (α _P1 ), and the pitting resistance index of the austenite phase (α _P1 ) γ _P1 ) are all within the scope of the present invention.
Specimens Nos. 21 to 23 of the present invention shown in Table 10 were prepared from welded steel pipes having an outer diameter of inches.

For comparison, an original plate having a chemical composition and production conditions within the scope of the present invention was welded using a flux a shown in Table 6 and wires A to D shown in Table 8, Volume fraction of ferrite phase of the weld metal (α _f ),
One of the index indicating 0.2% proof stress (σ _0.2 ), the pitting corrosion resistance index of the ferrite phase (α _P1 ) and the pitting corrosion resistance index of the austenite phase (γ _P1 ) is out of the scope of the present invention. Comparative specimens Nos. 25 to 28 shown in Table 10 were prepared, each of which was made of an outer diameter welded steel pipe.

[0070]

[Table 10]

The specimens Nos. 21 to 23 of the present invention and the comparative specimens Nos. 25 to 28 were heated at a temperature of 1050 ° C.
A solution heat treatment comprising cooling at a rate of 20 ° C./sec in a temperature range of 0 to 850 ° C. was performed. From each of the test pieces subjected to the solution heat treatment, a tensile test piece,
Charpy impact test specimen, general corrosion test specimen (thickness 3mm × width
20mm x 30mm length), pitting test specimen (thickness 3mm x width 40mm x length 40mm) and stress corrosion cracking test specimen (thickness 2mm x width 10mm)
× length 115 mm) was collected so that the center of each test piece was the center of the weld metal. In addition, a test piece similar to the above was collected from the base material of each specimen.

The tensile strength,-
The absorption energy at 20 ° C., the overall corrosion resistance, the pitting resistance, and the stress corrosion cracking resistance were measured, and the measurement results are shown in Table 11. The stress corrosion cracking resistance was measured as follows. That is, the two glass restraining tools 2 and 2 for restraining the longitudinal ends of the test piece A from the upper surface, and the longitudinal middle portion of the test piece A, as shown in the schematic front view in FIG. A four-point bending jig 1 composed of two glass push-up tools 3 and 3 spaced apart from each other at a predetermined interval was used. 4 above
The test piece A is bent at four points by the point bending jig 1 so that the test piece A
0.5% strain was applied to the surface of
Was evaluated by the presence or absence of stress corrosion cracking that occurred in test piece A when immersed in 150% of 20% NaCl-10 atm CO ₂ for 720 hours (hereinafter referred to as a four-point bending test).

[0073]

[Table 11]

As is clear from Tables 10 and 11,
The pitting corrosion resistance index (α _P1 ) of the ferrite phase of the weld metal or the pitting corrosion resistance index (γ _P1 ) of the austenitic phase is determined by the α _{P1 of the} base metal.
And test specimen N smaller than the lower value of γ _P1
In o.25 and No.28, pitting occurred in the weld metal, and the pitting resistance was poor. Volume fraction of ferrite phase of weld metal (α
_f ) is less than the range of the present invention, and comparative specimen No. 26
In the four-point bending test, cracks occurred in the weld metal, and the stress corrosion cracking resistance was poor. In Comparative Sample No. 27, in which the volume fraction (α _f ) of the ferrite phase of the weld metal was out of the range of the present invention, the absorption energy at −20 ° C. was inferior.

On the other hand, the test specimens Nos. 21 to 23 of the present invention
Tensile strength, -20 ° C absorbed energy, general corrosion resistance,
The pitting corrosion resistance and stress corrosion cracking resistance were all excellent.

[0076] As described above, according to the method of the present invention, excellent have and toughness 13Cr steel equivalent 0.2% proof stress (or 400 MPa), and, in an environment containing chloride or C0 ₂ SUS316 Industrially useful effects can be brought about, which enable economical production of a duplex stainless steel welded pipe having corrosion resistance higher than steel, overall corrosion resistance, pitting corrosion resistance and stress corrosion cracking resistance.

[Brief description of the drawings]

Figure 1: Nickel content, pitting potential and 0 ° C, -50
It is a graph which shows the relationship with the absorption energy of ° C.

Fig. 2 Volume fraction of ferrite phase (α _f ), 0 ℃, -50
It is a graph which shows the relationship between the absorption energy of ° C, elongation, and the end surface crack of a steel plate.

FIG. 3 is a graph showing the relationship between the oxygen content of a weld metal, the absorbed energy of the weld metal at −20 ° C., and the corrosion rate of a weld in a pitting corrosion test.

FIG. 4 is a graph showing a relationship between a heating temperature during solution heat treatment, a corrosion rate in a Hughy test of a welded portion, an absorbed energy of a weld metal at −20 ° C., and a volume fraction (α _f ) of a ferrite phase.

FIG. 5 is a graph showing a relationship between a cooling rate in a temperature range of 500 to 850 ° C. and a corrosion rate of a weld metal in a Huey test, that is, intergranular corrosion resistance.

FIG. 6 is a schematic front view of a four-point bending jig used for measuring stress corrosion cracking resistance.

[Explanation of symbols]

 1 4-point bending jig, 2 hold-down jig, 3 push-up jig, A specimen.

──────────────────────────────────────────────────続 き Continuation of the front page (51) Int.Cl. ⁶ Identification symbol FI C22C 38/58 C22C 38/58 (56) References JP-A-6-240411 (JP, A) JP-A-2-298235 (JP) , A) (58) Fields investigated (Int. Cl. ⁶ , DB name) C21D 8/10 C21D 9/08 C22C 38/00 302

Claims (2)

(57) [Claims] 1. Carbon (C): 0.05 wt.% Or less, silicon (Si): 1.5 wt.% Or less, manganese (Mn): 2.0 wt.% Or less, nickel (Ni): 3.0 to 5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, and the balance: Fe and unavoidable impurities, the following equation (1): σ _0.2 = 20 × Cr (%) + 11 × Mn (%) − Ni (%) + 133 × N (%) − 38 The index (σ _0.2 ) indicating the 0.2% proof stress obtained by ───── (1) is 400
A steel ingot or a slab having the chemical composition described above is prepared, and the steel ingot or the slab is hot-rolled to prepare a steel sheet. An original plate is prepared by performing an annealing consisting of heating at a temperature within the range of T (° C.) or less and then cooling, T = 71 × Ni (%) + 6 × Mn (%) − 36 × Cr (%) − 42 × Si (%) + 1037 × C (%) + 1113 × N (%) + 1608───── (2) Next, the original plate was formed into a tube to prepare a base tube, and the base tube was prepared. And the weld metal of the seam is as follows: carbon (C): 0.05 wt.% Or less, silicon (Si): 1.5 wt.% Or less, manganese (Mn): 2.0 wt.% Or less, nickel ( Ni): 3.0 to 5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, Oxygen (O): 0.035 wt.% Or less, and balance: Fe and Unavoidable impurities, and 0.
A welded steel pipe having a chemical composition having an index of 2% proof stress (σ _0.2 ) of 400 or more is prepared.Then, the welded steel pipe is heated at a temperature in the range of 900 to 1150 ° C., and 500 to 850 ° C. A solution heat treatment comprising cooling at a rate of 1 ° C./sec or more in a temperature range within the range of 0 ° C., and thus, the volume fraction α _f of the ferrite phase of the base material of the welded steel pipe is in the range of 0.40 to 0.60. And the pitting corrosion resistance index (α _P1 ) of the ferrite phase determined by the following equation (3) and the pitting corrosion resistance index (γ _P1 ) of the austenite phase determined by the following equation (4). 23.5
Not less _{than, α P1 = 23 × Cr (} %) / (3 × α f + 20) ──────────────── (3) γ P1 = 20 × Cr (%) / (3 × α _f +20) −16 × N (%) / (α _f −1) ──── (4) Then, the volume fraction of the ferrite phase in the seam weld metal of the welded steel pipe (α _f ) is in the range of 0.30 to 0.50, and the pitting corrosion resistance index (α _P1 ) of the ferrite phase and the pitting corrosion resistance index (γ
_P1 ) producing a duplex stainless steel welded pipe having a strength larger than any one of the lower values of the α _P1 and the γ _P1 of the base material, characterized by having excellent strength, toughness and corrosion resistance. Method for manufacturing stainless steel welded steel pipe. 2. Carbon (C): 0.05 wt.% Or less, silicon (Si): 1.5 wt.% Or less, manganese (Mn): 2.0 wt.% Or less, nickel (Ni): 3.0 to 5.0 wt.%, Chromium (Cr): 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, and the balance: Fe and unavoidable impurities, the following equation (1): σ _0.2 = 20 × Cr (%) + 11 × Mn (%) - Ni ( %) + 133 × N (%) - 38────── (1) index (sigma _0.2) showing a 0.2% proof stress obtained by 400
A steel ingot or a steel slab having the chemical composition described above is prepared, and the steel ingot or the steel slab is heated at a temperature equal to or lower than T (° C.) determined by the following equation (2), and then 900 ° C. An original plate is prepared by performing hot rolling including finishing at the above temperature, and T = 71 × Ni (%) + 6 × Mn (%) − 36 × Cr (%) − 42 × Si (%) + 1037 × C (%) + 1113 × N (%) + 1608 ───── (2) Then, the raw plate was formed to prepare a raw tube, and a seam portion of the raw tube was welded. , Carbon (C): 0.05 wt.% Or less, Silicon (Si): 1.5 wt.% Or less, Manganese (Mn): 2.0 wt.% Or less, Nickel (Ni): 3.0-5.0 wt.%, Chromium (Cr) : 21.0 to 25.0 wt.%, Nitrogen (N): 0.25 wt.% Or less, Oxygen (O): 0.035 wt.% Or less, and the balance: Fe and unavoidable impurities.
A welded steel pipe having a chemical composition having an index (σ _0.2 ) indicating a 2% proof stress of 400 or more is prepared, and then the welded steel pipe is heated at a temperature in the range of 900 to 1150 ° C., and 500 to 850 ° C. A solution heat treatment comprising cooling at a rate of 1 ° C./sec or more in a temperature range within the range of 0 ° C., and thus, the volume fraction α _f of the ferrite phase of the base material of the welded steel pipe is in the range of 0.40 to 0.60. And the pitting corrosion resistance index (α _P1 ) of the ferrite phase determined by the following equation (3) and the pitting corrosion resistance index (γ _P1 ) of the austenite phase determined by the following equation (4). 23.5
Not less _{than, α P1 = 23 × Cr (} %) / (3 × α f + 20) ──────────────── (3) γ P1 = 20 × Cr (%) / (3 × α _f +20) −16 × N (%) / (α _f −1) ──── (4) Then, the volume fraction of the ferrite phase in the seam weld metal of the welded steel pipe (α _f ) is in the range of 0.30 to 0.50, and the pitting corrosion resistance index (α _P1 ) of the ferrite phase and the pitting corrosion resistance index (γ
_P1 ) producing a duplex stainless steel welded pipe having a strength larger than any one of the lower values of the α _P1 and the γ _P1 of the base material, characterized by having excellent strength, toughness and corrosion resistance. Method for manufacturing stainless steel welded steel pipe.