JP2007260715A - Method for producing superhigh strength welded steel pipe - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing a superhigh strength welded steel pipe having a tensile strength of >800 MPa which has excellent brittle fracture propagation stop properties and is suitable as the one for transporting natural gas and crude oil. <P>SOLUTION: A steel sheet having a tensile strength of ≥800 MPa which has excellent brittle fracture propagation stop properties is subjected to cold working, so as to be formed into a pipe shape, and thereafter, the butted parts thereof 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, and particularly relates to an excellent brittle crack propagation characteristic and suitable for transportation of natural gas and crude oil.

  In recent years, line pipes used for transportation of natural gas and crude oil have been strengthened year by year in order to improve transport efficiency by increasing pressure and to improve local welding efficiency by reducing wall thickness, and already have X100 grade of API standard. The line pipe has been put into practical use, and an X120 grade exceeding a tensile strength of 900 MPa is also demanded.

  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.
JP 2003-293089 A JP 2002-173710 A JP 2000-355729 A JP-A-8-311548

  However, line pipes and line pipe steels described in the above-mentioned patent documents are premised on longitudinal seam welding by high heat input submerged arc welding, and depending on the plate thickness, the HAZ part is greatly softened at the longitudinal seam welded part. However, when a water pressure test using a real pipe is performed, there is a concern that the HAZ part with low strength breaks down.

  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.

  Accordingly, an object of the present invention is to provide a method for producing an ultra-high strength welded steel pipe that prevents HAZ softening of a longitudinal seam weld and is excellent in brittle crack propagation stopping characteristics.

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 base material having low tensile cracking resistance and low tensile cracking resistance and excellent tensile cracking strength of 800 MPa or more.
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 an excellent tensile strength 800MPa or more steel brittle crack propagation stop characteristics were molded into a tubular by cold working, the butted portion, CO gas shielded arc using laser and Ar-CO ₂ gas shielded with ₂ gas shielded A method for producing an ultra-high strength welded steel pipe, characterized by welding by a hybrid welding method combining welding.
2. 2. The method for producing an ultra high strength welded steel pipe according to 1, wherein the 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 tensile strength of 800 MPa or more, which is excellent in the brittle crack propagation stopping property,
% By mass
C: 0.03-0.12%
Si: ≦ 0.5%
Mn: 1.5 to 3.0%
P ≦ 0.010, S ≦ 0.002
Al: 0.01 to 0.08%
Cu: ≦ 0.8%
Ni: 0.01-3.0%
Cr: ≦ 0.8%
Mo: ≦ 0.8%
Nb: 0.01 to 0.08%
V: ≦ 0.10%
Ti: 0.005-0.025%
B: ≦ 0.003%
Ca: ≦ 0.01%
REM: ≦ 0.02%
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 in the temperature range of 950 ° C or lower is ≧ 67%, and the cooling rate is 20 to 80 ° C / s from the temperature range of 600 ° C or higher after the end of rolling. A steel plate obtained by starting cooling, stopping cooling in a temperature range of 250 ° C. or less, and immediately reheating to a temperature of 300 ° C. or more and 400 ° C. or less at a temperature rising rate of 5 ° C./s or more,
The chemical composition of the weld metal
% By 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 1 to 3, wherein the balance is Fe and inevitable 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%).

  According to the present invention, it is possible to produce a high-strength welded steel pipe having a tensile strength of 800 MPa or more that has a joint strength at the longitudinal seam portion that is equal to or higher than the tensile strength of the base metal and excellent in brittle crack propagation stopping characteristics, which is extremely useful industrially.

The present invention, after the excellent tensile strength 800MPa or more steel brittle crack propagation stop characteristics were molded into a tubular by cold working, laser and Ar-CO gas shielded arc using _two gas shielded with CO ₂ gas shielded A welded steel pipe is formed by welding the butt portion by a hybrid welding method combining welding.

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 of a steel sheet suitable as a steel sheet having a tensile strength of 800 MPa or more and excellent in brittle crack propagation stopping characteristics 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, and brings about softening resistance of HAZ by forming Nb and V carbides as described later. 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 increases remarkably and cold cracking occurs. In order to facilitate, 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: 1.5 to 3.0%
Mn acts as a hardenability improving element. In particular, in order to obtain a low temperature transformation structure necessary for achieving high strength in HAZ, addition of 1.5% or more is necessary. However, in the continuous casting process, the concentration in the central segregation part is significantly increased, exceeding 3.0%. If added, it causes delayed fracture at the segregation part, so the upper limit was 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: ≦ 0.8%, Cr: ≦ 0.8%, Mo: ≦ 0.8%
Cu, Cr, and Mo all act as a hardenability improving element. These are used to replace a large amount of Mn, and obtain a low-temperature transformation structure similar to Mn to contribute to the strengthening of the base metal and HAZ. However, these are expensive elements, and each is added in an amount of 0.8% or more. Even so, the effect of increasing the strength is saturated, so the upper limit was made 0.8%.

Ni: 0.1 to 3.0%
Ni is also a useful element that acts as a hardenability improving element and does not cause toughness deterioration even when added. In order to obtain this effect, addition of 0.1% or more is necessary, but since it is an expensive element, the upper limit was made 3.0%.

Nb: 0.01 to 0.08%
Nb is an element necessary for obtaining a desired HAZ strength by forming carbides and preventing temper softening of a weld heat-affected zone (HAZ) subjected to two or more thermal cycles.

  It also has the effect of expanding the austenite non-recrystallized region during hot rolling, and in order to make the non-recrystallized region up to 950 ° C., which is effective for fine graining, addition of 0.01% or more is necessary. It is. 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%
V is precipitation-hardened during multiple welding heat cycles due to the combined addition with Nb and contributes to the prevention of HAZ softening, but if added over 0.1%, precipitation hardening significantly reduces the HAZ toughness, so the upper limit is 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: ≦ 0.003%
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 0.003%, the effect is saturated, so the upper limit was made 0.003%.

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%.

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 correlates with a preheating temperature for preventing cold cracking in the HAZ part as a weld cracking sensitive composition. FIG. 1 shows the PCMB values of the low temperature cracking prevention preheating conditions of the HAZ part obtained by a low temperature cracking test conducted after giving various preheating temperatures to steels having various chemical compositions.

  From the figure, it is necessary to set the PcmB value to 0.22 or less in order to prevent HAZ cracking when the preheating temperature is allowed to 75 ° C. in the first layer welding during circumferential welding. 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.

Next, the reason for limiting the manufacturing method of the steel sheet will be described.
Heating temperature: 1000-1200 ° C
In performing hot rolling, the lower limit temperature for complete austenite is 1000 ° C. 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 upper limit was set 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 this temperature range, austenite grains expand and become fine grains in the thickness direction. The toughness of bainite steel obtained by accelerated cooling in this state is good. If the reduction amount is less than 67%, the effect of refining is insufficient and the toughness of bainite steel cannot be improved. Therefore, the lower limit of the cumulative reduction amount is set to 67%. In addition, the suitable range for aiming at a remarkable toughness improvement is 75% or more.

Cooling start temperature of accelerated cooling ≧ 600 ° C
After hot rolling, if the temperature at which accelerated cooling starts is low, proeutectoid ferrite is generated from the austenite grain boundaries during the air cooling process. When the cooling start temperature is lowered to about 600 ° C. or lower, most of the microstructure becomes ferrite and the strength is remarkably reduced. Therefore, the cooling start is performed at least from 600 ° C. or higher.

Accelerated cooling rate: 20-80 ° C / s
In order to suppress the ferrite transformation that causes a decrease in strength, it is necessary to perform accelerated cooling at 20 ° C./s or more. On the other hand, when the cooling rate exceeds 80 ° C./s, martensite transformation occurs particularly near the steel plate surface, and the strength of the steel plate increases. However, the toughness deterioration, particularly the Charpy absorption energy decrease, is significant, so the upper limit of the cooling rate is 80 C./s.

Cooling stop temperature for accelerated cooling: ≤250 ° C
In order to increase the strength of the steel sheet, it is necessary to lower the stop temperature of accelerated cooling to form a bainite or martensite structure that transforms at a low temperature. However, when the cooling stop temperature exceeds 250 ° C., the upper bainite structure has low toughness, so the cooling stop temperature is set to 250 ° C. or lower.

  Steel plates that have been subjected to low temperature transformation by accelerated cooling to increase strength will cause diffusible hydrogen in the steel to remain and cause cut cracks if air-cooled. 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 a temperature drop during the air cooling process, so it is desirable to start heating within 300 seconds, more preferably within 100 seconds. The reheating method may be either 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., and the DWTT characteristic is deteriorated. On the contrary, since it is possible to suppress the coarsening of cementite by increasing the heating rate, the heating rate during reheating was set to 5 ° C./s or more.

Reheating temperature: 300 ° C to 400 ° 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 400 ° C., the upper limit was set to 400 ° C. because the strength was 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 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 portion, the weld metal portion 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 of tensile strength ≧ 800 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 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 to make the microstructures into acicular ferrite. However, if a bainite structure is used to increase the tensile strength of 800 MPa or more, welding is required. When the amount of B in the metal exceeds 0.0030%, a martensitic 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 is an index of weldability of the weld metal, and it is good as the hardness (hereinafter referred to as T-cross hardness) after the seam welded portion of the pipe is subjected to the thermal effect when the pipe is circumferentially welded. Has a correlation.

  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 plate 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.

  Using steels (A to K) having the chemical composition shown in Table 1, steel plate No. 1 was used under the hot rolling / accelerated cooling and reheating conditions shown in Table 2. 1-15 were produced. 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, a full thickness tensile test piece and a DWTT test piece conforming to API-5L were sampled from each steel plate, and a JIS Z2202 (1980) V-notch Charpy impact test piece was taken from the central position of the plate thickness. , DWTT test (−30 ° C.) and Charpy impact test (−30 ° C.) were conducted 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 the welding wire and welding method 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 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 (−20 ° C.) was performed 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 controlled at a temperature of 30 ° C. and a humidity of 80%. A test bead was welded to a test body with a preheating temperature of 75 ° C. using a hand-welding rod for 100 kgf class high-strength steel left in this 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 so as to be orthogonal to the welded joint, and a T-cross hardness test was performed on the manufactured specimen.

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

  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 200 J and a DWTT ductile fracture surface ratio exceeding 85%. .

  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, in which seam welding was performed with SAW single-layer welding on both the inner and outer surfaces as usual, the welding heat was high, the HAZ softened, the HAZ fractured when the joint was pulled, and the joint strength was lower than the base metal.
In Comparative Example 2-3, in which the shield gas of the gas arc torch 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.

Further, 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 considered to be a blowhole, so it was broken by the weld metal when the joint was pulled. Charpy absorbed energy decreased.

  In Comparative Example 5-2, in which the PcmW value of the weld metal exceeded the upper limit, the weld joint strength and toughness were good, but the T-cross hardness exceeded 300 points.

  In Comparative Example 9 in which reheating was not performed after sheet rolling and accelerated cooling, cracks occurred in the sheet cutting experiment. Further, in Comparative Example 10 in which the cumulative reduction amount of the non-recrystallized region during rolling was below the lower limit, the base metal Charpy and DWTT characteristics were reduced. Furthermore, Comparative Example 11 in which the cooling stop temperature exceeded the upper limit could not achieve the target strength of 800 MPa or more.

  Further, in Comparative Example 12 in which the base metal C amount exceeded the upper limit, the base metal Charpy value, the weld metal Charpy value, and the HAZ Charpy value were deteriorated, and further, a low temperature crack was generated in the y-type weld crack test.

  Similarly, Comparative Example 13 in which the PcmB value exceeded the upper limit and Comparative Example 15 in which the amount of Mn exceeded the upper limit also caused low temperature cracking. On the other hand, in Comparative Example 14 in which the amount of the base material Si exceeded the upper limit, the base material Charpy and DWTT deteriorated.

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, where (a) is a single layer outer surface side welding of laser / arc hybrid welding, (b) is a single layer welding 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 (4)

After an excellent tensile strength 800MPa or more steel brittle crack propagation stop characteristics were molded into a tubular by cold working, the butted portion, CO gas shielded arc using laser and Ar-CO ₂ gas shielded with ₂ gas shielded A method for producing an ultra-high strength welded steel pipe, characterized by welding by a hybrid welding method combining welding.   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 tensile strength of 800 MPa or more, which is excellent in the brittle crack propagation stopping property,
% By mass
C: 0.03-0.12%
Si: ≦ 0.5%
Mn: 1.5 to 3.0%
P ≦ 0.010, S ≦ 0.002
Al: 0.01 to 0.08%
Cu: ≦ 0.8%
Ni: 0.01-3.0%
Cr: ≦ 0.8%
Mo: ≦ 0.8%
Nb: 0.01 to 0.08%
V: ≦ 0.10%
Ti: 0.005-0.025%
B: ≦ 0.003%
Ca: ≦ 0.01%
REM: ≦ 0.02%
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 in the temperature range of 950 ° C or lower is ≧ 67%, and the cooling rate is 20 to 80 ° C / s from the temperature range of 600 ° C or higher after the end of rolling. A steel plate obtained by starting cooling, stopping cooling in a temperature range of 250 ° C. or less, and immediately reheating to a temperature of 300 ° C. or more and 400 ° C. or less at a temperature rising rate of 5 ° C./s or more,
The chemical composition of the weld metal
% By 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 balance is Fe and inevitable 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%).

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