JP4811166B2 - Manufacturing method of super high strength welded steel pipe exceeding tensile strength 800MPa - Google Patents

Description

  The present invention relates to a method for manufacturing an ultra-high strength welded steel pipe having a tensile strength exceeding 800 MPa, and in particular, toughness, cut cracking resistance, and low temperature cracking susceptibility of base metal, weld heat affected zone (HAZ) and longitudinal seam welded portion. The present invention relates to a material suitable 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. .
JP 2003-293089 A JP 2002-173710 A JP 2000-355729 A

  However, in the line pipes and line pipe steels described in Patent Documents 1 to 3 described above, when vertical seam welding is performed by high heat input submerged arc welding, the HAZ portion is greatly softened at the vertical seam weld depending on the plate thickness. However, since there is a concern that the HAZ part having a low strength is broken in a water pressure test using a real pipe, the welding heat input must be lowered, and the productivity is lowered.

  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 weld by multi-layer welding with low heat input that joins the pipes, and particularly increases the sensitivity to cold cracking at the first layer weld. For this reason, it is necessary to manage the preheating of the groove and the workability at the site is reduced.

  Furthermore, when high strength is achieved by lowering the cooling stop temperature after hot rolling and changing the microstructure to a low temperature transformation structure, when the steel sheet as cooled is cut into the required size by shearing, Due to the remaining diffusible hydrogen, cracks parallel to the plate surface (hereinafter referred to as cut cracks) occur.

  On the other hand, when heat treatment such as tempering is performed after accelerated cooling, hydrogen in the steel is sufficiently diffused, so that cut cracks can be suppressed, but cementite precipitates and coarsens in the microstructure during the heat treatment process, reducing toughness. In particular, the DWTT (Drop Weight Tear Test) characteristic for evaluating the brittle crack propagation stop characteristic deteriorates.

  Then, this invention aims at providing the manufacturing method of the welded steel pipe for line pipes using the steel for line pipes which solved the said subject in tensile strength 800MPa or more.

  In order to solve the above-mentioned problems, the present inventors have conducted intensive studies on the following items with respect to the base material characteristics 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 manufacturing process of a steel sheet having excellent deformation performance, DWTT characteristics, cut cracking resistance, and low temperature cracking sensitivity at a tensile 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. The microstructure is a mixed structure of ferrite, tempered martensite, and lower bainite, and ACR obtained from Ca, O, S in the steel satisfies 0-2, tensile strength of 800 MPa or more, YR 85% or less, and uniform elongation 5% or more after the steel sheet was formed into a tubular by cold working, welding the butted portions, the hybrid welding in combination with 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.
However, ACR = (Ca− (0.18 + 130 × Ca) × O) / (1.25 × S), and Ca, O, and S indicate the content (%) in steel.
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). The steel plate is
% By mass
C: 0.04 to 0.12%
Si: 0.01 to 0.5%
Mn: 1.80 to 2.50%
Al: 0.01 to 0.08%
P ≦ 0.010%, S ≦ 0.002%
Cu: 0.01 to 0.8%
Ni: 0.1 to 1.0%
Cr: 0.01 to 0.8%
Mo: 0.01 to 0.8%
Nb: 0.01 to 0.08%
V: 0.01-0.10%
Ti: 0.005-0.025%
Ca: 0.0005 to 0.01%
N: 0.001 to 0.006%
PcmB ≦ 0.22
Steel consisting of the balance Fe and inevitable impurities,
After heating to 1000 to 1200 ° C., hot rolling is started, rolling is performed so that the rolling end temperature is in the temperature range of Ar ₃ transformation point or higher and Ar ₃ + 100 ° C. or lower, and then Ar ₃ −50 ° C. or higher, Cooled from the temperature range below the Ar ₃ transformation point to the cooling stop temperature in the temperature range above the martensite transformation start temperature Ms and above 300 ° C. at a cooling rate above the martensite formation critical cooling rate Vcrm satisfying the equation (1). Thereafter, a steel plate obtained by holding for 60 s to 300 s within a cooling stop temperature ± 50 ° 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 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
ACR = (Ca− (0.18 + 130 × Ca) × O) / (1.25 × S)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
And each element shows content (%).
logVcrm = 2.94−0.75 × (β−1) (1)
Here, β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo
5. Furthermore, after the steel sheet is held for 60 s to 300 s within the cooling stop temperature ± 50 ° C., it is rapidly heated from the temperature to a temperature range of 450 ° C. or higher and below the Ac ₁ transformation point at a rate of 1 ° C./s or higher. The method for producing an ultra-high strength welded steel pipe according to 4, wherein the method is obtained by tempering.

  INDUSTRIAL APPLICABILITY According to the present invention, it is possible to produce an ultra-high strength welded steel pipe having a tensile strength of 800 MPa or more, which is excellent in safety and suitable for transportation of natural gas and crude oil, and is extremely useful industrially.

In the present invention, the microstructure is a mixed structure of ferrite, tempered martensite, and lower bainite, and the ACR obtained from Ca, O, and S in the steel satisfies 0 to 2, the tensile strength is 800 MPa or more, and YR is 85% or less and uniform. after a 5% elongation or more steel sheets were molded into a tubular by cold working, butt portion by a hybrid welding method in combination with gas shielded arc welding using a laser and Ar-CO ₂ gas shielded with CO ₂ gas shielded This is characterized in that a welded steel pipe is obtained.

  YR is a value expressed in% by multiplying the value obtained by dividing the yield strength at a nominal strain of 0.5% by the tensile strength by 100.

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. (A) shows the case where the plate thickness is thin, outer surface side single layer welding of the laser / arc hybrid welding 9, (b) shows the case where the plate thickness is thicker, and inner / outer surface side single layer welding of the laser / arc hybrid welding 9, (C) shows a thicker case, where the inner surface side is laser / arc hybrid welding 9 and the outer surface side is submerged arc welding 10.

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, in the present invention, the microstructure is a mixed structure of ferrite, tempered martensite, and lower bainite, and the ACR obtained from Ca, O, and S in the steel satisfies 0 to 2, the tensile strength is 800 MPa or more and YR is 85% or less and A steel sheet having a uniform elongation of 5% or more will be described.

  The microstructure ensures a tensile strength of 800 MPa or more, YR 85% or less and uniform elongation of 5% or more, and to obtain excellent deformation performance and DWTT characteristics. Ferrite, tempered martensite and lower bainite are excellent in strength, toughness and deformation performance. A mixed tissue is used. Moreover, ACR in steel composition shall be 0-2.

  ACR is defined by (Ca− (0.18 + 130 × Ca) × O) / (1.25 × S), and is a parameter related to MnS. When the range is from 0 to 2, CaS is generated, MnS that is harmful to toughness and serves as a trapping site for diffusible hydrogen can be reduced, and a steel sheet having excellent cut cracking resistance and low-temperature cracking sensitivity can be obtained. In addition, Ca, O, and S show content (%) in steel.

  The reason for limiting the components suitable for the steel sheet of the present invention will be described. In the description,% is mass%.

C: 0.04 to 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, 0.04% or more of addition is necessary. However, if added over 0.12%, the hardness of the circumferential welded portion of the pipe increases significantly, and low temperature cracking is likely to occur. Therefore, the upper limit was made 0.12%.

Si: 0.01 to 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 to 2.5%
Mn acts as a hardenability improving element. In particular, in order to obtain a low temperature transformation structure to achieve high strength in HAZ, addition of 1.8% or more is necessary. However, in the continuous casting process, the concentration in the central segregation part is remarkably increased, and the addition exceeds 2.5%. This causes delayed fracture at the segregation part, so the upper limit was set to 2.5%.
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.

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.

Cu: 0.01-0.8%, Cr: 0.01-0.8%, Mo: 0.01-0.8%
Cu, Cr, and Mo all act as hardenability improving elements, but the effect cannot be obtained at 0.01% or less. These are used for replacement of a large amount of Mn, and similarly contribute to increase the strength of the base material and HAZ by obtaining a low-temperature transformation structure, but they are expensive elements and each is 0.8% or more. Even if added, the effect of increasing the strength is saturated, so the upper limit was made 0.8%.

Ni: 0.1 to 1.0%
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, since it is an expensive element, the upper limit was made 1.0%.

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

  In addition, there is an effect of expanding the austenite non-recrystallized region during 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: 0.01 to 0.1%
V is precipitated and hardened during multiple welding thermal 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%.

Ca: 0.0005 to 0.01%
Ca is an element effective for controlling the form of sulfides in steel, and when added, the formation of MnS harmful to toughness is suppressed, but if it is less than 0.0005%, the effect cannot be obtained. 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%.

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 an effective pinning effect, 0.001% or more must be present in the steel. However, if it exceeds 0.006%, HAZ heated to 1450 ° C or higher near the weld, particularly in the vicinity of the melting line. When TiN decomposes at, the upper limit was made 0.006% because the solute N had a bad influence.

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

  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 during circumferential welding of pipes, the PcmB value must be 0.22 or less to prevent HAZ cracking when the pipe preheating temperature is allowed to 75 ° C, so the upper limit was set to 0.22. .

  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 reasons for limiting the manufacturing method of the steel sheet will be explained.
[Production conditions]
The molten steel having the above composition is melted by a normal melting means such as a converter or an electric furnace, and is made into a steel material such as a slab by a normal casting method such as a continuous casting method or an ingot-bundling method. Is preferred.

  There are no particular restrictions on the steel making method, but from the economical point of view, it is desirable to perform the steel making process by the converter method and the slab casting by the continuous casting process. The melting method and the casting method are not limited to the methods described above.

1. Slab heating-rolling conditions The steel material is heated to a temperature at which it becomes an austenite single phase structure. The heating temperature of the steel material is preferably 1000 to 1200 ° C. in order to austenite the steel material. When the heating temperature of the steel material is less than 1000 ° C., the hot deformation resistance is too high, and the rolling reduction per time cannot be taken high, and the productivity is lowered.

  Further, when a precipitate-forming element such as V or Nb is contained, these elements are not sufficiently dissolved in austenite, and it is difficult to sufficiently exhibit the effects of these elements. On the other hand, when the heating temperature exceeds 1200 ° C., the crystal grains become coarse, and the amount of scale loss and the frequency of furnace repairs increase. For this reason, the heating temperature of the steel material was limited to a range of 1000 to 1200 ° C.

The heated steel material is subjected to hot rolling at a rolling end temperature of not less than Ar ₃ transformation point and not more than Ar ₃ + 100 ° C.

2. Heat treatment
After the end of rolling, cooling is performed at a cooling rate of not less than the Martensite critical cooling rate Vcrm from the temperature range of Ar ₃ −50 ° C. to the Ar ₃ transformation point to a temperature range of 300 ° C. or less below the Ms point. After holding for 60 s to 300 s within 50 ° C., cool to room temperature.

Starting temperature of quenching is less than Ar 3 _-50 ° C., since ferrite is increased significantly in the Quenching starting tissue, be subjected to a hardening process can not be obtained the desired microstructure, the desired high strength and high toughness Can not be secured.

If the quenching start temperature is higher than the Ar ₃ transformation point, pro-eutectoid ferrite cannot be obtained, and deformation performance satisfying YR 85% or less and uniform elongation 5% or more cannot be obtained. Therefore, to limit the cooling start temperature in the range of _Ar 3 -50 ° C. or more Ar ₃ or less transformation point.

The quenching cooling rate is a cooling rate equal to or higher than the martensite formation critical cooling rate Vcrm. In the present invention, the martensite formation critical cooling rate Vcrm is a cooling rate defined by the following equation (3).
logVcrm = 2.94-0.75 * (β-1)
(Β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo) (3)
(Here, Vcrm: Martensite formation critical cooling rate (° C./s) and martensite formation critical cooling rate Vcrm means a cooling rate containing a martensite structure in a fraction of 90% or more of the entire structure. .

  By applying a quenching process that cools to a quenching stop temperature in the temperature range of 300 ° C. or more at a martensite transformation start temperature Ms or less at a cooling rate of the martensite generation critical cooling rate Vcrm or more, it is partially at each position in the plate thickness direction. Martensite is generated first.

  By partially generating martensite, strain due to expansion during martensite transformation is generated at the interface between the martensite and untransformed austenite, and this strain energy transforms untransformed austenite into lower bainite. In addition to being easier, the lower bainite phase can be produced in a finer amount and in a larger amount than in the prior art.

  If the quenching cooling rate is less than the martensite formation critical cooling rate Vcrm, the amount of coarse bainite generated before the martensite transformation increases, and the strain generation due to the martensite transformation becomes insufficient, and the desired effect is achieved. I can't get it.

  In addition, when the quenching and cooling stop temperature exceeds the Ms point, the strain generation effect due to the formation of martensite cannot be expected, the transformation to the lower bainite phase is insufficiently promoted, and it is generated during isothermal holding or air cooling. Increases the amount of island martensite harmful to toughness.

  On the other hand, when the quenching and cooling stop temperature is less than 300 ° C., the diffusion of C becomes insufficient, and carbide effective for crack propagation resistance does not precipitate inside the bainitic ferrite. For this reason, the quenching and cooling stop temperature is set to a temperature in the temperature range of 300 ° C. or lower from the Ms point. In addition, Preferably, it is a temperature range below 350 degreeC below Ms point.

  Next, after the cooling is stopped at the quenching cooling stop temperature in the above-described range, the steel temperature is kept within the cooling stop temperature ± 50 ° C. for 60 s to 300 s, and then air-cooled to room temperature.

  By maintaining the quenching cooling stop temperature within ± 50 ° C for 60 s to 300 s, martensite is self-annealed, while the transformation of untransformed austenite to lower bainite is promoted, and a mixed structure of tempered martensite and lower bainite is obtained. be able to.

  Moreover, it becomes possible to reduce the amount of acicular island-shaped martensite harmful to toughness formed between laths of martensite.

  In the isothermal transformation within 60 s, the lower bainite transformation is not completed and high strength and high toughness cannot be obtained, and when it is kept for longer than 300 s, the structure is coarsened and the strength is lowered. For this reason, the holding time in the temperature range is set to 60 s to 300 s, preferably 60 s to 100 s.

Further, in the case of particularly improving toughness, the steel plate is held for 60 to 300 s within the cooling stop temperature ± 50 ° C., and then from the temperature to a temperature range of 450 ° C. or more and Ac ₁ transformation point or less ₁ ° C./s or more. Tempering is carried out by rapid heating at a heating rate of.

  It should be tempered at least within 300 s after cooling is stopped in order to suppress a decrease in strength due to coarsening of the structure.

If the heating temperature is lower than 450 ° C., the effect of improving toughness can not be almost obtained, since when the Ac ₁ transformation point or more temperature reduction in strength occurs, the heating temperature is less than Ac ₁ transformation point 450 ° C. or higher. Further, if the rate of temperature rise is less than 1 ° C./s, the toughness is improved, but the strength is remarkably reduced, so the rate of temperature rise is 1 ° C./s or more.

  While martensite is self-annealed, the transformation of untransformed austenite to lower bainite is promoted, and a mixed structure of tempered martensite and lower bainite is obtained. Thereby, toughness can be improved with almost no deterioration in strength.

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

  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 explained.
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 0.1%, the toughness of the reheated portion 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 (Hv5) of 300 points or less, the upper limit of the PcmW value of the weld metal was set to 0.2.

  Various steel plates were produced under the hot rolling / accelerated cooling and reheating conditions shown in Table 2 using steels A to I having chemical compositions shown in Table 1. Reheating was performed using an induction heating type heating device installed in the same line as the accelerated cooling equipment.

  The obtained steel plate was cut at 20 points with a shearing machine, and then the cut surface of the steel plate was examined by magnetic particle inspection to determine the number of cut end faces 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.

  And the case where the cutting crack was not recognized in all the cutting | disconnection locations was made favorable (the number of cutting crack generation | occurrence | production 0).

  From the obtained steel plate, a full-thickness tensile test piece and DWTT test piece in accordance with API-5L were collected, and a V-notch Charpy impact test piece of JIS Z2202 (1980) was taken from the central position of the plate thickness. A test (test temperature: −30 ° C.) and a Charpy impact test (test temperature: −30 ° C.) were conducted to evaluate strength and toughness.

  The welding methods shown in Table 3 were used to butt-weld steel sheets obtained by variously changing the welding wire and welding method to produce a welded joint. 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 and a V-notch Charpy impact test piece of JIS Z2202 were collected, and a tensile test and a Charpy impact test of the welded joint (collection position: 2 mm below the outer surface side surface) , Notch position: weld metal, HAZ, test temperature: −30 ° C.) to evaluate the strength and toughness of the weld. In addition, the groove shape of welding followed the groove shape at the time of a blank pipe manufacture.

  The low temperature cracking susceptibility was subjected to a y-type weld cracking test in accordance with JIS Z 3158. The test atmosphere was a temperature of 30 ° C., 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 for 1 hour in the environment.

  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. The T-cross hardness test was performed on a test body in which gas arc welding was performed so as to be orthogonal to the welded joint to create a T-cross portion.

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

  Fitting No. 1 to 10 are joints using the steel of the present invention, and the base materials constituting the joints do not generate any cracks in the plate cutting experiment, have a base metal tensile strength exceeding 800 MPa, and high Charpy absorbed energy exceeding 200 J. Further, the DWTT ductile fracture surface ratio exceeding 85% was exhibited, and deformation performance of YR 85% or less and uniform elongation 5% or more was satisfied.

  Furthermore, the joint strength was also equal to or higher than that of the base metal, and the weld metal and HAZ Charpy absorbed energy were excellent values exceeding 100 J. In addition, no cracks were observed in the y-type weld cracking test, and the T-cross test showed low weld hardness and excellent weldability.

On the other hand, when the gas arc torch shield gas during laser-arc hybrid welding was CO ₂ gas, the oxygen content of the weld metal exceeded the upper limit. No. 11 significantly decreased the Charpy absorbed energy of the weld metal.

Further, joint No. of shielding gas laser torch was a mixed gas of Ar and CO ₂ In No. 12, defects that were thought to be blowholes remained in the weld metal, which caused the weld metal to fracture during joint tension and reduced Charpy absorbed energy.

  Fitting No. whose weld metal PcmW value exceeded the upper limit. No. 13 had good weld joint strength and toughness, but its T-cross hardness (Hv5) exceeded 300 and was inferior in cold cracking susceptibility.

  On the other hand, a joint No. whose rolling finish temperature exceeded the upper limit of the present invention was used. In the base material No. 14, the austenite grains were insufficiently refined. As a result, the Charpy absorbed energy value of the base material was low, and the DWTT ductile fracture rate was 40% or less.

In addition, a joint No. whose cooling start temperature exceeded the upper limit of the present invention was used. The base material of No. 15 did not undergo ferrite transformation below the Ar ₃ transformation point, so the YR was high, the uniform elongation was less than 5%, and the deformation performance deteriorated.

  Fitting No. In the base material No. 16, the cooling stop temperature after rolling exceeds the upper limit of the present invention, martensite transformation does not occur, a bainite main structure is formed, and the bainite substructure is coarsened due to the cooling stop at a higher temperature. Joint tensile strength decreased with yield strength.

  In addition, the joint No. Although the base metal of No. 17 had a cooling stop temperature after rolling lower than the lower limit of the present invention, and the lower bainite main structure was not obtained and became a tempered martensite main structure, the strength was high, but Charpy absorbed energy The value was low, and the DWTT ductile fracture surface ratio was 55%.

  In addition, the joint No. in which the cooling rate after rolling was lower than the lower limit of the present invention was used. The strength of the 18 base material was significantly reduced. Fitting No. whose holding time at the cooling stop temperature ± 50 ° C. was below the lower limit of the present invention. Although the base material of 19 increased in strength due to an increase in the martensite structure, the volume fraction of the lower bainite structure was not sufficient, the Charpy absorbed energy value was low, and the DWTT ductile fracture surface ratio was as low as 55%.

  Fitting No. In the base material No. 20, the holding time at the cooling stop temperature ± 50 ° C. exceeded the upper limit of the present invention, the base metal strength was low, and the DWTT ductile fracture surface ratio was also reduced to 70%.

Fitting No. In the base material of 21, the heating temperature after stopping the cooling exceeded the upper limit of the present invention and exceeded the Ac ₁ transformation point of the steel sheet. As a result, α-γ reverse transformation occurred and a large amount of island martensite was generated. As a result of the decrease in the volume fraction of the bainite structure, the strength decreased.

  Fitting No. in which the heating rate during online heating after cooling stopped was below the lower limit of the present invention. The base material No. 22 showed a high base material strength, but the Charpy absorbed energy value was low, and the DWTT ductile fracture surface ratio was reduced to 20%.

  In addition, the joint No. The base material of No. 23 had an ACR value outside the range of the present invention, and as a result of an increase in MnS-based sulfides, the DWTT ductile fracture surface ratio decreased.

  Fitting No. In the base material No. 24, the amount of Mn added to the steel sheet exceeds the upper limit outside the range of the present invention, the base metal Charpy value, the weld metal Charpy value, and the HAZ Charpy value are deteriorated. Occurred.

  On the other hand, the joint No. in which the C addition amount of the steel sheet exceeded the upper limit of the present invention. Similarly, cold cracking occurred in the 25 base metal.

The correlation diagram of the cold crack prevention preheating temperature of steel and a PcmB value. 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, 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 10 Submerged arc welding

Claims (5)

The microstructure is a mixed structure of ferrite, tempered martensite, and lower bainite, and ACR obtained from Ca, O, S in the steel satisfies 0-2, tensile strength of 800 MPa or more, YR 85% or less, and uniform elongation 5% or more the steel sheet after forming the tubular by cold working, welding the butted portions, the hybrid welding in combination with 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
However, ACR = (Ca− (0.18 + 130 × Ca) × O) / (1.25 × S), and Ca, O, and S indicate the content (%) in steel.   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. The steel plate is
% By mass
C: 0.04 to 0.12%
Si: 0.01 to 0.5%
Mn: 1.80 to 2.50%
Al: 0.01 to 0.08%
P ≦ 0.010%, S ≦ 0.002%
Cu: 0.01 to 0.8%
Ni: 0.1 to 1.0%
Cr: 0.01 to 0.8%
Mo: 0.01 to 0.8%
Nb: 0.01 to 0.08%
V: 0.01-0.10%
Ti: 0.005-0.025%
Ca: 0.0005 to 0.01%
N: 0.001 to 0.006%
PcmB ≦ 0.22
Steel consisting of the balance Fe and inevitable impurities,
After heating to 1000 to 1200 ° C., hot rolling is started, rolling is performed so that the rolling end temperature is in the temperature range of Ar ₃ transformation point or higher and Ar ₃ + 100 ° C. or lower, and then Ar ₃ −50 ° C. or higher, Cooled from the temperature range below the Ar ₃ transformation point to the cooling stop temperature in the temperature range above the martensite transformation start temperature Ms and above 300 ° C. at a cooling rate above the martensite formation critical cooling rate Vcrm satisfying the equation (1). Thereafter, a steel plate obtained by holding for 60 s to 300 s within a cooling stop temperature ± 50 ° 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 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
ACR = (Ca− (0.18 + 130 × Ca) × O) / (1.25 × S)
Ms = 517-300C-11Si-33Mn-22Cr-17Ni-11Mo
And each element shows content (%).
logVcrm = 2.94−0.75 × (β−1) (1)
Here, β (%) = 2.7C + 0.4Si + Mn + 0.45Ni + 0.8Cr + Mo Furthermore, after the steel sheet is held for 60 s to 300 s within the cooling stop temperature ± 50 ° C., it is rapidly heated from the temperature to a temperature range of 450 ° C. or higher and below the Ac ₁ transformation point at a rate of 1 ° C./s or higher. The method for producing an ultra-high strength welded steel pipe according to claim 4, wherein the method is obtained by tempering.

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