JP2008248315A - Method for manufacturing ultrahigh-strength, high-deformability welded steel pipe having excellent toughness in base material and weld zone - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for manufacturing a welded steel pipe for line pipe having ≥900 MPa tensile strength by which cutting crack resistance can be improved without deteriorating toughness, particularly crack arrest property, while keeping high deformability and further a strength of a joint equal to or higher than that of a base material can be attained without decreasing toughness of a weld metal. <P>SOLUTION: A steel sheet having ≥900 MPa tensile strength and ≤85% yield ratio, which has specific components and also has a microstructure characterized as follows, is used: any of (ferrite plus bainite), (ferrite plus martensite) and (ferrite plus bainite plus martensite) includes ≥90% by area fraction; an area ratio of ferrite is 10 to 50%; and an average grain size of cementite in bainite and/or martensite is ≤0.5 μm. After the steel sheet is formed into a pipe shape by cold working, welding is carried out in such a way that a chemical composition of a weld metal consists of specific components by a hybrid welding process combining a laser using CO<SB>2</SB>gas shield with a gas shield arc welding using Ar-CO<SB>2</SB>gas shield. <P>COPYRIGHT: (C)2009,JPO&INPIT

Description

  The present invention relates to a manufacturing method of an ultra-high strength and high deformability line pipe having a tensile strength of 900 MPa or more and a yield ratio ≦ 85%, and does not require an additional step in refining the raw steel plate, and stops brittle crack propagation in the base metal part. Excellent properties, sufficient toughness in weld metal and weld heat affected zone (HAZ) of longitudinal seam welds, and satisfying the tensile strength that the joint tensile strength exceeds the base metal strength, and transport of natural gas and crude oil The present invention relates to a product that is suitable for use and has excellent safety.

  In recent years, line pipes used for transportation of natural gas and crude oil have been strengthened year by year in order to improve transportation efficiency by increasing pressure and to improve local welding work efficiency by reducing wall thickness. It is desirable to have a high deformability because cracks due to local buckling will not occur even if large deformations occur in the line pipe due to ground deformation at.

  So far, X100 grade line pipes have been put into practical use under the API standard, and further, the demand for X120 grade exceeding tensile strength 900 MPa is being realized.

  For example, in Patent Document 1, two-stage cooling is performed after hot rolling, and cooling of the second stage is stopped. A technique for achieving high strength by setting the temperature to 300 ° C. or lower is disclosed.

  Patent Document 2 discloses a technique relating to accelerated cooling and tempering conditions for obtaining high strength and high toughness while reducing the amount of expensive alloy element addition. Patent Document 3 discloses a steel pipe excellent in compression buckling resistance that exhibits a low yield ratio by giving an appropriate area ratio of the second phase structure according to the ratio of the tube thickness and the outer diameter. Yes.

Patent Document 4 discloses a technique related to component design for reducing the amount of alloying elements added to the base material and obtaining high strength and high toughness in the weld metal of the longitudinal seam weld.
JP 2003-293089 A JP 2002-173710 A Japanese Patent Laid-Open No. 09-184015 JP 2000-355729 A

  However, if the strength is increased by means such as accelerated cooling while keeping the amount of alloying elements in the base metal low, there will be a deviation from the strength of the weld heat affected zone (hereinafter abbreviated as HAZ).

  This is because the HAZ cooling rate remains constant unless the vertical seam welding method is changed, and the strength of the HAZ remains low. In distribution.

  As a result, even if the strength of only the base metal part is increased, if an actual pipe test such as a water pressure test is performed, the HAZ part with low strength will break, and in practical use, HAZ is considered from the viewpoint of safety. Partial softening is a problem to be solved.

  On the other hand, by reducing the cooling rate of accelerated cooling and increasing the amount of alloy elements in the base material, the strength difference between the base material and the HAZ part can be reduced, but in the case of circumferential welding where the pipes are joined together, the heat input is small. Since it is common to perform multi-layer welding, the HAZ hardness of the weld is rather increased, and the risk of cold cracking in the first layer weld is increased.

  Although increasing the strength of the weld metal in the longitudinal seam welds compensates for the lack of joint strength caused by this HAZ softened part, it cannot be said to be sufficient. Furthermore, increasing the strength of the weld metal without changing the welding method increases the sensitivity of weld metal defects such as cold cracking by increasing the amount of alloying elements in the weld metal and requires reworking. There is a concern that welding workability will be significantly deteriorated.

  Furthermore, when high strength is achieved by lowering the cooling stop temperature and forming a low-temperature transformation structure, the diffusible hydrogen remaining in the steel is the cause when sheared steel sheets are cut into the required size by shearing. Thus, a crack parallel to the plate surface (hereinafter referred to as a cutting crack) occurs.

  On the other hand, when heat treatment is performed after accelerated cooling, hydrogen in the steel is sufficiently diffused, so that cracking can be suppressed, but cementite precipitates and coarsens in the microstructure during the heat treatment process, resulting in reduced toughness, especially brittle crack propagation. DWTT (Drop Weight Tear Test) characteristics for evaluating stop characteristics deteriorate.

  In addition, the technology described in Patent Document 3 has a high deformability so that cracks do not occur even if a large deformation occurs in the line pipe due to a large earthquake or ground deformation in a frozen land zone. It is aimed at lowering the yield ratio (YR) divided by.

  In this technology, the steel pipe base material has the second phase, so the Charpy absorbed energy is low, and there is a concern that it will be inferior to the crack propagation stoppage property of ductile fracture caused by an extrinsic accident. Therefore, the tensile strength does not reach 900 MPa or more.

  Therefore, the present invention has a high deformation performance, improves toughness, in particular, crack propagation stopping characteristics without deteriorating the crack propagation stop property, and further improves the crack resistance of the weld metal and lowers the toughness of the weld metal. It aims at providing the manufacturing method of the welded steel pipe for line pipes of the tensile strength of 900 MPa or more which achieved joint strength.

As a result of intensive studies to solve the above problems, the present inventors have obtained the following knowledge.
1) In order to prevent buckling even if large deformation of the pipe occurs due to ground deformation, it is necessary to make the steel yield ratio 85% or less by making the steel a double phase structure, and high strength. In order to obtain it, the microstructure must be any of ferrite + bainite, ferrite + martensite, and ferrite + bainite + martensite, and the area ratio of ferrite must not exceed 50%.
2) Although the above-mentioned reverse phase structure steel can achieve high strength and low yield ratio, the same strength level of bainite or martensite single phase structure steel is used for Charpy absorbed energy, which is an index for evaluating the crack propagation stopping performance of ductile fracture. Although it tends to be lower, the desired Charpy absorption is achieved by appropriately controlling O, Ca, and S in the steel to control the form of sulfide inclusions in the steel, especially by reducing coarse MnS. Energy can be achieved.
3) When the steel with accelerated cooling after hot rolling to obtain bainite or martensite structure is cut by shearing in the above multiphase steel, it is cut due to diffusible hydrogen in the steel. Although cracks may occur on the surface, it can be prevented by setting the amount of hydrogen in the steel sheet before shearing to less than 2 ppm. For this purpose, dehydrogenation heat treatment at a temperature of at least 300 ° C. is required after accelerated cooling.

Specifically, after accelerating cooling is stopped, reheating is started immediately and the steel sheet temperature is raised to 300 ° C or higher to promote hydrogen diffusion. The amount falls below 2 ppm.
4) However, if cementite present in bainite or martensite becomes coarse due to reheating after accelerated cooling, the DWTT characteristic, which is an indicator of brittle crack propagation stoppage characteristics, deteriorates, but the heating rate for reheating increases. Even after the reheating is completed, the average particle size of cementite in the steel is prevented to 0.5 μm or less.
5) When a welded steel pipe is manufactured by welding the end after forming the steel plate with the above high strength, high toughness, and low yield ratio into a cylindrical shape, the strength decreases due to the effect of welding heat, and the joint strength decreases. In order to prevent this, a new welding method that combines laser and arc is effective as a welding method for increasing the cooling rate of the weld zone, and the welding efficiency is also excellent.

The present invention has been further studied based on the above knowledge, that is, the present invention,
1. % By mass
C: 0.03-0.12%
Si: 0.01 to 0.5%
Mn: 1.5 to 3%
Al: 0.01 to 0.08%
Nb: 0.01 to 0.08%
Ti: 0.005-0.025%
N: 0.001 to 0.01%
B: ≦ 0.0003%
Ca: 0.0005 to 0.01%
O: ≦ 0.003%
S: ≦ 0.001%
Furthermore, Cu: 0.01 to 0.5%
Ni: 0.01 to 1%
Cr: 0.01 to 0.5%
Mo: 0.01 to 0.5%
V: 0.01 to 0.1%
And the content of Ca, O, S satisfies the formula (1), the PcmB value calculated by the formula (2) satisfies PcmB ≦ 0.22, and the balance Fe and The composition of inevitable impurities,
Either ferrite + bainite, ferrite + martensite, or ferrite + bainite + martensite has an area fraction of 90% or more, and the area ratio of ferrite is 10 to 50%, and bainite and / or martensite A steel sheet having a microstructure with an average particle size of cementite of 0.5 μm or less and a tensile strength of 900 MPa or more and a yield ratio ≦ 85% is formed into a tubular shape by cold working, and then a laser using a CO ₂ gas shield and Ar— By the hybrid welding method combined with gas shielded arc welding using a CO ₂ gas shield, the chemical composition of the weld metal is
C: 0.05-0.09%
Si: 0.1 to 0.4%
Mn: 1.0-2.0%
Al: ≦ 0.015%
Cu: ≦ 0.5%
Ni: ≦ 3.0%
Cr: ≦ 1.0%
Mo: ≦ 1.0%
V: ≦ 0.1%
Ti: 0.003-0.10%
B: ≦ 0.0030%
O: ≦ 0.03%
N: ≦ 0.008%
Containing
The PcmW value calculated by Equation (3) satisfies PcmW ≦ 0.2,
In addition, a method for producing an ultra-high-strength, high-deformability welded steel pipe excellent in the toughness of the base metal and the welded portion, wherein the butt portion is welded so as to become the remaining Fe and inevitable impurities.
1 ≦ (1-130 × [O]) × [Ca] / (1.25 × [S]) ≦ 3 (1)
PcmB = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 * B (2)
PcmW = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 * B-12 * N-4 * O (3)
In the manufacturing method of the super-high-strength, high-deformability welded steel pipe excellent in the base metal and weld zone toughness described in 2.1, butt welding is performed, and hybrid welding is performed by combining laser and gas shield arc welding for each inner and outer surfaces A manufacturing method of ultra-high strength and high-deformability welded steel pipe with excellent base metal and weld toughness.
In the manufacturing method of super high strength and high deformability welded steel pipe excellent in base material and weld toughness described in 3.1, after butt welding and hybrid welding combining laser and gas shielded arc welding on the inner surface side, A method of manufacturing an ultra-high-strength, high-deformability welded steel pipe excellent in base metal and weld toughness, characterized by submerged arc welding of the outer surface.

  According to the present invention, the joint strength of the longitudinal seam portion is equal to or higher than the tensile strength of the base material, and even if a water pressure test or the like is performed, the seam welded portion does not break. It is possible to manufacture high strength welded steel pipes and is extremely useful industrially.

  According to the present invention, even when the welding method that has been used for the production of low-strength grades is used, the joint strength of the longitudinal seam is sufficiently high, the tensile strength is 900 MPa or more, and the yield ratio is 85%. The following high-strength, high-deformability welded steel pipes can be manufactured.

Hereinafter, the present invention will be described in detail by component composition, structure and manufacturing method.
[Ingredient composition of base material]
First, the component composition of the base steel sheet according to the present invention will be described. “%” Means “% by mass”.

C: 0.03-0.12%
C contributes to strength increase by supersaturated solid solution in the low-temperature transformation structure. To obtain this effect, 0.03% or more is necessary, but if it exceeds 0.12%, , The hardness of the circumferential welded part of the pipe is significantly increased, and cold cracking is likely to occur. Therefore, the C content is 0.03 to 0.12%.

Si: 0.01 to 0.5%
Si is an element that acts as a deoxidizer and increases the strength of steel by solid solution strengthening. However, if its amount is less than 0.01%, its effect cannot be obtained. It drops significantly. For this reason, Si content shall be 0.01 to 0.5%.

Mn: 1.5 to 3%
Mn acts as a hardenability improving element. In particular, 1.5% or more is necessary to obtain a low temperature transformation structure to achieve high strength in HAZ. However, in the continuous casting process, the concentration of the central segregation part is significantly increased. Cause delayed destruction. For this reason, Mn content shall be 1.5 to 3%.

Al: 0.01 to 0.08%
Al acts as a deoxidizing element. When the content is 0.01% or more, a sufficient deoxidation effect can be obtained. However, if the content exceeds 0.08%, the cleanliness in the steel decreases and the toughness deteriorates. For this reason, Al content shall be 0.01-0.08%.

Nb: 0.01 to 0.08%
Nb has the effect of expanding the austenite non-recrystallized region during hot rolling. In particular, Nb is contained in an amount of 0.01% or more in order to make the non-recrystallized region up to 950 ° C. However, if the amount exceeds 0.08%, the toughness of HAZ is significantly impaired. For this reason, Nb content shall be 0.01-0.08%.

Ti: 0.005-0.025%
Ti forms nitrides and is effective in reducing the amount of solute N in the steel. Precipitated TiN prevents the austenite grains from becoming coarse by the pinning effect, contributing to improved toughness of the base metal and HAZ. To do.

  In order to obtain the necessary pinning effect, it is necessary to contain 0.005% or more, but if added over 0.025%, carbides are formed, and the toughness is significantly deteriorated by precipitation hardening. , Ti content is 0.005 to 0.025%.

N: 0.001 to 0.006%
N is usually present as an inevitable impurity in steel, but TiN is formed by adding Ti as described above to suppress austenite coarsening.

  In order to obtain the required pinning effect, the content needs to be 0.001% or more, but if it exceeds 0.006%, it is heated to 1450 ° C or more near the weld, especially in the vicinity of the melting line. In addition, since TiN decomposes in HAZ and the adverse effect of solute N becomes significant, the N content is made 0.001 to 0.01%.

B: ≦ 0.0003%
When B is contained in excess of 0.0003%, it segregates at the austenite grain boundaries during hot rolling, and has the effect of suppressing the formation of ferrite transformation, and hinders the control of the multiphase structure of steel to achieve high deformability. Therefore, the B content is 0.0003% or less.

One or more of Cu, Ni, Vr, Mo, and V Cu, Cr, Mo, and V all act as a hardenability improving element. Therefore, one or two of these elements is used for the purpose of increasing the strength. The above is contained in the range shown below.

Cu: 0.01 to 0.5%
Cu contributes to improving the hardenability of steel at 0.01% or more. However, if the content exceeds 0.5%, the structure of HAZ formed at a high cooling rate by laser-arc hybrid welding, which will be described later, becomes martensite and causes deterioration of HAZ toughness. Has a Cu content of 0.01 to 0.5%.

Ni: 0.01 to 1%
Containing 0.01% or more of Ni contributes to improving the hardenability of steel. Especially when added in a large amount, it does not cause toughness deterioration, so it is effective for toughening. However, it is an expensive element and the effect is saturated even if it exceeds 1%. The content of is 0.01 to 1%.

Cr: 0.01 to 0.5%
When Cr is contained in an amount of 0.01% or more, it contributes to improving the hardenability of the steel. However, if the content exceeds 0.5%, the HAZ structure formed at a high cooling rate by laser-arc hybrid welding can be obtained. In order to become martensite and cause deterioration of HAZ toughness, when adding Cr, the content of Cr is set to 0.01 to 0.05%.

V: 0.01 to 0.1%
V strengthens precipitation by forming carbonitrides, and contributes particularly to the prevention of softening of the heat affected zone. This effect is obtained at 0.01% or more. However, if it exceeds 0.1%, precipitation strengthening significantly reduces toughness. Therefore, when V is added, its content is set to 0.01 to 0.00%. 1%.

Ca: 0.005 to 0.01%
Ca is an effective element for controlling the morphology of sulfides in steel, and its addition suppresses the formation of MnS, which is harmful to toughness. However, if added over 0.01%, CaO-CaS clusters are formed and the toughness deteriorates, so the Ca content is made 0.005 to 0.01%.

O: 0.003% or less, S: 0.001% or less In the present invention, O and S are inevitable impurities and define the upper limit of the content. The content of O is 0.003% or less from the viewpoint of suppressing the formation of inclusions that are coarse and adversely affect toughness.
Moreover, although the production | generation of MnS is suppressed by adding Ca, when there is much content of S, since MnS cannot be suppressed even by the form control by Ca, it is 0.001% or less.
(1-130 × [O]) × [Ca] / (1.25 × [S])
This parameter formula: (1-130 × [O]) × [Ca] / (1.25 × [S]) is for 1 ≦ (1-130 × [O]) × [ It is defined so as to satisfy Ca] / (1.25 × [S]) ≦ 3.

  By satisfying the above range, the formation of MnS is suppressed and the coarsening of CaO · CaS generated by adding excessive Ca is also suppressed, thereby improving the cleanliness of steel and obtaining high Charpy absorbed energy. Is possible.

  That is, Ca has the ability to form sulfides, and by adding, it suppresses the generation of MnS that lowers Charpy absorbed energy in molten steel during steelmaking, and forms CaS that is relatively harmless to toughness instead.

However, since Ca is also an oxide-forming element, it must first be added in anticipation of consumption as an oxide. From the viewpoint of suppressing inclusion formation that is coarse and adversely affects toughness, the effective Ca amount (Ca *) excluding CaO generation was determined after setting O ≦ 0.003% and S ≦ 0.005%. Calculate using the following equation (4) by regression,
Ca * = (1-130 × [O]) × [Ca] (4)
Furthermore, when Ca is added so that the value obtained by dividing the effective Ca amount (Ca *) by the stoichiometric ratio of Ca and S satisfies the following formula (5), all S in the steel forms CaS. It is.
[S] ≦ Ca * / 1.25 (5)
On the other hand, it was also found that when a large amount of Ca was added, the CaO · CaS produced was coarsened, and instead the Charpy absorbed energy was reduced. From the results of laboratory studies, 3 · [S] ≧ Ca * / 1.25 (6) in order to suppress this Ca coarsening.
It is desirable that Therefore, the range of this parameter formula: (1-130 × [O]) × [Ca] / (1.25 × [S]) should be 1 or more and 3 or less as the range between (5) and (6). Stipulate.

PcmB
PcmB is calculated by the following equation (7) as a weld cracking susceptibility composition, and correlates with a preheating temperature for preventing cold cracking in the HAZ part.

Fig. 2 shows the PCMB values for the low temperature cracking prevention preheating conditions of the HAZ part obtained by the low temperature cracking test conducted after applying various preheating temperatures for steels having various chemical compositions.
In the first layer welding at the time of circumferential welding between pipes, in order to prevent HAZ cracking when the pipe preheating temperature is allowed to 75 ° C., the PcmB value needs to be 0.22 or less. 22

In consideration of workability at the pipeline laying site, it is desirable that the pipe preheating temperature is low. From this viewpoint, the preferred range of PcmB is 0.20 or less.
PcmB = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 * B (7)
[Micro structure]
The microstructure of the base steel sheet is a double phase structure with an area fraction of 90% or more due to soft ferrite and hard phase. Furthermore, the area fraction of ferrite, which is a soft phase, and bainite and / or martensite constituting the hard phase Specifies the average cementite particle size.

  In the microstructure, the tensile strength is high because the rephase structure of ferrite + bainite, ferrite + martensite, ferrite + bainite + martensite accounts for 90% or more in area ratio, preferably 95% or more. Yield strength is low, and it is possible to achieve both high strength and low yield ratio.

  In order to obtain a tensile strength of 900 MPa or more, the hard phase is bainite, martensite, or a mixed structure thereof.

  From the viewpoint of toughness, it is desirable that the bainite and / or martensite constituting the hard phase has a structure transformed from fine austenite having a thickness in the thickness direction of 30 μm.

  The area fraction of ferrite, which is a soft layer, is 10 to 50%. If the ferrite content is less than 10%, the behavior is the same as that of a bainite or martensite single phase structure or a mixed structure of both, and the yield strength is high and it is difficult to achieve a desired low yield ratio.

  On the other hand, when ferrite exceeds 50%, soft ferrite is the main component, the tensile strength is greatly reduced, and it is difficult to achieve high strength exceeding 900 MPa. Preferably it is 10 to 30%.

  From the viewpoint of improving toughness, it is preferable that the ferrite has an average particle diameter of 20 μm or less. In the base material (steel material) according to the present invention, the presence of residual γ, island martensite, pearlite, etc. of less than 10% is allowed in the area fraction.

  Furthermore, the cementite average particle size in the bainite and / or martensite constituting the hard phase: ≦ 0.5 μm or less.

  In order to form bainite and / or martensite as a hard phase in the microstructure of steel, when accelerated cooling is performed as described later, cut cracks may occur during subsequent cutting of the steel sheet.

  To prevent this, when tempering is performed immediately after accelerated cooling, cementite precipitates in the bainite and / or martensite structure.

  When these cementites grow to a coarse size exceeding 0.5 μm by tempering, the DWTT characteristics deteriorate and the Charpy absorption energy decreases, so the average particle size of cementite in innite and / or martensite is 0.5 μm. The following.

  In particular, when the average particle size of cementite is less than 0.2 μm, the Charpy absorbed energy can be further increased.

The average particle size of cementite is measured using the following method. A sample for microstructural observation is taken in parallel to the cross section in the rolling direction of the steel sheet, mirror-polished, speed-etched, and then observed with a scanning electron microscope, and an electron micrograph is taken with 10 random fields of view. From the obtained photograph, the equivalent circle diameter of each cementite particle is calculated by image analysis, and the average value is calculated.
[Production conditions]
In the present invention, the base material only needs to have the above-described composition, microstructure, and strength characteristics. A manufacturing method is not particularly defined, but a preferable manufacturing method is as follows.
(1) Hot rolling heating temperature: 1000 to 1200 ° C
When performing hot rolling, the rolling slab is heated to 1000 ° C. or higher in order to completely austenite. On the other hand, when the steel slab is heated to a temperature exceeding 1200 ° C., even if TiN pinning is performed, the austenite grain growth is remarkable and the base material toughness deteriorates, so the heating temperature is set to 1000 to 1200 ° C.

Cumulative reduction at 950 ° C or lower ≧ 67%
As described above, 950 ° C. or lower is an austenite non-recrystallized region due to Nb addition. By accumulating large pressures in the temperature range, the austenite grains expand and become finer, especially in the thickness direction, and the toughness of the bainite steel obtained by accelerated cooling in this state is improved.

  However, if the cumulative rolling amount is less than 67%, the effect of refining is insufficient and the effect of improving the toughness of the steel is difficult to obtain, so the cumulative rolling amount is set to 67% or more. Note that the preferred range for significantly improving toughness is 75% or more.

Rolling end temperature: Ar3 point or higher, Ar3 point + 100 ° C. or lower When the rolling end temperature is lower than Ar3 point, rolling is performed in the ferrite transformation temperature range, the transformation-generated ferrite is greatly processed, and Charpy absorbed energy is reduced.

On the other hand, if the rolling is finished at a high temperature exceeding the Ar3 point + 100 ° C, the effect of refining by the austenite non-recrystallization zone rolling becomes insufficient, so the rolling finishing temperature is set to the Ar3 point or higher and Ar3 point + 100 ° C or lower. . The Ar3 point can be calculated by the following equation (8) using the value of the chemical composition of the base steel plate.
Ar3 = 910-310 × [C] −80 × [Mn] −20 × [Cu] −15 × [Cr] −55 × [Ni] −80 × [Mo] (8)
Accelerated cooling start temperature: Ar3 point -50 ℃ or more and less than Ar3 point to achieve a low yield ratio, a soft ferrite structure is generated. However, accelerated cooling suppresses ferrite transformation, After hot rolling, ferrite is transformed during the air cooling process until accelerated cooling starts. For this reason, the start temperature of accelerated cooling is made less than the Ar3 point.

  On the other hand, if the cooling start temperature is less than Ar3 point-50 ° C, the ferrite area fraction exceeds 50% and a tensile strength of 900 MPa or more cannot be obtained, so the lower limit is Ar3 point-50 ° C.

Accelerated cooling rate: 20-80 ° C / s
In order to obtain a hard phase composed of bainite and / or martensite, accelerated cooling is performed at 20 ° C./s or more. On the other hand, even if the cooling rate exceeds 80 ° C./s, the obtained structure does not change and the material improvement is saturated, so the upper limit is set to 80 ° C./s.

Cooling stop temperature for accelerated cooling: ≤250 ° C
In order to increase the strength of the steel sheet, the stop temperature of accelerated cooling is lowered to produce bainite and martensite that transform at low temperatures. When the cooling stop temperature exceeds 250 ° C., the accelerated cooling is stopped with insufficient transformation, and the rough structure generated from the untransformed lowers the toughness. Therefore, the cooling stop temperature is set to 250 ° C. or less.

  If a steel sheet that has been transformed at a low temperature by accelerated cooling is air-cooled as it is, diffusible hydrogen in the steel remains, which increases the possibility of cutting cracks.

  Therefore, reheating is performed immediately after cooling is stopped. If the time until reheating is long, it becomes difficult for hydrogen to diffuse due to the temperature drop during the air cooling process, so it is desirable to start heating within 300 seconds. (Preferably within 100 seconds.) The reheating method may be furnace heating or induction heating.

Heating rate during reheating: ≧ 5 ° C / s
When the heating rate at the time of reheating is less than 5 ° C./s, cementite is generated and coarsened during heating to a temperature exceeding 300 ° C. in particular, resulting in deterioration of DWTT characteristics. On the other hand, since it is possible to suppress the cementite coarsening by increasing the rate of temperature rise, the rate of temperature rise during reheating is set to 5 ° C./s or more.

Reheating temperature: 300 ° C to 450 ° C
When the reheating temperature is less than 300 ° C., hydrogen does not diffuse sufficiently and cutting cracks cannot be prevented. Therefore, the reheating temperature is set to 300 ° C. or higher. On the other hand, when heated to a temperature exceeding 450 ° C., the upper limit is set to 450 ° C. because the strength is significantly reduced due to softening by tempering.

  The steel making method is not particularly limited, but from the economical viewpoint, it is desirable to carry out the steel making process by the converter method and the slab casting by the continuous casting process.

There are no particular limitations on the method of forming the steel sheet produced by the above method into a steel pipe, and any of the conventionally used UOE forming, press bend forming, and roll forming can be used. Next, the reason for limiting the chemical composition in the weld metal will be explained.
[Chemical composition of weld metal]
C: 0.05-0.09%
In weld metal, C is an important element as a strengthening element for steel. In particular, in order to achieve overmatching of the joint part, the weld metal part also needs to have a tensile strength ≧ 900 MPa, and in order to obtain this strength, it is necessary to contain 0.05% or more.

  On the other hand, if it exceeds 0.09%, hot cracking of the weld metal tends to occur, so the C content of the weld metal is set to 0.05 to 0.09%.

Si: 0.1 to 0.4%
Si is necessary for deoxidizing the weld metal and ensuring good workability. If it is less than 0.1%, a sufficient deoxidation effect cannot be obtained. On the other hand, if it exceeds 0.4%, welding workability deteriorates. Therefore, the Si content of the weld metal is set to 0.1 to 0.4%.

Mn: 1.0-2.0%
Mn is an important element for increasing the strength of the weld metal. In particular, high strength such as tensile strength ≧ 900 MPa cannot be achieved by conventional acicular ferrite organization, and can be achieved by containing a large amount of Mn to form a bainite structure. In order to acquire this effect, it is necessary to make it contain 1.0% or more, but since weldability will deteriorate when it exceeds 2.0%, Mn content of a weld metal shall be 1.0-2.0%.

Al: ≦ 0.015%
Al acts as a deoxidizing element, but in the weld metal part, the effect of improving toughness due to deoxidation by Ti is rather large, and when the inclusion of Al oxide system increases, the absorbed energy of weld metal Charpy decreases, The Al content of the weld metal is set to 0.015% or less.

Cu: ≦ 0.5%, Ni: ≦ 3.0%, Cr: ≦ 1.0%, Mo: ≦ 1.0%
Like the base material, Cu, Ni, Cr, and Mo improve the hardenability even in the weld metal, so are included for bainite organization. However, when the amount increases, the amount of alloying elements added to the welding wire increases, resulting in a significant increase in wire strength, resulting in an obstacle in wire feedability during welding. For this reason, the upper limits of the Cu, Ni, Cr, and Mo contents of the weld metal are 0.5%, 3.0%, 1.0%, and 1.0%, respectively.

V: ≦ 0.1%
An appropriate amount of V is an effective element because it increases strength without degrading toughness and weldability. However, if it exceeds 0.10%, the toughness of the reheated portion of the weld metal is significantly degraded. For this reason, the V content of the weld metal is set to 0.1% or less.

Ti: 0.003-0.10%
Ti acts as a deoxidizing element in the weld metal and is effective in reducing oxygen in the weld metal. In order to obtain this effect, 0.003% or more is necessary, but when it exceeds 0.10%, excess Ti forms carbides and deteriorates the toughness of the weld metal. For this reason, the Ti content of the weld metal is set to 0.003 to 0.10%.

B: ≦ 0.0030%
In welded pipes for line pipes with low strength grades, B is effective to make the microstructure of weld metal into acicular ferrite. When the amount of B in the metal exceeds 0.0030%, a martensite structure is generated and the toughness is lowered. For this reason, the B content of the weld metal is set to 0.0030% or less.

O: ≦ 0.03%
Reduction of the amount of oxygen in the weld metal has an effect of improving toughness, and is particularly improved by making it 0.03% or less, so the O content in the weld metal is made 0.03% or less.

N: ≦ 0.008%
Reduction of solute N in the weld metal also has an effect of improving toughness, and is particularly remarkably improved by setting it to 0.008% or less. Therefore, the N content in the weld metal is set to 0.008% or less.

PcmW ≦ 0.2
PcmW is an index of the weldability of the weld metal calculated by the equation (9). Particularly, the seam welded portion 3 of the pipe (outer surface side weld metal 31, inner surface side weld, as in the T-cross portion 1 shown in FIG. The metal 32) shows a good correlation with the hardness (T-cross hardness, the dotted line 4 is the measurement position) after being subjected to the thermal effect that is received when the circumferential welding 2 between the pipes is performed.

FIG. 3 shows the result of arranging the hardness (T-cross hardness) after being affected by heat by PcmW. From FIG. 3, when PcmW is large and T-cross hardness is high, cold cracking is likely to occur in the pipe seam welded portion 3 (outer surface side weld metal 31, inner surface side weld metal 32) during circumferential welding. In order to satisfy the Vickers hardness of 300 points or less, which is a measure for preventing cracking, the upper limit of the PcmW value of the weld metal is set to 0.2.
PcmW = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 * B-12 * N-4 * O (9)
[Welding method]
Next, the reasons for limiting the end welding method after forming the steel pipe will be described. 5, in the present invention, hybrid welding method combining a gas shielded arc welding using a laser and Ar-CO ₂ gas shielded with CO ₂ gas shielded are shown schematically.

  In the hybrid welding method, a laser torch 6 and a gas arc welding torch 7 are arranged, and a weld bead 9 is formed by welding in one pool 8 by both of them. This can significantly improve the cooling rate of the weld.

  This is probably because the steel plate can be easily melted by applying high-density heat input to a narrow region by the preceding laser torch 6, and the weld metal can be sufficiently deposited even at the heat input level of the subsequent gas arc welding.

By using CO ₂ gas for the laser shield, the generation of blowholes and blowholes peculiar to laser welding is remarkably suppressed, and the gas arc welding shield is made of a mixed gas of Ar and CO ₂ to reduce the amount of oxygen in the weld metal. It can be kept low.

  FIG. 6 schematically shows a cross section of a weld using laser / arc hybrid welding according to the present invention. In the figure, 10 indicates a laser-arc hybrid weld, and 12 indicates a submerged arc weld.

  (A) is a case of thin-walled steel pipe, in the case of outer surface single layer welding by laser / arc hybrid welding, (b) is in the case of thick plate, in case of inner / outer surface single layer welding by laser / arc hybrid welding, (c) Laser-arc hybrid welding is used for inner surface single layer welding and outer surface is subjected to single-layer welding by conventional submerged arc welding.

  As the hardness of the HAZ part by laser-arc hybrid welding is sufficiently hard compared to conventional SAW, if the tube thickness is thick and one-layer laser-arc hybrid welding is not possible, as shown in Fig. 6 (b) Even when laser-arc hybrid welding is performed one layer at a time on the inner and outer surfaces of the pipe, the decrease in joint strength is small.

  In addition, as shown in FIG. 6 (c), even if the outer surface is subjected to conventional single layer welding by SAW welding, a sufficient strength is ensured in the HAZ portion on the inner surface, and the joint strength is equal to or higher than that of the base material. Can be satisfied.

  With laser-arc hybrid welding, it is possible to increase the strength of the seam weld while preventing cold cracking of the seam weld and cracking due to the seam weld metal being circumferentially welded and hardened. .

  Steel plates A to K were produced using the steel having the chemical composition shown in Table 1 under the hot rolling / accelerated cooling and reheating conditions shown in Table 2. The reheating was performed using an induction heating type heating device installed in the same line field as the accelerated cooling equipment.

  First, each steel plate was cut at 20 points by a shearing machine, and then the cut surface of the steel plate was examined by magnetic particle flaw detection to determine the number of cut end surfaces where cut cracks were observed. Here, even when a plurality of cracks could be confirmed in one end face, the number of cut cracks was set to 1 because there was only one end face. When no cut cracks were observed at all cut locations, (the number of cut crack occurrences 0) was considered good.

Next, specimens for microstructural confirmation were collected from each steel sheet, and the microstructural fraction and cementite grain size were investigated. Furthermore, full thickness tensile test specimens and DWTT test specimens conforming to API-5L, JIS Z2202 (1980) V-notch Charpy impact test specimens were collected from the center of the thickness, and steel sheet tensile tests, DWTT tests and Charpy impact tests were conducted. Tests were conducted to evaluate strength and toughness.

  In addition, the welding methods shown in Table 3 were mainly used for butt welding of steel plates with various changes in welding wires and welding methods to produce welded joints. Analytical samples were collected from the weld metal parts of each joint and subjected to chemical analysis. The analysis results are also shown in Table 3.

  In addition, a joint tensile test piece (with surplus) conforming to API-5L, a JIS Z2202 V-notch Charpy impact test piece taken at the position where the notch hits the weld metal and HAZ, A Charpy impact test was conducted to evaluate the strength and toughness of the weld.

  Furthermore, in accordance with JIS Z 3158, the y-type weld cracking test was conducted with the test environment controlled at a temperature of 30 ° C and a humidity of 80%. A test bead was welded to a specimen with a preheating temperature of 75 ° C using a hand-welded rod for 100 kgf-class high-strength steel that was left in the environment for 1 hour. Weld crack susceptibility was evaluated by the cross-sectional crack rate obtained from the observation results of the cross-section orthogonal to the test bead.

  In addition, gas arc welding was performed perpendicular to the welded joint, and a T-cross hardness test was performed on the prepared specimen.

  Table 4 summarizes the results of the microstructure investigation, strength / toughness investigation, weld joint strength / toughness investigation, weld crack susceptibility evaluation, and T-cross hardness results.

  None of the steels conforming to the present invention has cracks in the plate cutting experiment, has a base metal tensile strength exceeding 900 MPa, has a high base metal Charpy absorbed energy exceeding 200 J, and a DWTT ductile fracture surface ratio exceeding 85%. Satisfied.

  Furthermore, the joint strength was also equal to or greater than that of the base metal, and the weld metal and HAZ Charpy absorbed energy were also high values exceeding 100J. In addition, excellent weldability was exhibited in the y-type weld crack test and the T-cross test.

  On the other hand, in Comparative Example 2-2 where the same inner and outer surface SAW welding was performed, HAZ softening was remarkable, and as a result of fracture at the HAZ part in the joint tensile test, the joint tensile strength was significantly lower than the base metal tensile strength.

In Comparative Example 2-3, where the gas arc torch shield gas during laser-arc hybrid welding was CO ₂ gas and the oxygen content of the weld metal exceeded the upper limit, the Charpy absorbed energy of the weld metal was significantly reduced.

In Comparative Example 2-4, in which the laser torch shield gas was a mixed gas of Ar and CO ₂ , the weld metal had a defect that was thought to be a blowhole. Absorbed energy decreased.

  Comparative Example 6-2, in which the PcmW value of the weld metal exceeded the upper limit, had good weld joint strength and toughness, but exhibited low weldability with a T-cross hardness exceeding 300 points.

  In Comparative Example 10 where reheating was not performed after sheet rolling and accelerated cooling, cracks occurred in the sheet cutting experiment. In Comparative Example 11, where the heating rate during reheating after accelerated cooling was below the lower limit, the cementite in bainite and martensite was coarsened to 0.7 μm on average during heating, resulting in the base material Charpy absorbed energy and DWTT. The characteristic fell below the target.

  Furthermore, in Comparative Example 12 in which the cooling start temperature exceeded the upper limit, the ferrite transformation in the air cooling process before accelerated cooling was not sufficient, and as a result of the ferrite area ratio falling below 10%, the yield strength increased and the low yield ratio Could not be achieved.

  Further, in Comparative Example 13 where the amount of the base material C was below the lower limit, the base material tensile strength was below the target. In Comparative Example 14 in which the amount of base metal B exceeded the upper limit, ferrite transformation was suppressed by segregation of B austenite grain boundaries, resulting in a ferrite fraction of 0% in the microstructure of the steel, yield strength increasing, The yield ratio could not be achieved.

  In Comparative Example 15 in which the values calculated using the parameter formulas according to the present invention for the base materials Ca, O, and S were below the lower limit, the morphology control of MnS by Ca addition was insufficient, and the cleanliness by coarse MnS was insufficient. As a result of deterioration, the Charpy absorbed energy of the base material decreased.

  In Comparative Example 16, in which the PcmB value calculated by the chemical composition of the base metal exceeded the upper limit, high strength was achieved in the base metal and the weld zone, but weldability deteriorated because cracks occurred in the y-type weld crack test. .

The figure which shows the outer surface side hardness distribution of the welded steel pipe which performed inner and outer surface 1 layer submerged arc welding. Correlation diagram between cold crack prevention preheating temperature and Pcm value of steel. The correlation diagram of the weld metal HAZ hardness and PcmW value which were obtained by the T cross test. The figure explaining a T cross part. The schematic diagram explaining the laser arc hybrid welding method. Schematic diagram of laser-arc hybrid welded joint cross-section, (a) for thin-walled steel pipe, outer-layer single-layer welding by laser-arc hybrid welding, (b) for laser-arc hybrid welding when plate thickness is thick In the case of inner / outer surface single layer welding, (c) is a diagram showing a case where inner surface single layer welding is performed by laser / arc hybrid welding, and the outer surface is subjected to single layer welding by conventional submerged arc welding.

Explanation of symbols

DESCRIPTION OF SYMBOLS 1 T cross part 2 Circumferential welding 3 Seam welding part 31 Outer surface side welding metal 32 Inner surface side welding metal 4 Measurement position 6 Laser torch 7 Gas arc welding torch 8 Pool 9 Welding bead 10 Hybrid welding part 11 Submerged arc welding part

Claims (3)

% By mass
C: 0.03-0.12%
Si: 0.01 to 0.5%
Mn: 1.5 to 3%
Al: 0.01 to 0.08%
Nb: 0.01 to 0.08%
Ti: 0.005-0.025%
N: 0.001 to 0.01%
B: ≦ 0.0003%
Ca: 0.0005 to 0.01%
O: ≦ 0.003%
S: ≦ 0.001%
Furthermore, Cu: 0.01 to 0.5%
Ni: 0.01 to 1%
Cr: 0.01 to 0.5%
Mo: 0.01 to 0.5%
V: 0.01 to 0.1%
And the content of Ca, O, S satisfies the formula (1), the PcmB value calculated by the formula (2) satisfies PcmB ≦ 0.22, and the balance Fe and The composition of inevitable impurities,
Either ferrite + bainite, ferrite + martensite, or ferrite + bainite + martensite has an area fraction of 90% or more, and the area ratio of ferrite is 10 to 50%, and bainite and / or martensite A steel sheet having a microstructure with an average particle size of cementite of 0.5 μm or less and a tensile strength of 900 MPa or more and a yield ratio ≦ 85% is formed into a tubular shape by cold working, and then laser and Ar− using a CO ₂ gas shield are used. By the hybrid welding method combined with gas shielded arc welding using CO ₂ gas shield, the chemical composition of the weld metal is mass%,
C: 0.05-0.09%
Si: 0.1 to 0.4%
Mn: 1.0-2.0%
Al: ≦ 0.015%
Cu: ≦ 0.5%
Ni: ≦ 3.0%
Cr: ≦ 1.0%
Mo: ≦ 1.0%
V: ≦ 0.1%
Ti: 0.003-0.10%
B: ≦ 0.0030%
O: ≦ 0.03%
N: ≦ 0.008%
Containing
The PcmW value calculated by Equation (3) satisfies PcmW ≦ 0.2,
In addition, a method for producing an ultra-high-strength, high-deformability welded steel pipe excellent in the toughness of the base metal and the welded portion, wherein the butt portion is welded so as to become the remaining Fe and inevitable impurities.
1 ≦ (1-130 × [O]) × [Ca] / (1.25 × [S]) ≦ 3 (1)
PcmB = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 * B (2)
PcmW = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 * B-12 * N-4 * O (3)   In the manufacturing method of the ultra-high strength and high deformability welded steel pipe excellent in base material and welded portion toughness according to claim 1, butt welding is performed, and hybrid welding is performed by combining laser and gas shield arc welding for each inner and outer surface layers. A manufacturing method of ultra-high-strength, high-deformability welded steel pipe with excellent base metal and weld toughness.   In the manufacturing method of the super-high strength and high deformability welded steel pipe excellent in the base material and weld zone toughness according to claim 1, after performing butt welding and hybrid welding in which the inner surface side is combined with laser and gas shielded arc welding, A method of manufacturing an ultra-high-strength, high-deformability welded steel pipe with superior base metal and weld toughness, characterized by submerged arc welding of the outer surface.

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