JP2004099930A - High-strength welded steel pipe having excellent toughness of weld zone, and method for manufacturing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-strength welded steel pipe which has low-temperature toughness, facilitates on-site welding, and has tensile strength above ≥800 MPa (over X100 stipulated in API (Americal Petroleum Institute)) and a method for manufacturing the same. <P>SOLUTION: The high-strength welded steel pipe having the excellent toughness of a weld zone is ≥800 MPa in the tensile strength of a base metal and weld metal, ≥230 Hv in the Vickers hardness in the coarse grain region of a weld heat affected zone, and is 0.5 to 1 in the ratio to Hv calculated by equation (1): Hv=270+1300C (1) from the amount of C [mass%] of the base metal. The manufacturing method for the steel pipe comprises forming a steel sheet to a tubular form, subjecting the butt zone thereof to submerged arc welding from the inside and outside surfaces and expanding the pipe, in which the weld zone is cooled at ≥1°C/s from 600°C up to 400°C. <P>COPYRIGHT: (C)2004,JPO

Description

【００６５】
【表２】

【００６６】
【表３】

【００６７】
【表４】

【００６８】
【表５】

【００６９】
【発明の効果】
本発明により低温靱性および現地溶接性の優れた高強度ラインパイプ（引張強さ８００ＭＰａ以上、ＡＰＩ規格Ｘ１００超）用溶接鋼管が安定して大量に製造できるようになった。これにより、パイプラインの輸送効率、施工能率の飛躍的な向上が可能となり、産業上の貢献が極めて高い。
【図面の簡単な説明】
【図１】溶接熱影響部粗粒域の模式図。
【図２】Ａｃ_１近傍に再熱された粗粒ＨＡＺ部のミクロ組織の模式図。
【図３】Ａｃ_３近傍に再熱された粗粒ＨＡＺ部のミクロ組織の模式図。
【図４】溶接部の−３０℃でのシャルピー吸収エネルギ−［Ｊ］と溶接後の６００℃から４００℃までの冷却速度［℃／ｓ］の関係を示す図。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a low-temperature toughness (hereinafter, weld toughness) and a field weldability of a weld metal and a weld heat affected zone (hereinafter, HAZ) having a tensile strength of 800 MPa or more and suitable for a crude oil / natural gas line pipe or the like. The present invention relates to a high-strength steel pipe excellent in heat resistance and a method for producing the same.
[0002]
[Prior art]
In recent years, pipelines that transport crude oil and natural gas over long distances have been required to increase the efficiency of 1) higher transport efficiency by increasing the pressure and 2) improve the efficiency of on-site construction by reducing the outer diameter and weight of line pipes. High strength steel pipes for line pipes are used. Up to now, line pipes up to X80 (yield strength of 551 MPa or more, tensile strength of 620 to 827 MPa) according to the American Petroleum Institute of Technology (API) standard have been put to practical use, but further higher strength is required.
[0003]
Since a high-strength line pipe of X100 (tensile strength of 800 MPa) or more withstands a pressure approximately twice as high as that of X65, it is possible to increase the pressure and reduce the outer diameter and the wall thickness. Therefore, by adopting a high-strength line pipe of X100 or more, the material cost, transportation cost, and on-site welding construction cost can be suppressed as compared with X65, so that the pipeline laying cost can be significantly reduced.
[0004]
Until now, a high-strength steel pipe of X100 or more has been developed, and its manufacturing method is disclosed in Patent Document 1. In this method, the steel pipe is welded after being formed, and then the entire steel pipe is subjected to aging treatment to increase the strength by precipitation strengthening of Cu, resulting in an increase in manufacturing cost. Patent Document 2 discloses a high-strength steel pipe in which HAZ softening is suppressed by optimizing the components of a weld metal.
[0005]
However, in high-strength steel pipes exceeding X100, it is very difficult to ensure the low-temperature toughness of the HAZ, and in particular, the toughness of the HAZ that has been subjected to multi-pass welding in two or more passes is significantly reduced. For example, Charpy at −30 ° C. Some have an absorption energy of less than 50 J. This is due to the coarsening of the particle size of the HAZ reheated to a high temperature by two or more passes of welding.
[0006]
As described above, as a problem associated with increasing the strength of steel pipes for line pipes, weld toughness and on-site weldability are particularly important, and there is a demand for an early development of an epoch-making high-strength steel pipe (over X100) that overcomes these. Have been.
[Patent Document 1]
Japanese Patent Application Laid-Open No. H8-311549 [Patent Document 2]
JP-A-10-306347
[Problems to be solved by the invention]
The present invention provides a high-strength welded steel pipe having an excellent balance between strength and low-temperature toughness, good weld toughness, and easy on-site weldability and a tensile strength of 800 MPa or more (API standard X100 or more) and a method for producing the same. Things.
[0008]
[Means for Solving the Problems]
In order to develop a high-strength welded steel pipe having a tensile strength of 800 MPa or more and excellent weld toughness and on-site weldability, the present inventors clarify the cause of the decrease in the low-temperature toughness of the weld and the HAZ. I did diligent research for that. First, it was found that the low-temperature toughness of the weld and the HAZ was in the HAZ coarse grain region shown in FIG. FIG. 1 schematically shows a welded portion obtained by performing two-pass welding. The HAZ coarse-grained region is a portion where the HAZ obtained by the first-pass welding is reheated to a high temperature by the second-pass welding. And the particle size is coarse.
[0009]
Furthermore, in these HAZ coarse-grained regions, as shown in FIG. 2 at the prior austenite grain boundaries, and within the grains of the prior austenite, as shown in FIG. 3, as shown in FIG. ) Was formed, and it was found that this was the starting point of brittle fracture. 2 and 3 schematically show the inside of one old austenite grain and the grain boundary. The inside of the prior austenite grains and the grain boundaries near the crystal grains in the center of the figure have the same fine structure, but are omitted except for the inside of the crystal grains and the grain boundaries in the center of the figure in order to clearly show the features.
[0010]
The present inventors have conducted a more detailed study, and at least after the final welding, have been cooled from 600 ° C. to 400 ° C. at 1 ° C./s, thereby suppressing the generation of MA in the weld heat affected zone coarse grain region, We succeeded in improving the toughness and invented a new high-strength welded steel pipe with excellent weld toughness and on-site weldability. That is, the gist of the present invention is as follows.
(1) The tensile strength of the base metal and the weld metal is 800 MPa or more, the Vickers hardness of the weld heat-affected zone coarse-grained area is 230 Hv or more, and the C content [mass%] of the base metal is calculated by the formula (1). A high strength welded steel pipe excellent in weld toughness, characterized in that the calculated ratio to Hv is 0.5 to 1.
Hv = 270 + 1300C (1)
(2) The base material is, by mass%, C: 0.02 to 0.10%, Si: 0.6% or less, Mn: 1.5 to 2.5%, P: 0.015% or less, S : 0.003% or less, Ni: 0.1 to 2.0%, Mo: 0.1 to 0.6%, Nb: 0.001 to 0.10%, Ti: 0.030% or less, Al: B: 0.0020% or less, N: 0.006% or less, V: 0.10% or less, Cu: less than 1.0%, Cr: 1.0% or less, One or more of Ca: 0.01% or less, REM: 0.02% or less, Mg: 0.006% or less, the balance consisting of iron and inevitable impurities, and the mass of the weld metal is %, C: 0.02 to 0.14%, Si: 0.05 to 0.4%, Mn: 1.2 to 2.2%, P: 0.010% or less, S: 0.010% Less than, i: 1.3 to 3.2% or less, B: 0.005% or less, and one or more of Cr, Mo, and V in the range of Cr + Mo + V: 1.0 to 2.5%. The high-strength welded steel pipe according to (1), wherein the high-strength welded steel pipe contains iron and unavoidable impurities.
(3) The Ni content of the weld metal is 1% by mass or more higher than the base metal, and the above formula (1) is obtained from the Vickers hardness of the weld metal and the weld heat affected zone and the C content [mass%] of the weld metal and the base material. (1) or (2), wherein the ratio of Hv calculated according to (1) or (2) is high.
(4) In a method for manufacturing a steel pipe in which a steel sheet is formed into a tubular shape and the butted portion is subjected to submerged arc welding from the inner and outer surfaces and expanded, at least after final welding, the welded portion is at least 1 ° C / s from 600 ° C to 400 ° C. The method for producing a high-strength welded steel pipe having excellent weld toughness according to any one of (1) to (3), characterized in that the steel pipe is cooled by a cooling method.
(5) In mass%, C: 0.01 to 0.12%, Si: 0.05 to 0.3%, Mn: 1.2 to 2.4%, Ni: 4.0 to 8.5%. And a welding wire and a sintering mold or a molten metal containing one or more of Cr, Mo, and V in the range of Cr + Mo + V: 3.0 to 5.0%, with the balance being iron and unavoidable impurities. (4) The method for producing a high-strength welded steel pipe excellent in toughness of a weld portion according to (4), wherein the welding is performed using a mold flux.
[0011]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the contents of the present invention will be described in detail.
[0012]
The present inventors conducted a tensile test of a base metal and a weld metal of a welded steel pipe in accordance with ASTM E8, and conducted a detailed study on the weld toughness of a steel pipe having a tensile strength of 800 MPa or more. First, a V-notch test piece of JIS Z 2202 was collected from the HAZ coarse grain area shown in FIG. 1 and subjected to a Charpy impact test at −30 ° C. in accordance with JIS Z 2242.
[0013]
Since the test piece having a Charpy absorbed energy of less than 50 J was partially brittle, the starting point was investigated. As a result, it became clear that the starting point of the brittle fracture can be classified into the following two types according to the strength of the steel pipe.
(1) MA shown in FIG. 2 which exists in the HAZ coarse-grained region where the HAZ heated to just below the melting point by the first pass welding is further reheated to the vicinity of Ac ₁ by the second pass welding. The temperature immediately below the melting point is a temperature of 1100 ° C. or more and less than the melting point, and the temperature near Ac ₁ is in a range of 650 to 750 ° C. In the HAZ coarse-grained region, depending on the heat input of welding, massive MA having a length of about several tens of μm exists along a coarse old austenite grain boundary composed of upper bainite of about 100 to 300 μm. This is often seen as a starting point of brittle fracture of HAZ in a high-strength steel pipe of X100 or less.
(2) MA shown in FIG. 3 in which the HAZ heated just below the melting point by the first pass welding is present in the HAZ coarse-grained region reheated to the vicinity of Ac ₃ by the second pass welding. The temperature in the vicinity of Ac ₃ is in the range of 850 to 1000 ° C. In this HAZ coarse grain region, although different depending on the heat input of welding, about 10 μm old austenite grains are mixed in the grain boundaries of coarse old austenite grains of about 100 to 300 μm, and in the coarse old austenite grains, Lumpy MA having a length of about several tens of μm exists. The inside of the former austenite grains is granular bainite. This is often seen as a starting point of brittle fracture of HAZ in a high-strength steel pipe of X120 or more.
[0014]
The reason why old austenite grains of about 10 μm are mixed in the grain boundaries of coarse old austenite grains of about 100 to 300 μm will be described. At the time of cooling after welding in the first pass, residual austenite is generated in the grains, which grow and coalesce during heating by welding in the second pass. Further, the growth of the austenite grains newly generated at the grain boundaries is suppressed by the fine Nb carbide. As a result, a structure in which austenite grains having a size of about 10 μm are formed at grain boundaries of coarse austenite grains having a size of about 100 to 300 μm is formed, and is cooled to become austenite grains. Moreover, because it is heated to Ac ₃ near the temperature higher than Ac ₁ near, B, etc. to improve the hardenability to form precipitates, the amount of solid solution decreases. As a result, the quenchability is reduced, the material is transformed into granular bainite upon cooling, and MA is generated in the grains.
[0015]
In addition, when it does not distinguish upper bainite, lower bainite, and granular bainite, it is called bainite. Although bainite contains retained austenite and martensite, it is difficult to distinguish between bainite and martensite and to observe retained austenite in an optical microscope structure.
[0016]
The MA formed in the grain boundaries and in the grains of the upper bainite and the granular bainite is coarse and harder than the upper bainite and the granular bainite, so that it is a starting point of the brittle fracture. Therefore, it was considered that toughness would be improved if the cooling rate after welding was increased to suppress the formation of MA, and if the microstructure was mainly composed of lower bainite and martensite and the remainder was retained austenite. Therefore, the influence of the cooling rate after welding on the weld toughness of a 0.07% C-1.9% Mn-based welded steel pipe in which the tensile strength of the base metal and the weld metal is 800 MPa or more was investigated in detail. The control of the cooling rate was performed in a range from 600 ° C. at which the bainite transformation started at the time of cooling to 400 ° C. at which the bainite transformation was completed by about 50%.
[0017]
FIG. 4 shows the relationship between the cooling rate from 600 ° C. to 400 ° C. after welding and the Charpy absorbed energy at −30 ° C. As a result, it was found that it is necessary to cool from 600 ° C. to 400 ° C. at 1 ° C./s or more in order to make the Charpy absorbed energy at −30 ° C. to be 50 J or more. Further, it was found that the Charpy absorbed energy should be 100 J or more, 150 J or more, and 200 J or more by setting the cooling rate to 5 C / s or more, 10 C / s or more, and 30 C / s or more, respectively.
[0018]
Next, the reasons for limiting the component elements will be described below.
[0019]
C content is limited to 0.02 to 0.10%. C is extremely effective in improving the strength of steel, and in order to obtain the target strength, an amount of C of 0.02% or more is necessary, and it is preferable to contain 0.04% or more. However, if the C content is too large than 0.10%, the low-temperature toughness and on-site weldability of the base material and HAZ are significantly deteriorated, so the upper limit is made 0.10% or less. Further, a preferable upper limit of C is 0.08% or less.
[0020]
Si is an element added for deoxidation and improvement of strength, but if added in excess of 0.6%, the HAZ toughness and on-site weldability are significantly deteriorated, so the upper limit was made 0.6% or less. Steel can be deoxidized with Al and Ti, and Si need not always be added, but is contained as an impurity in an amount of 0.01% or more.
[0021]
Mn is an element indispensable for ensuring the balance between excellent strength and low-temperature toughness with the microstructure of the steel of the present invention as lower bainite, and the lower limit thereof is 1.5% or more. However, if the Mn content is more than 2.5%, the hardenability of the steel increases to deteriorate the HAZ toughness and the on-site weldability, and also promotes the center segregation of the continuously cast steel slab and the low-temperature toughness of the base metal. Is also deteriorated, so the upper limit is made 2.5% or less.
[0022]
P and S are impurity elements, and the upper limits are set to 0.015% or less and 0.003% or less, respectively. The main reason for this is to further improve the low-temperature toughness of the base material and the HAZ. The reduction of the P content reduces the center segregation of the continuously cast slab, and also prevents the intergranular fracture and improves the low-temperature toughness. Further, the reduction of the amount of S has an effect of reducing MnS to be elongated by hot rolling and improving ductility. Note that the lower limits of P and S are 0.003% or more and 0.0001% or more in the current technology, respectively.
[0023]
The purpose of adding Ni is to improve the strength of the steel of the present invention having a small C content without deteriorating the low-temperature toughness and the on-site weldability. Compared with Mn, Cr or Mo addition, Ni addition not only causes less formation of a hardened structure harmful to low-temperature toughness in the rolled structure, but addition of a small amount of 0.1% or more Ni improves HAZ toughness. Is also effective. In order to improve the HAZ toughness, the amount of Ni added is preferably set to 0.3% or more. However, if the addition amount is more than 2.0%, not only economy but also HAZ toughness and on-site weldability deteriorate, so the upper limit is made 2.0% or less. Ni addition is also effective in preventing Cu cracking during continuous casting and hot rolling. In this case, it is necessary to add the Ni amount at least 1/3 of the Cu amount.
[0024]
The reason for adding Mo is to improve the hardenability of steel and make the microstructure lower bainite. In the B-added steel, the effect of improving the quenching of Mo is enhanced, and Mo coexists with Nb to suppress recrystallization of austenite during controlled rolling, and is also effective in refining the austenite structure. In order to obtain such an effect, it is necessary to add Mo by 0.1% or more. However, excessive Mo addition exceeding 0.6% deteriorates HAZ toughness and on-site weldability, and further impairs the effect of improving the hardenability of B, so the upper limit was made 0.6% or less.
[0025]
Nb coexists with Mo to suppress the recrystallization of austenite during controlled rolling, not only to refine the structure, but also to contribute to precipitation hardening and hardenability, and toughen the steel. In particular, when Nb and B coexist, the effect of improving hardenability increases synergistically. Since the Nb content is required to be 0.001% or more to obtain the effect, the lower limit is set to 0.001%. However, if the added amount of Nb is more than 0.10%, the HAZ toughness and on-site weldability are adversely affected, so the upper limit is set to 0.10% or less.
[0026]
The addition of Ti forms fine TiN, suppresses the coarsening of crystal grains and the austenite grains of HAZ during reheating of the slab, refines the microstructure, and improves the low-temperature toughness of the base material and HAZ. To obtain such an effect, the lower limit of the Ti content is preferably set to 0.001%. Further, it also has a role of fixing solid solution N harmful to the effect of improving the hardenability of B as TiN. For this purpose, it is preferable to add 3.4 N or more of Ti. When Al is less than 0.004%, Ti forms an oxide, acts as an intragranular ferrite generation nucleus in the HAZ, and has an effect of reducing the HAZ toughness. However, if the amount of Ti is too large than 0.03%, coarsening of TiN and a precipitation effect by TiC occur, deteriorating low-temperature toughness. Therefore, the upper limit is limited to 0.03%.
[0027]
Al is an element usually contained in steel as a deoxidizing material, and also has an effect on refining the structure. However, if the amount of Al exceeds 0.07%, Al-based nonmetallic inclusions increase and impair the cleanliness of the steel, so the upper limit was made 0.07% or less. Deoxidation can be performed with Ti or Si, and Al need not always be added. However, in the current technology, the content is 0.001% or more as an impurity.
[0028]
Next, the purpose of adding B, N, V, Cu, Cr, Ca, REM, and Mg will be described.
[0029]
The main purpose of further adding these elements to the basic components is to further improve the strength and low-temperature toughness and expand the size of a steel material that can be manufactured without impairing the excellent characteristics of the steel of the present invention. .
[0030]
B is a very effective element for dramatically increasing the hardenability of steel in a trace amount and for making the microstructure lower bainite. Further, B enhances the effect of improving the hardenability of Mo, and synergistically increases the hardenability together with Nb. To obtain this effect, it is preferable to add B in an amount of 0.0003% or more. On the other hand, if it is added in excess of 0.0020%, not only the low-temperature toughness is deteriorated, but also the effect of improving the hardenability of B may be lost, so the upper limit was made 0.0020% or less.
[0031]
N forms TiN and suppresses the coarsening of the austenite grains of the HAZ during reheating of the slab and improves the low-temperature toughness of the base material and the HAZ. However, if the N content is more than 0.006%, HAZ toughness is degraded due to slab surface flaws and solid solution N, and the hardenability of B is reduced. Therefore, the upper limit of N is made 0.006% or less. There is a need. Since the lower the N content, the better, the lower limit is not specified, but usually contains 0.0015% or more as an impurity.
[0032]
V has almost the same effect as Nb, and the combined addition of Nb and V further enhances the excellent characteristics of the steel of the present invention. In order to sufficiently exhibit this effect, it is preferable to add V at 0.03% or more. The upper limit is acceptable up to 0.10% or less from the viewpoint of HAZ toughness and on-site weldability, but the addition of 0.03 to 0.08% is particularly preferable.
[0033]
Although Cu increases the strength of the base material and the welded portion, it is preferable to add 0.01% or more of Cu in order to sufficiently exhibit this effect. On the other hand, if Cu is added in an amount of 1.0% or more, HAZ toughness and on-site weldability are significantly deteriorated. Therefore, the upper limit of the amount of Cu is set to less than 1.0%.
[0034]
Although Cr increases the strength of the base metal and the welded portion, it is preferable to add 0.01% or more of Cr in order to sufficiently exhibit this effect. On the other hand, if the Cr content is more than 1.0%, the HAZ toughness and the on-site weldability are significantly deteriorated. For this reason, the upper limit of the Cr content is set to 1.0% or less.
[0035]
Ca and REM control the sulfide (MnS) morphology and improve low temperature toughness. In order to sufficiently exhibit this effect, it is preferable to add 0.0001% or more of Ca and REM. If the Ca content exceeds 0.01% and the REM exceeds 0.02%, CaO-CaS or REM-CaS is generated in large amounts to form large clusters and large inclusions, which not only impairs the cleanliness of the steel, It also has an adverse effect on local weldability. Therefore, the upper limit of the amount of Ca added is limited to 0.01% or less, and the upper limit of the amount of REM added is limited to 0.02% or less. Note that the S and O contents were reduced to 0.003% or less and 0.002% or less, respectively, and ESSP = (Ca) [1-124 (O)] / 1.25S was set to 0.5 ≦ ESSP ≦ 10. It is particularly effective to set it to 0.
[0036]
Mg forms a finely dispersed oxide, suppresses coarsening of the HAZ, and improves low-temperature toughness. In order to sufficiently exhibit this effect, it is preferable to add Mg in an amount of 0.0001% or more. On the other hand, if the amount of Mg exceeds 0.006%, a coarse oxide is generated, and on the contrary, the toughness is deteriorated. Therefore, the upper limit is made 0.006% or less.
[0037]
Next, the reasons for limiting the weld metal will be described.
[0038]
C content is limited to 0.02 to 0.14%. C is extremely effective in improving the strength of steel, and is required to be 0.02% or more in order to obtain a target strength. However, if the C content is more than 0.14%, low-temperature welding cracks are liable to occur, and the highest hardness of the so-called T-cross HAZ portion where seam welding intersects the circumferentially welded portion of the steel pipe subjected to field welding. , The upper limit is set to 0.14% or less. Note that a preferable upper limit of C is 0.10% or less.
[0039]
Si is required to be 0.05% or more in order to prevent blowholes, but if it exceeds 0.4%, the low-temperature toughness, particularly the low-temperature toughness in the HAZ coarse-grain region, is remarkably deteriorated. Therefore, the range of Si is set to 0.05 to 0.4%.
[0040]
Mn is an indispensable element for securing a balance between excellent strength and low-temperature toughness, and its lower limit is 1.2%. However, if Mn is more than 2.2%, segregation is promoted and not only the low-temperature toughness is deteriorated, but also the production of a welding material becomes difficult, so the upper limit was made 2.2%.
[0041]
P and S are impurity elements, and in order to reduce the low-temperature toughness and low-temperature cracking susceptibility of the weld metal, both upper limits are set to 0.010% or less. Note that the lower limits of P and S are 0.003% or more and 0.0001% or more in the current technology, respectively. Ni is an element that enhances hardenability to secure strength and further improves low-temperature toughness. However, if it is less than 1.3%, the target strength and low-temperature toughness cannot be obtained, so the lower limit is 1.3% or more. . On the other hand, if the Ni content is more than 3.2%, there is a risk of hot cracking, so the upper limit was made 3.2%.
[0042]
B is a very small element that enhances hardenability and is effective in improving the low-temperature toughness of the weld metal. However, if the content is more than 0.005%, the low-temperature toughness of the weld metal decreases. Therefore, the upper limit of the B content is set to 0.005% or less. Note that B is preferably added at 0.0003% or more.
[0043]
Cr, Mo, and V are all elements that enhance hardenability, and one or more of them are added to obtain high strength. This effect is not sufficient when Cr + Mo + V is less than 1.0%, and when Cr + Mo + V is added in an amount larger than 2.5%, low-temperature cracking tends to occur. Therefore, the range of Cr + Mo + V is set to 1.0 to 2.5%.
[0044]
In addition, the weld metal may contain elements such as Ti, Al, Zr, Nb, and Mg that are added as necessary to improve the refining and solidification during welding. Next, the welding wire will be described.
[0045]
C is set to 0.01 to 0.12% in consideration of dilution with the base metal component and mixing of C from the atmosphere in order to obtain a range of the amount of C required for the weld metal.
[0046]
Si is set to 0.05 to 0.3% in consideration of dilution with a base metal component in order to obtain a range of Si amount required for the weld metal.
[0047]
Mn was set to 1.2 to 2.4% in consideration of dilution by a base metal component in order to obtain a range of Mn amount required for the weld metal.
[0048]
Ni was set to 4.0 to 8.5% in consideration of dilution with a base metal component in order to obtain a range of the amount of Ni required for the weld metal.
[0049]
One or two or more types of Cr + Mo + V are added. However, in order to obtain a range of the amount of Cr + Mo + V required for the weld metal, the content is set to 3.0 to 5.0% in consideration of dilution with a base metal component.
[0050]
In addition, it is desirable that the impurities of P and S are as small as possible, and B can be added for securing the strength. Further, Ti, Al, Zr, Nb, Mg and the like are used for the purpose of deoxidation.
[0051]
In addition, since the weld metal has a solidified structure and is inferior in toughness to the base metal, it is preferable to increase Ni, which improves the toughness, to be higher than the base material. In order to sufficiently exhibit this effect, it is preferable that the Ni content of the weld metal is 1% or more higher than that of the base metal.
[0052]
Next, the Vickers hardness will be described.
[0053]
In order to improve the low-temperature toughness of the HAZ coarse-grain region, the Vickers hardness of the HAZ coarse-grain region shown in FIG. 1 needs to be 230 Hv. If the Vickers hardness of the HAZ coarse-grained area is less than 230 Hv, the microstructure of the HAZ coarse-grained area is granular bainite or upper bainite, and -30 ° C because MA exists in the grains and / or at the grain boundaries. , The Charpy absorbed energy of the sample falls to less than 50 J. Further, in order to improve the low-temperature toughness of the HAZ and make the Charpy absorbed energy at −30 ° C. to be 100 J or more, it is preferable that the Vickers hardness of the HAZ coarse-grained region be 250 Hv or more. The upper limit of the Vickers hardness in the HAZ coarse-grain area is Hv calculated from the C content of the base material by Hv = 270 + 1300C.
[0054]
The measured value of the Vickers hardness of the HAZ is such that the ratio to the Hv calculated from the C content of the base material by Hv = 270 + 1300C is in the range of 0.5 to 1. Satisfying this, it is preferable that the Vickers hardness of the HAZ coarse-grained area is 230 Hv or more, and that no polygonal ferrite is generated by an optical micrograph, a scanning electron micrograph, or a hyperelectron micrograph. As a result, the HAZ coarse-grained region has a microstructure mainly composed of lower bainite and martensite, and the remainder is composed of retained austenite. Main lower bainite and martensite means that the lower bainite and martensite have a microstructure of 90 to 100% in area ratio, and the rest is retained austenite.
The Vickers hardness of the HAZ coarse grain area is measured according to JIS Z2244. In addition, a small piece of a portion containing the weld metal and the HAZ is cut out from the sample with the cross section perpendicular to the welding direction as an observation surface, mirror-polished, and nital-etched. The microstructure of this sample is observed with an optical microscope, and the Vickers hardness of the HAZ coarse grain region is measured. The sample may be etched by repeller etching. The measurement is preferably performed by cutting out a plurality of samples and calculating as an average value of about 3 to 5 points. The Vickers hardness is measured with a test force in the range of 0.09807 to 980.7 N. If the test force is small, the indentation is small and the accuracy is reduced. Therefore, it is preferable to set the range of 0.9807 to 98.07N.
[0055]
Regarding the Vickers hardness of the weld metal and the HAZ, the ratio of the measured value of the Vickers hardness of the weld metal and the HAZ to the Hv calculated from the C content of the weld metal and the base material by Hv = 270 + 1300C is 0.5 to 1%. Is preferably within the range. It is preferable that this is satisfied, and that no polygonal ferrite is formed by an optical micrograph, a scanning electron micrograph, or a hyperelectron micrograph. This corresponds to the fact that the weld metal and the HAZ have an area ratio of lower bainite and martensite of 90 to 100%, and the balance has a microstructure consisting of retained austenite. The Vickers hardness of the weld metal and the HAZ is preferably calculated as an average value of about 3 to 5 points by preparing a plurality of samples in the same manner as the Vickers hardness of the HAZ coarse grain region. The measurement of the C content of the weld metal and the HAZ is performed by taking a sample from the weld metal and the HAZ and conforming to JIS G1211.
[0056]
Next, the manufacturing conditions will be described.
[0057]
A steel plate is formed into a tube, and the butted portion is subjected to submerged arc welding from the inner and outer surfaces, and then expanded. The welding of the butt portion may be performed by MAG arc welding on the outer surface and then by submerged arc welding on the inner and outer surfaces.
[0058]
After welding, it is extremely important that the cooling rate from 600 ° C. at which bainite transformation starts at the time of cooling to 400 ° C. at which bainite transformation ends by about 50% be at least 1 ° C./s. Thereby, the HAZ coarse-grained region becomes a microstructure composed of lower bainite composed of lath-like ferrite and fine cementite and a microstructure composed of martensite and retained austenite, thereby improving toughness. On the other hand, when the cooling rate is less than 1 ° C./s, upper bainite composed of lath-like ferrite and MA at the lath boundary is generated, and when the cooling rate is further reduced, granular bainite in which the upper bainite has collapsed is generated and the toughness is reduced. . Since granular bainite is formed at 500 to 600 ° C, upper bainite is formed at 450 to 500 ° C, and lower bainite is formed at 400 to 450 ° C, it is preferable to cool the temperature range from 600 ° C to 450 ° C particularly quickly. Furthermore, in order to improve the toughness of the HAZ coarse-grained region and make the Charpy absorbed energy at −30 ° C. 100 J or more, the cooling rate is preferably 5 ° C./s or more. The higher the cooling rate, the more the toughness of the HAZ coarse grain region can be further improved. That is, in order to improve the Charpy absorbed energy at −30 ° C. to 150 J or more and 200 J or more, the cooling rate is preferably set to 10 ° C./s or more and 30 ° C./s or more, respectively. Although the upper limit of the cooling rate is not particularly defined, there is a limit due to technical restrictions and it depends on the sheet thickness, but at present, it is difficult to cool faster than 300 ° C./s.
[0059]
The forced cooling for increasing the cooling rate after welding may be forced air cooling by a fan, but air, nitrogen, helium, argon or other gas, water, mist or dry ice can be blown. The forced cooling needs to be performed at least after the final welding, but may be performed after each pass welding.
[0060]
【Example】
Steel having the components shown in Table 1 was melted in a converter, and cast into 240 mm thick slabs by continuous casting. These slabs were rolled into steel sheets of 14 to 25 mm under the conditions shown in Table 2. Further, after these steel sheets were subjected to UO forming, submerged arc welding was performed from the inner surface and the outer surface using wires and fluxes having the components shown in Table 3. After the outer surface welding, forced air cooling by a fan or forced cooling by blowing gas such as air was performed. The measurement of the cooling rate at this time was performed by a thermocouple attached to a portion 3 mm from the outer surface of the steel pipe of the weld metal after the inner surface welding. Thereafter, the pipe was expanded to a steel pipe having an outer diameter of 711 to 1219 mm.
[0061]
An analysis sample was taken from a part of the welded portion of these steel pipes, and the result of component analysis is shown in Table 4. Test pieces were taken from the base material and the weld metal of the steel pipe, and a tensile test was performed in accordance with ASTM E8, and it was confirmed that the tensile strength of the base material and the weld metal was 800 MPa or more. Further, a V-notch test piece of JIS Z 2202 was taken from the welded portion of the steel pipe and the HAZ, and a Charpy impact test was performed at −30 ° C. in accordance with JIS Z 2242 to evaluate the Charpy absorbed energy. In addition, the notch position was HAZ of 1 mm on the base material from the meeting part of the base metal and the weld metal in the center part of the plate thickness and the meeting part.
[0062]
Also, a plurality of small pieces were cut out from the HAZ, mirror-polished and nital-etched, and the Vickers hardness of the HAZ coarse-grained area was measured at 3 to 5 points at a test force of 980.7 N in accordance with JIS Z 2244, and the average value was obtained. Calculated. Table 5 shows the test results together with the cooling rates from 600 ° C to 400 ° C. Note that vE- ₃₀ in Table 5 is the Charpy absorbed energy at -30 ° C, FL is a notch position where the base metal and the weld metal meet, and FL + 1 mm is a 1 mm HAZ from the meeting portion to the base metal. Means that.
[0063]
Production No. in which the weld was heat-treated according to the present invention. In Nos. 1 to 20, the Charpy absorbed energy at -30 ° C exceeded 50 J, which was extremely good. On the other hand, the production No. In Nos. 21 to 32, since the cooling rate after welding is not within the range of the present invention, the Vickers hardness in the HAZ coarse-grained area is reduced, and the low-temperature toughness of the weld, particularly the HAZ indicated by FL + 1 mm, is significantly reduced. And the Charpy absorbed energy at −30 ° C. is less than 50 J.
[0064]
[Table 1]

[0065]
[Table 2]

[0066]
[Table 3]

[0067]
[Table 4]

[0068]
[Table 5]

[0069]
【The invention's effect】
INDUSTRIAL APPLICABILITY According to the present invention, a high-strength welded steel pipe for a high-strength line pipe (tensile strength 800 MPa or more, API standard X100 or more) excellent in low-temperature toughness and on-site weldability can be stably mass-produced. As a result, the transportation efficiency and construction efficiency of the pipeline can be dramatically improved, and the industrial contribution is extremely high.
[Brief description of the drawings]
FIG. 1 is a schematic diagram of a welding heat affected zone coarse grain region.
Schematic diagram of a microstructure of FIG. 2 Ac ₁ coarse-grained HAZ portion is reheated in the vicinity.
[3] Ac ₃ schematic diagram of the microstructure of the coarse-grained HAZ portion is reheated in the vicinity.
FIG. 4 is a graph showing the relationship between the Charpy absorbed energy at -30 ° C. [J] of a welded portion and the cooling rate [° C./s] from 600 ° C. to 400 ° C. after welding.

Claims (5)

The tensile strength of the base metal and the weld metal is 800 MPa or more, the Vickers hardness of the weld heat-affected zone coarse grain area is 230 Hv or more, and Hv calculated from the C content [mass%] of the base material by the formula (1). The high strength welded steel pipe excellent in weld toughness characterized by having a ratio of 0.5 to 1.
Hv = 270 + 1300C (1) The base material is mass%
C: 0.02 to 0.10%,
Si: 0.6% or less,
Mn: 1.5 to 2.5%,
P: 0.015% or less,
S: 0.003% or less,
Ni: 0.1 to 2.0%,
Mo: 0.1 to 0.6%,
Nb: 0.001 to 0.10%,
Ti: 0.030% or less,
Al: not more than 0.07%,
B: 0.0020% or less,
N: 0.006% or less,
V: 0.10% or less,
Cu: less than 1.0%,
Cr: 1.0% or less,
Ca: 0.01% or less,
REM: 0.02% or less,
Mg: one or more of 0.006% or less, the balance being iron and unavoidable impurities, and the weld metal in mass%
C: 0.02 to 0.14%,
Si: 0.05 to 0.4%,
Mn: 1.2 to 2.2%,
P: 0.010% or less,
S: 0.010% or less,
Ni: 1.3 to 3.2% or less,
B: contains 0.005% or less, and further contains one or more of Cr, Mo, and V as Cr + Mo + V: 1.0 to 2.5%
The high-strength welded steel pipe having excellent weld toughness according to claim 1, characterized in that the content is within the range described above, and the balance consists of iron and inevitable impurities. The Ni content of the weld metal was 1% by mass or more higher than that of the base metal, and was calculated from the Vickers hardness of the weld metal and the weld heat-affected zone and the C content [mass%] of the weld metal and the base material according to the above equation (1). The high strength welded steel pipe having excellent weld toughness according to claim 1 or 2, wherein the ratio of Hv is 0.5 to 1. In a method of manufacturing a steel pipe in which a steel plate is formed into a tubular shape and the butted portion is subjected to submerged arc welding from the inner and outer surfaces and expanded, at least after final welding, the welded portion is cooled at a rate of 1 ° C / s or more from 600 ° C to 400 ° C. The method for producing a high-strength welded steel pipe having excellent weld toughness according to any one of claims 1 to 3, characterized in that: In mass%,
C: 0.01 to 0.12%,
Si: 0.05-0.3%,
Mn: 1.2 to 2.4%,
Ni: 4.0 to 8.5%
And one or more of Cr, Mo and V are Cr + Mo + V: 3.0 to 5.0%
The high-strength welding according to claim 4, characterized in that the welding is performed using a welding wire composed of iron and unavoidable impurities and a sintering type or a molten type flux. Manufacturing method of steel pipe.

Images