JP2007260716A - Method for producing ultrahigh strength welded steel pipe having excellent deformability - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing an ultrahigh strength welded steel pipe having excellent deformability which has a tensile strength of >800 MPa and is suitable as the one for transporting natural gas and crude oil. <P>SOLUTION: A steel sheet having a composition comprising, by mass, 0.03 to 0.12% C, ≤0.5% Si, 1.8 to 3.0% Mn, ≤0.010% P, ≤0.002% S, 0.01 to 0.08% Al, ≤0.7% Cu, 0.01 to 3.0% Ni, ≤1.0% Cr, ≤1.0% Mo, 0.01 to 0.08% Nb, ≤0.10% V, 0.005 to 0.025% Ti, ≤0.005% B, ≤0.01% Ca, ≤0.02% REM, ≤0.03% Zr, ≤0.01% Mg, 0.001 to 0.006% N and ≤0.22 PcmB, and the balance Fe with inevitable impurities, and having a yield ratio of ≤80% and a tensile strength of ≥800 MPa is subjected to cold working, so as to be formed into a pipe shape. Thereafter, the butted parts are welded by a hybrid welding process in which laser welding using CO<SB>2</SB>gas shield and gas shielded arc welding using Ar-CO<SB>2</SB>gas shield are combined. <P>COPYRIGHT: (C)2008,JPO&INPIT

Description

    The present invention relates to a method for producing an ultra-high strength welded steel pipe having a tensile strength of more than 800 MPa, particularly when the yield ratio of the base metal (ratio of yield strength to tensile strength) is as low as 80% or less and undergoes deformation such as compression and bending. And has sufficient toughness in the weld metal and weld heat-affected zone (HAZ) of the longitudinal seam weld, and the joint tensile strength exceeds the base metal strength. It relates to a good drop for transportation of natural gas and crude oil.

  In recent years, line pipes used for transportation of natural gas and crude oil have to solve various problems such as improvement of transportation efficiency by increasing pressure, improvement of on-site welding efficiency by thinning, and preservation of the global environment.

  In order to improve transport efficiency by increasing the pressure, it is effective to increase the strength of the steel sheet. To improve the efficiency of on-site welding by reducing the wall thickness, X100 grade line pipes have already been put into practical use under the API standard with excellent cold cracking susceptibility. An X120 grade with a tensile strength exceeding 900 MPa has also been prototyped.

  Regarding the method for producing a welded steel pipe for high-strength line pipes and a high-strength thick steel plate as the material, for example, in Patent Document 1, two-stage cooling is performed after hot rolling, and the second stage cooling stop temperature is set to 300 ° C. or less. Thus, a technique for achieving high strength 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 technology related to the component design for reducing the amount of alloying elements added to the base metal as in Patent Document 2 and obtaining high strength and high toughness in the weld metal of the longitudinal seam weld. .

  Patent Document 4 describes a steel sheet for ultra-high-strength steel pipe obtained by aging treatment after hot rolling, cooling a low C-high Cu steel.

  In order to preserve the global environment, even if a large deformation occurs in a line pipe due to a large earthquake or ground deformation in a frozen land zone, a high deformability is required to prevent cracking. The yield ratio is used as an index of high deformability, and the lower the yield ratio, the higher the critical strain for crack initiation.

  It is known that a low yield ratio is obtained by making the microstructure of a steel material a two-phase structure of a soft ferrite phase and a hard phase in which hard bainite and martensite are appropriately dispersed. As a production method for obtaining a structure in which the hard phase is appropriately dispersed in the soft phase as described above, quenching from a two-phase region of ferrite and austenite (Q ′) between quenching (Q) and tempering (T). ) Is disclosed.

Patent Document 6 discloses that a low yield ratio can be achieved by a ferrite + bainite + martensite structure in which a soft phase is processed ferrite.
JP 2003-293089 A JP 2002-173710 A JP 2000-355729 A JP-A-8-311548 JP-A-55-97425 Japanese Patent Laid-Open No. 08-209291

  However, the line pipes and line pipe steels described in Patent Documents 1 to 4 described above have a large softening of the HAZ part at the longitudinal seam welded part depending on the plate thickness when longitudinal seam welding is performed by high heat input submerged arc welding. However, since there is a concern that the HAZ part having a low strength is broken in a water pressure test using an actual pipe, it is necessary to reduce the welding heat input, resulting in a reduction in productivity.

  In order to compensate for the lack of joint strength due to HAZ softening, it is effective to increase the strength of the weld metal in the longitudinal seam weld. However, the amount of alloy elements in the weld metal is increased, and weld metal defects such as cold cracking occur. It becomes easier and the welding workability such as reworking is remarkably deteriorated.

  In addition, increasing the amount of alloying elements in the base metal increases the HAZ hardness at the circumferential welded part by multi-layer welding with low heat input that joins the pipes, and particularly increases the sensitivity to cold cracking at the first layer welded part. For this reason, it is necessary to manage the preheating of the groove and the workability at the site is reduced.

  In the technique described in Patent Document 5 that achieves a low yield ratio without impairing the high strength of the base material, it is necessary to perform heat treatment many times, resulting in a decrease in productivity and an increase in manufacturing cost.

  The toughening by the technique described in Patent Document 6 is obtained by actively developing a ferrite texture to improve the ductility-brittle fracture surface ratio. Occurs and the absorbed energy (Charpy impact value) in the Charpy impact test is rather lowered.

  Accordingly, an object of the present invention is to provide a method for manufacturing an ultra-high strength welded steel pipe that prevents HAZ softening of a longitudinal seam weld and has excellent deformability.

In order to solve the above-mentioned problems, the present inventors have conducted intensive studies on the following items regarding the base material and the longitudinal seam welding method.
1) Application of laser-arc hybrid welding to longitudinal seam welding: Maintaining the welding efficiency of conventional high heat input submerged arc welding while improving the cooling rate of the weld zone and increasing the strength of HAZ and weld metal Let welding conditions.

2) A process for producing a steel sheet having a tensile strength of 800 MPa or more and excellent toughness and having a microstructure satisfying a low yield ratio of 80% or less and excellent in low-temperature cracking sensitivity.
3) A weld metal component that suppresses weld cracking during longitudinal seam welding and achieves weld metal strength and toughness at a high cooling rate.

The present invention has been made on the basis of the knowledge obtained as a result of the above studies, that is, the present invention
1. After 80% yield ratio less and a tensile strength 800MPa or more steel sheets were molded into a tubular by cold working, the butted portion, the gas-shielded arc welding using a laser and Ar-CO ₂ gas shielded with CO ₂ gas shielded A method for producing an ultra-high strength welded steel pipe characterized by welding by a combined hybrid welding method.
2. 2. The method for producing an ultra-high strength welded steel pipe according to 1, wherein inner and outer surfaces of the butt portion are welded by the hybrid welding.
3. 2. The method of manufacturing an ultra high strength welded steel pipe according to 1, wherein an inner surface of the butt portion is welded by the hybrid welding and an outer surface is welded by submerged arc welding.
4). A steel plate having a yield ratio of 80% or less and a tensile strength of 800 MPa or more,
% By mass
C: 0.03-0.12%
Si: ≦ 0.5%
Mn: 1.8-3.0%
P ≦ 0.010%, S ≦ 0.002%
Al: 0.01 to 0.08%
Cu: ≦ 0.7%
Ni: 0.01-3.0%
Cr: ≦ 1.0%
Mo: ≦ 1.0%
Nb: 0.01 to 0.08%
V: ≦ 0.10%
Ti: 0.005-0.025%
B: ≦ 0.005%
Ca: ≦ 0.01%
REM: ≦ 0.02%
Zr: ≦ 0.03%
Mg: ≦ 0.01%
N: 0.001 to 0.006%
PcmB ≦ 0.22
Steel consisting of the balance Fe and inevitable impurities,
After reheating to 1000 to 1200 ° C, hot rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C or lower, and at a cooling rate of 20 to 80 ° C / s from the temperature range of 700 ° C or higher after the end of rolling. A steel plate obtained by starting cooling, stopping cooling in a temperature range of 400 to 650 ° C., immediately reheating to 600 to 700 ° C., and then air cooling to room temperature,
The chemical composition of the weld metal at the butt 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%
PcmW ≦ 0.2
The method for producing an ultra-high strength welded steel pipe according to any one of claims 1 to 3, wherein the remaining Fe and unavoidable impurities.

However, PcmB = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B
PcmW = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 × B-12 × N-4 × O
And each element shows content (mass%).
5). MA in the bainite of the said steel plate is 5 to 20% by an area rate, The manufacturing method of the ultra high strength welded steel pipe of 4 characterized by the above-mentioned.

  INDUSTRIAL APPLICABILITY According to the present invention, it is possible to produce an ultra-high strength welded steel pipe having a joint strength at a longitudinal seam portion that is equal to or higher than the tensile strength of the base material and excellent in deformability, and is extremely useful industrially.

The present invention provides a laser using a CO ₂ gas shield and an Ar—CO ₂ gas after forming a steel plate with a tensile strength of 800 MPa or more having excellent brittle crack propagation stopping characteristics and a yield ratio of 80% or less into a tubular shape by cold working. A welded steel pipe is formed by welding the butt portion by a hybrid welding method combined with gas shielded arc welding using a shield.

Figure 4 is a schematic view illustrating a hybrid welding method in combination with gas shielded arc welding using laser welding and Ar-CO ₂ gas shielded with CO ₂ gas shielded, hybrid welding 5, the welding direction, A laser torch 6 is arranged ahead of the gas arc welding torch 7.

  The laser torch 6 and the gas arc welding torch 7 are arranged so as to form a bead 9 as one pool welding in which the weld pools 8 formed by the respective weldings are combined into one. As a result, it is possible to weld the steel plate butt portion at a welding speed comparable to that of conventional submerged arc welding, and the cooling rate of the welded portion is significantly improved.

  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.

  The welding heat input when welding the base metal of the same thickness by the hybrid welding and the submerged arc welding is about ½ of the submerged arc welding by the hybrid welding.

  Therefore, when the pipe thickness is thick and penetration welding is not possible with one layer of laser / arc hybrid welding, even if laser / arc hybrid welding is performed on each of the inner and outer surfaces of the pipe, the decrease in joint strength is small. Moreover, even if one-layer welding is performed on the outer surface side by conventional SAW welding, a sufficient strength is ensured in the HAZ portion on the inner surface side, and a joint strength equal to or higher than that of the base material can be satisfied.

  FIG. 5 schematically shows a longitudinal seam welding method in the method for producing an ultra high strength welded steel pipe according to the present invention. When the plate thickness is thin, laser-arc hybrid welding outer surface side single layer welding (a), when it is thicker, laser-arc hybrid welding inner surface outer layer single layer welding (b), and when thicker, the inner surface side is laser-arced. Hybrid welding, the outer surface side is submerged arc welding (c).

Note that the use of CO ₂ gas as the laser welding shield gas significantly suppresses the generation of blowholes, and the gas arc welding shield gas is a mixed gas of Ar and CO ₂ to keep the amount of oxygen in the weld metal low. be able to.

  Next, the reasons for limiting the components suitable for a steel sheet having a yield ratio of 80% or less and a tensile strength of 800 MPa or more excellent in deformation performance in the present invention will be described. % Means mass%.

C: 0.03-0.12%
C contributes to an increase in strength by being supersaturated in a low temperature transformation structure. In order to obtain these effects, 0.03% or more of addition is necessary. However, if added over 0.12%, the hardness of the circumferential welded part of the pipe is remarkably increased and cold cracking is likely to occur. Therefore, the upper limit was made 0.12%.

Si: ≦ 0.5%
Since Si strengthens the solid solution regardless of the transformation structure, it is effective in increasing the strength of the base material and HAZ. However, if added over 0.5%, the toughness is significantly reduced, so the upper limit was made 0.5%.

Mn: 1.8-3.0%
Mn acts as a hardenability improving element. Furthermore, adding a large amount has the effect of reducing the amount of C that can be dissolved in the ferrite phase, and when the bainite transformation is performed from the austenite region of steel by accelerated cooling, the C concentration in the untransformed austenite region is increased. , MA production can be increased.

  As will be described later, in order to make the area ratio of MA 5% or more, it is necessary to add at least 1.8% or more. On the other hand, in the continuous casting process, the concentration in the center segregation part is remarkably increased, and if it exceeds 3.0%, it causes delayed fracture in the segregation part, so the upper limit is made 3.0%.

Al: 0.01 to 0.08%
Al acts as a deoxidizing element. Sufficient deoxidation effect can be obtained with addition of 0.01% or more, but if added over 0.08%, the cleanliness in the steel is lowered and the toughness is deteriorated, so the upper limit is 0.08%. It was.

P: ≦ 0.010%, S: ≦ 0.002%
Both P and S are present as inevitable impurities in the steel. In particular, the segregation at the center segregation portion is an element, and the upper limit is set to 0.010% and 0.002%, respectively, in order to suppress the decrease in toughness due to the segregation portion of the base material.

Cu: ≦ 0.7%
Cu acts as a hardenability-enhancing element and can replace a large amount of Mn addition. However, when added over 0.7%, Cu dissolved in supersaturation precipitates during reheating after accelerated cooling, and the yield strength of steel increases due to precipitation hardening, making it difficult to achieve a low yield ratio. Therefore, the upper limit is set to 0.7%.

Ni: 0.1 to 3.0%
Ni is also a useful element because it acts as a hardenability improving element and does not cause toughness deterioration when added. In order to obtain this effect, addition of 0.1% or more is necessary. However, when a large amount is added, Ni is concentrated on the scale generated on the surface of the slab during hot rolling slab heating, resulting in a change in scale properties. The upper limit is set to 3.0% because the scale bites easily during rolling and causes surface scratches.

Cr: ≦ 1.0%
Cr also acts as a hardenability-enhancing element and can be used as a substitute for adding a large amount of Mn. However, if added over 1%, the HAZ toughness deteriorates significantly, so when added, the upper limit is made 1%.

Mo: ≦ 1.0%
Mo also acts as a hardenability improving element, and can be used as a substitute for adding a large amount of Mn. However, since it is an expensive element and the increase in strength is saturated even if it is added in excess of 1%, when it is added, the upper limit is made 1%.

Nb: 0.01 to 0.08%
Nb has the effect of expanding the austenite non-recrystallized region at the time of hot rolling, and in order to make the non-recrystallized region up to 950 ° C., addition of 0.01% or more is necessary. On the other hand, if added over 0.08%, the toughness of the HAZ is remarkably impaired, so the upper limit is made 0.08%.

V: ≦ 0.1%
When V is added in combination with Nb, it precipitates and hardens during multiple welding heat cycles and contributes to the prevention of HAZ softening. However, if added over 0.1%, precipitation hardening significantly reduces the HAZ toughness. Has an upper limit of 0.1%.

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. Addition of 0.005% or more is necessary to obtain the required pinning effect, but if added over 0.025%, carbides are formed, and the toughness deteriorates significantly due to precipitation hardening. The upper limit was 0.025%.

B: ≦ 0.005%
B segregates at the austenite grain boundaries and suppresses ferrite transformation, thereby contributing particularly to the prevention of HAZ strength reduction. However, even if added over 0.005%, the effect is saturated, so the upper limit was made 0.005%.

Ca: ≦ 0.01%
Ca is an element effective in controlling the form of sulfide in steel, and when added, suppresses the generation of MnS harmful to toughness. However, if added over 0.01%, a CaO-CaS cluster is formed and the toughness is deteriorated, so the upper limit was made 0.01%.

REM: ≦ 0.02%
REM is also an effective element for controlling the form of sulfide in steel, and when added, it suppresses the generation of MnS harmful to toughness. However, since it is an expensive element and the effect is saturated even if added over 0.02%, the upper limit was made 0.02%.

Zr: 0.0005 to 0.03%
Zr forms carbonitrides in steel and brings about a pinning effect that suppresses the coarsening of austenite grains, particularly in the weld heat affected zone. In order to obtain a sufficient pinning effect, addition of 0.0005% or more is necessary. However, if over 0.03% is added, the cleanliness in the steel is remarkably lowered, leading to a reduction in toughness. , When added, the upper limit is made 0.03%.

Mg: 0.0005 to 0.01%
Mg is produced as fine oxides in the steel during the steel making process, and has a pinning effect that suppresses austenite grain coarsening, particularly in the weld heat affected zone. In order to obtain a sufficient pinning effect, addition of 0.0005% or more is necessary. However, if added over 0.01%, the cleanliness in the steel is lowered and the toughness is lowered. When doing so, the upper limit is made 0.01%.

N: 0.001 to 0.006%
N is usually present as an inevitable impurity in steel, but TiN that suppresses austenite coarsening is formed by adding Ti as described above. In order to obtain the required pinning effect, 0.001% or more must be present in the steel. However, if it exceeds 0.006%, it was heated to 1450 ° C or more in the weld zone, particularly in the vicinity of the melting line. When TiN decomposes in HAZ, the upper limit was made 0.006% because the adverse effect of solute N was significant.

PcmB ≦ 0.22
PcmB (= C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B) correlates with a preheating temperature for preventing cold cracking in the HAZ part as a weld cracking susceptibility composition.

  FIG. 1 is a summary of PCMB values of cold cracking prevention preheating conditions of the HAZ part obtained by cold cracking tests conducted after giving various preheating temperatures to 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. It was set to 22.

  In consideration of workability at the pipeline laying site, it is desirable that the pipe preheating temperature is low. From this viewpoint, the preferable range of PcmB is 0.20 or less.

P: ≦ 0.010%, S: ≦ 0.002%
Both P and S are present as inevitable impurities in the steel. In particular, the segregation at the center segregation portion is an element, and the upper limit is set to 0.010% and 0.002%, respectively, in order to suppress the decrease in toughness due to the segregation portion of the base material.

Next, the reason for limiting the manufacturing method of the raw steel plate will be described.
Heating temperature: 1000-1200 ° C
When hot rolling is performed, the steel slab is heated to 1000 ° C. or more in order to austenite. On the other hand, if the steel slab is heated to a temperature exceeding 1200 ° C., even if pinning is performed with TiN, the austenite grain growth is remarkable and the base material toughness deteriorates, so the upper limit is set to 1200 ° C.

Cumulative reduction at 950 ° C or lower ≥ 70%
In the present invention, since Nb addition is 950 ° C. or less in the austenite non-recrystallized region, the austenite grains are stretched in the temperature region and the austenite grains are stretched, and the bainite that is transformed by the subsequent accelerated cooling is fine. And toughness is improved.

  In the present invention, since hard MA is dispersed in order to achieve a low yield ratio, the toughness of the parent phase bainite portion needs to be sufficiently high. If the cumulative reduction amount is less than 70%, the fine reduction is insufficient and the toughness is reduced due to the influence of island martensite (also referred to as “MA”), so the cumulative reduction amount is 70% or more. Let's say. Preferably, a cumulative reduction amount of 75% or more is required.

Cooling start temperature of accelerated cooling ≧ 700 ° C
After hot rolling, if the temperature at which accelerated cooling is started is low, proeutectoid ferrite is generated from the austenite grain boundaries during the air cooling process, which causes a reduction in the strength of the base metal. 700 ° C.

Accelerated cooling rate: 20-80 ° C / s
In order to achieve a high strength of 800 MPa or more, the microstructure needs to be a bainite-based structure. For this reason, accelerated cooling is performed after hot rolling. When the cooling rate is less than 20 ° C./s, transformation is performed at a relatively high temperature, so that sufficient strength cannot be obtained.

  On the other hand, when the cooling rate exceeds 80 ° C./s, it is difficult to control the cooling stop temperature described later, and martensite transformation occurs particularly near the surface, and the base material toughness is significantly reduced. Let s.

Cooling stop temperature for accelerated cooling: 400-650 ° C
In the present invention, the cooling stop temperature management of accelerated cooling is an important manufacturing condition. In the present invention, since the C-concentrated untransformed austenite existing after reheating is transformed into MA during the subsequent air cooling, it is necessary to stop the cooling in the temperature range where untransformed austenite exists during the bainite transformation.

  If the cooling stop temperature is less than 400 ° C., the bainite transformation is completed, so MA is not generated during air cooling, and a low yield ratio cannot be achieved. On the other hand, if it exceeds 650 ° C., C is consumed in the pearlite that precipitates during cooling and MA is not generated, so the upper limit is set to 650 ° C.

Reheating temperature after stopping cooling: 600-700 ° C
By reheating immediately after accelerated cooling, C can be concentrated in untransformed austenite and MA can be generated in the subsequent air cooling process.

  When the time until the start of reheating is long, untransformed austenite decreases due to the temperature drop during that time, and the amount of MA generated in the air cooling process after heating decreases, so it is desirable to perform reheating within 300 seconds. Preferably, it is within 100 seconds.

  Furthermore, if the reheating temperature is less than 600 ° C., C concentration to austenite does not occur sufficiently, and the required MA amount cannot be ensured. On the other hand, if the reheating temperature exceeds 700 ° C., the bainite transformed by accelerated cooling becomes austenite again and sufficient strength cannot be obtained, so the reheating temperature is specified to be 600 ° C. or more and 700 ° C. or less.

  There is no need to set the temperature holding time at the reheating temperature. Further, in the cooling process after reheating, MA is generated regardless of the cooling rate, so cooling after reheating is not particularly defined, but basically it is preferably air cooling.

  The microstructure of the base material obtained by limiting the above components and the steel sheet production conditions is a structure in which hard MA is formed as the second phase between the bainite laths of the bainite structure that is transformed by continuous cooling.

  By stopping the cooling in the middle of transformation and immediately reheating it, the strength of the bainite phase is moderately reduced due to the tempering effect.

  At the same time, when the cooling is stopped, the untransformed austenite is enriched with C in the steel and a large amount of MA is likely to be formed. As a result, a soft tempered bainite phase and a hard MA are dispersed into a microstructure. Decreases.

  Since MA is finely and uniformly dispersed, the toughness is not lowered, and the matrix is tempered, which is rather improved.

  When the structure is mainly composed of ferrite due to insufficient cooling rate of accelerated cooling, it becomes difficult to achieve a tensile strength of 800 MPa or more. On the other hand, when martensite is organized, the toughness decreases although sufficient strength can be secured.

  It is desirable that MA in bainite is dispersed in an area ratio of 5 to 20%. If the area ratio is less than 5%, the yield ratio is not sufficiently lowered. Therefore, the MA area ratio is preferably at least 5% or more.

  On the other hand, when the area ratio exceeds 20%, the toughness of the base metal is remarkably deteriorated, so the upper limit is desirably 20% or less.

  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 reasons for limiting the additive elements of the weld metal will be described.
C: 0.05-0.09%
Also in the weld metal, C is an important element as a steel strengthening element. In particular, in order to achieve overmatching of the joint part, the weld metal part also needs to have a tensile strength ≧ 800 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 upper limit was made 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 upper limit was made 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 ≧ 800 MPa cannot be achieved by the conventional acicular ferrite structure, 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 if it exceeds 2.0%, weldability deteriorates, so the upper limit was made 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, Do not add aggressively, and set the upper limit to 0.015%.

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, as the amount increases, the amount of alloying elements added to the welding wire increases, resulting in a significant increase in wire strength. As a result, the wire feedability during welding is impaired. It was set to 0%, 1.0%, and 1.0%.

V: ≦ 0.1%
An appropriate amount of V is an effective element because it increases the strength without degrading toughness and weldability. However, if it exceeds 0.10%, the toughness of the reheated part of the weld metal is significantly degraded, so the upper limit is set to 0. 0.1%.

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, a content of 0.003% or more is necessary. However, if it exceeds 0.10%, the excess Ti forms carbides and degrades the toughness of the weld metal, so the upper limit. Was 0.03%.

B: ≦ 0.0030%
In welded pipes for line pipes with low strength grades, it is effective to add B in order to make the microstructures into acicular ferrite. However, in order to increase the tensile strength to 800 MPa or higher, when using a bainite structure, weld metal When the amount of B exceeds 0.0030%, a martensite structure with low toughness is generated, so the upper limit was made 0.0030%.

O: ≦ 0.03%
Reduction of the amount of oxygen in the weld metal has an effect of improving toughness. In particular, the upper limit is set to 0.03% because it is remarkably improved by setting it to 0.03% or less.

N: ≦ 0.008%
Reduction of solute N in the weld metal also has an effect of improving toughness. In particular, the upper limit is set to 0.008% because it can be remarkably improved by setting it to 0.008% or less.

PcmW ≦ 0.2
PcmW (= C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 × B−12 × N−4 × O) is an indicator of weldability of the weld metal, and the seam weld of the pipe is a circle between the pipes. It has a good correlation with the hardness (hereinafter referred to as T-cross hardness) after being subjected to the thermal effect when circumferential welding is performed.

  3 shows the T-cross part 1, 2 is circumferential welding, 3 is vertical seam welding, 31 is the outer surface side of the vertical seam welding 3, 32 is the inner surface side of the vertical seam welding 3, and 4 is T-cross hardness. The T-cross hardness is measured in the tube thickness direction at a position 2 mm on the HAZ side from the bond part of the circumferential weld 2 and the maximum hardness of the obtained hardness distribution is obtained. Define.

  Fig. 2 shows the relationship between T-cross hardness and PcmW. When PcmW is large and T-cross hardness is high, cracking is likely to occur at the pipe seam weld during circumferential welding. In order to satisfy the standard Vickers hardness of 300 points or less, the upper limit of the PcmW value of the weld metal was set to 0.2.

  Steel plates AK 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.

  A sample for microstructural observation was taken from the central part of the plate width of the obtained steel plate, and a plate thickness section parallel to the rolling longitudinal direction was mirror-polished, and then MA was revealed using a two-step etching method. Thereafter, a 10-view microstructure photograph was taken at a magnification of 2000 using a surface scanning microscope (SEM), and the area ratio of MA in the photograph was measured with an image analyzer.

  Next, after re-polishing the sample, the microstructure was revealed by using the nital etching method, and the area ratio of bainite and the presence of other phases were confirmed by the same method as MA.

  From each steel plate, a full thickness tensile test piece conforming to API-5L and a V-notch Charpy impact test piece of JIS Z2202 (1980 revision) from the central position of the plate thickness are collected, and a steel plate tensile test and Charpy impact test (test temperature) -30 ℃) to evaluate strength and toughness.

  In addition, the welding methods shown in Table 3 were mainly used for butt welding of steel sheets with various changes in welding wires and welding methods. 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 V-notch Charpy impact test piece of JIS Z2202 at the position where the notch hits the weld metal and HAZ, and a weld joint tensile test and A Charpy impact test (test temperature -20 ° C) was conducted to evaluate the strength and toughness of the weld.

  Furthermore, a y-type weld cracking test was performed in accordance with JIS Z 3158. The test environment was a temperature of 30 ° C. and a humidity of 80%, and a test bead was welded to a test body with a preheating temperature of 75 ° C. using a hand-welded rod for 100 kgf class high-strength steel 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.

  Moreover, the gas arc welding which simulated the circumference welding so that it was orthogonal to a welded joint was implemented, and the T-cross hardness test was done with the produced test body. A hardness test (Hv 5 kg) was performed in the plate thickness direction at a position 2 mm on the HAZ side from the bond portion of the circumferential weld 2, and the maximum hardness of the obtained hardness distribution was obtained.

  Table 4 summarizes the microstructure, strength and toughness investigation results of the base metal, the strength and toughness investigation results of welded joints, the evaluation of weld crack sensitivity, and the T-cross hardness results.

  The examples of the present invention include a base metal tensile strength of 800 MPa or more, a yield ratio of 80% or less, a Charpy absorbed energy of 250 J or more, a joint tensile strength of 800 MPa or more, a weld metal and HAZ Charpy absorbed energy of 100 J or more, and an oblique y crack test (preheating). The target value was no cracking at a temperature of 75 ° C. and T-cross hardness (Hv5) of 300 or less.

  None of the steels suitable for the present invention cracked in the plate cutting experiment, had a base metal tensile strength exceeding 800 MPa, and satisfied a high Charpy absorbed energy exceeding 250 J.

  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. Further, no cracks were observed in the y-type weld cracking test, and the hardness of the T-cross test was 300 or less, indicating excellent low temperature cracking sensitivity.

  On the other hand, in the comparative example in which any of the component composition of the steel sheet, the manufacturing conditions, or the chemical component of the weld metal is outside the scope of the present invention, the strength / toughness investigation result of the base metal, the strength / toughness investigation result of the welded joint, and the diagonal y Either the crack test result or the T-cross hardness test result did not satisfy the target value of the present invention.

The correlation diagram of the 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 hardness test. The schematic diagram explaining laser arc hybrid welding. BRIEF DESCRIPTION OF THE DRAWINGS FIG. 1 is a diagram for explaining a longitudinal seam weld according to the present invention, wherein (a) is a single layer welding on the outer surface side of laser / arc hybrid welding, (b) is a single layer welding on the inner / outer surface side of laser / arc hybrid welding, and (c) is an inner surface side. Shows the case of laser-arc hybrid welding and the outer surface side of submerged arc welding.

Explanation of symbols

DESCRIPTION OF SYMBOLS 1 T-cross part 2 Circumferential welding 3 Longitudinal seam welding 31 Outer surface side of vertical seam welding 32 Inner surface side of vertical seam welding 4 Measurement position of hardness test for obtaining T-cross hardness 5 Hybrid welding 6 Laser torch 7 Gas arc welding torch 8 Weld pool 9 Bead

Claims (5)

After 80% yield ratio less and a tensile strength 800MPa or more steel sheets were molded into a tubular by cold working, the butted portion, the gas-shielded arc welding using a laser and Ar-CO ₂ gas shielded with CO ₂ gas shielded A method for producing an ultra-high strength welded steel pipe, characterized by welding by a combined hybrid welding method.   The method for manufacturing an ultra high strength welded steel pipe according to claim 1, wherein the inner and outer surfaces of the butt portion are welded by the hybrid welding.   The method for producing an ultra-high strength welded steel pipe according to claim 1, wherein an inner surface of the butt portion is welded by the hybrid welding and an outer surface is welded by submerged arc welding. A steel plate having a yield ratio of 80% or less and a tensile strength of 800 MPa or more,
% By mass
C: 0.03-0.12%
Si: ≦ 0.5%
Mn: 1.8-3.0%
P ≦ 0.010%, S ≦ 0.002%
Al: 0.01 to 0.08%
Cu: ≦ 0.7%
Ni: 0.01-3.0%
Cr: ≦ 1.0%
Mo: ≦ 1.0%
Nb: 0.01 to 0.08%
V: ≦ 0.10%
Ti: 0.005-0.025%
B: ≦ 0.005%
Ca: ≦ 0.01%
REM: ≦ 0.02%
Zr: ≦ 0.03%
Mg: ≦ 0.01%
N: 0.001 to 0.006%
PcmB ≦ 0.22
Steel consisting of the balance Fe and inevitable impurities,
After reheating to 1000 to 1200 ° C, hot rolling is performed so that the cumulative reduction amount ≧ 70% in the temperature range of 950 ° C or lower, and at a cooling rate of 20 to 80 ° C / s from the temperature range of 700 ° C or higher after the end of rolling. A steel plate obtained by starting cooling, stopping cooling in a temperature range of 400 to 650 ° C., immediately reheating to 600 to 700 ° C., and then air cooling to room temperature,
The chemical composition of the weld metal at the butt 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%
PcmW ≦ 0.2
The method for producing an ultra-high strength welded steel pipe according to any one of claims 1 to 3, wherein the remaining Fe and unavoidable impurities.
However, PcmB = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B
PcmW = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 60 × B-12 × N-4 × O
And each element shows content (mass%). MA in the bainite of the said steel plate is 5 to 20% by an area rate, The manufacturing method of the ultra high strength welded steel pipe of Claim 4 characterized by the above-mentioned.

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