JP2012193446A - Steel plate for high-ductile high-strength welded steel pipe, steel pipe, and method for manufacturing the same - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a steel plate for a fully high-ductile ultrahigh strength welded pipe equivalent to an X150 grade, and a steel pipe and a method for manufacturing the same.SOLUTION: The steel plate for a high-ductile ultrahigh-strength welding steel pipe contains, by mass%, C: 0.09-0.11%, Si: 0.05-0.20%, Mn: 1.0-1.5%, Al: 0.01-0.08%, Cu: 2.0-4.0%, Nb: 0.05-0.07%, Ti: 0.015-0.025%, one or more selected from Cr: 0.05-0.6%, Mo: 0.05-0.6%, V: 0.01-0.1%, and B: 0.0005-0.003%, and the balance being Fe and inevitable impurities. The metal structure is bainite. In addition, the product of the tensile strength (MPa) in the rolling direction and the uniform elongation (%) in one direction is 8,500 or more.

Description

The present invention relates to a steel sheet and a steel pipe for a high ductility welded steel pipe having a particularly high uniform elongation in a steel sheet and a steel pipe for an ultra-high strength line pipe exceeding a tensile strength of 1150 MPa, and a method for producing the same. This line pipe is used for transportation of natural gas and crude oil.

  In recent years, line pipes used for transportation of natural gas and crude oil have become stronger year by year in order to improve transport efficiency by increasing pressure and to improve local welding efficiency by reducing wall thickness. So far, X100 grade line pipes have been put into practical use according to the API standard, and further up to X120 grade, which has a tensile strength exceeding 900 MPa, has been desired to be further strengthened.

  Regarding the manufacturing method of such a high-strength line pipe welded steel pipe and a high-strength thick steel plate as a raw material thereof, for example, in Patent Document 1, in consideration of earthquake resistance, the microstructure is ferrite + bainite, or ferrite + martensite, Or the technique which achieves high strength and a low yield ratio by making it ferrite + bainite + martensite and improves the deformation performance of a welded steel pipe is disclosed. Patent Document 2 discloses a technique related to an ultra-high strength thick steel plate that satisfies a tensile strength of 1150 MPa.

JP 2006-307334 A JP 2006-45644 A

  However, for example, when trying to put into practical use an ultra-high strength steel pipe having a tensile strength exceeding 1150 MPa such as X150 grade, the combination of a soft ferrite structure and a hard bainite structure and / or martensite structure described in Patent Document 1 High deformability can be imparted with a tensile strength of up to about 1000 MPa. On the other hand, in Patent Document 2, it is necessary to obtain a substantially 100% martensite structure in order to obtain a high strength of 1150 MPa or more, and 0.15% or more of C is required to ensure the strength. In such a high C design, not only low temperature but also high temperature weld cracking susceptibility is required, and an additional process for preventing cracking is required in the welded steel pipe manufacturing process, and the ductility is poor due to the single phase structure. It is difficult to ensure the safety of the line.

  The present invention relates to a steel sheet and a steel pipe for an ultra-high strength welded steel pipe corresponding to X150 grade, which has sufficient ductility even if the strength is increased to a tensile strength exceeding 1150 MPa without considering the addition of high C in consideration of weldability. It aims at providing the manufacturing method.

  The inventors first studied various steel strengthening methods in order to increase the C content to about 0.11% and increase the tensile strength to 1150 MPa, and actively performing precipitation strengthening. It was found that the strength can be secured even at low C. Furthermore, after the steel microstructure is bainite, a large amount of Cu is supersaturated and dissolved, and then dispersed and precipitated in bainite finely and in a large amount, thereby inhibiting dislocation motion and increasing the work hardening rate. It has been found that uniform elongation is remarkably improved because Cu itself has ductility.

  Furthermore, in order to clarify the optimum Cu precipitation form, the inventors set the basic chemical composition of steel to 0.11% C, 0.20% Si, 1.45% Mn, 0.02% Al in mass%. , 0.05% Nb, 0.020% Ti, and small steel ingots with various changes in Cu from 1.5% to 5.0% were prepared, subjected to hot rolling and accelerated cooling. 15mm thick steel plate with a bainite structure as a comparative material, on the other hand, after hot rolling and accelerated cooling, followed by accelerated cooling followed by rapid heating and tempering to minimize temper softening of bainite Prepared. In rapid heating and tempering, the heating temperature was variously changed to change the Cu precipitation state in bainite.

  From these steel sheets, a full thickness tensile test piece based on API-5L is taken in parallel with the rolling longitudinal direction, a tensile test is performed, and the obtained tensile strength and uniform elongation are subjected to rapid heat treatment for Cu precipitation. The amount of increase in tensile strength and uniform elongation due to Cu precipitation was calculated by comparing the case of performing the above and the case of accelerated cooling.

Next, three thin film samples for a transmission electron microscope each having a thickness of 0.3 mm were taken in parallel with the thickness direction from the rolling longitudinal cross section of the steel plate subjected to the rapid heat treatment for Cu precipitation, and transmitted. The Cu precipitates in the bainite lath were observed with an electron microscope at a magnification of 100,000. Photographs of Cu deposits taken for 3 fields of view for each thin film sample, a total of 9 fields, are image-analyzed, and the number of precipitates per unit observation area (1 μm ² ) and the average value of the diameters of each precipitate are measured with an image analyzer. Quantified. And the relationship with the above-mentioned amount of tensile strength increase and the amount of uniform elongation was derived from the number of Cu precipitates and the precipitate size in the bainite of the steel sheet.

As a result, as shown in FIG. 1, the amount of increase in tensile strength when Cu was precipitated was 1.0 × 10 ³ (number of particles / μm ² ) per unit observation area (1. In the above case, it was found that an increase in strength of 100 MPa or more can be obtained particularly when the number of Cu precipitated particles having a particle diameter of 40 nm or less is 1.0 × 10 ³ or more. At this time, as shown in FIG. 2, an increase in uniform elongation was also observed, and it was found that the increase in uniform elongation was larger as the average particle size of the Cu precipitated particles was larger. The present invention is obtained by further examining these findings, and the gist thereof is as follows.

  In the first invention, the steel composition is in mass%, C: 0.09 to 0.11%, Si: 0.05 to 0.20%, Mn: 1.0 to 1.5%, Al: 0 0.01 to 0.08%, Cu: 2.0 to 4.0%, Nb: 0.05 to 0.07%, Ti: 0.015 to 0.025%, and Cr: 0.005%. Containing one or more selected from 05 to 0.6%, Mo: 0.05 to 0.6%, V: 0.01 to 0.1%, B: 0.0005 to 0.003%, High ductility ultra-high strength, characterized by comprising the balance Fe and inevitable impurities, the metal structure being bainite, and the product of tensile strength (MPa) and uniform elongation (%) in the rolling direction being 8500 or more It is a steel plate for welded steel pipes.

In the second invention, the average particle size of Cu precipitates precipitated in the bainite of the metal structure is 40 nm or less, and the number of Cu precipitate particles per unit area of the steel plate cross section is 1.0 × 10 ³ particles / μm ² or more. Further, the steel sheet for high ductility ultra-high strength welded steel pipe according to the first invention, characterized in that the tensile strength in the rolling direction exceeds 1150 MPa.

  The third invention further includes, in mass%, Ca: 0.0005 to 0.01%, REM: 0.0005 to 0.02%, Zr: 0.0005 to 0.03%, Mg: 0.0005. It is a steel plate for high ductility ultra high strength welded steel pipes according to the first or second invention, characterized by containing one or more selected from -0.02%.

  The fourth invention further comprises Ni: 1.0 to 4.0% by mass%, and the high ductility ultra high strength welded steel pipe according to any one of the first to third inventions Steel plate.

  5th invention is manufactured using the steel plate in any one of 1st thru | or 4th invention, The product of pipe longitudinal direction tensile strength (MPa) and uniform elongation (%) is 8500 or more. It is a high ductility ultra high strength welded steel pipe characterized by

  A sixth aspect of the present invention is the high ductility ultra high strength welded steel pipe according to the fifth aspect of the present invention, wherein the tensile strength in the longitudinal direction of the pipe exceeds 1150 MPa.

7th invention starts hot rolling, after heating the steel slab which has the steel composition in any one of 1st thru | or 4th invention to 1100-1200 degreeC, The cumulative reduction rate in 950 degrees C or less Is 50% or more, the rolling is finished at the Ar ₃ transformation point or higher, and then accelerated cooling is performed at a cooling rate of 40 to 80 ° C./sec and a cooling stop temperature of 400 to 500 ° C. The steel sheet for high ductility ultra-high strength welded steel pipe, which is rapidly cooled to 550 to 650 ° C. at a temperature rising rate of 5 ° C./sec or more from the temperature range of (cooling stop temperature −50 ° C.) and then air-cooled. It is a manufacturing method.

  An eighth invention is characterized in that the steel plate according to any one of the first to third inventions is cold-formed into a cylindrical shape, the opposite end faces are butt welded, and then expanded or reduced in diameter. It is a manufacturing method of a high ductility super high strength welded steel pipe.

  According to the present invention, a steel plate and a steel pipe for an ultra-high-strength welded steel pipe corresponding to the X150 grade having sufficient ductility even if the tensile strength exceeds 1150 MPa without increasing the addition of high C in consideration of weldability, and its It has become possible to provide a manufacturing method.

It is a figure explaining the relationship between the number of Cu precipitation particles, an average particle diameter, and TS increase amount. It is a figure explaining the relationship between the number of Cu precipitation particles, an average particle diameter, and the amount of uniform elongation increase.

  Hereinafter, the present invention will be described in detail.

1. About component composition Below, the component composition of this invention is demonstrated. In addition,% in a component composition shall be mass% altogether.

C: 0.09 to 0.11%
C contributes to the formation of cementite and alloy carbides with Nb, Ti, V, and Mo in the bainite structure, and increases the strength of bainite. If the content is less than 0.09%, the effect of increasing the strength is insufficient, so 0.09% or more is to be contained in order to obtain a high strength with a tensile strength exceeding 1150 MPa. On the other hand, if the content exceeds 0.11%, the hot cracking of the weld metal becomes significant due to dilution of C into the weld metal of the pipe, so the C content is in the range of 0.09 to 0.11%. .

Si: 0.05-0.20%
Si containing 0.05% or more strengthens the solid solution regardless of the transformation structure, and is effective in increasing the strength of the base material and HAZ. However, if the content exceeds 0.20%, the toughness is remarkably reduced, so the Si content is in the range of 0.05 to 0.20%.

Mn: 1.0 to 1.5%
Mn acts as a hardenability improving element. In order to suppress the ferrite transformation and make the metal structure of the base material mainly composed of bainite, it is necessary to contain 1.0% or more. On the other hand, since the effect is saturated even if the content exceeds 1.5%, the amount of Mn is set to a range of 1.0 to 1.5%.

Al: 0.01 to 0.08%
Al acts as a deoxidizing element. A sufficient deoxidation effect can be obtained when the content is 0.01% or more. However, if the content exceeds 0.08%, the cleanliness in the steel is lowered and the toughness is deteriorated. The range is 01 to 0.08%.

Cu: 2.0 to 4.0%
Cu can be precipitated in a large amount as metal Cu in bainite, thereby improving the balance between strength and ductility by precipitation strengthening without reducing ductility. In order to obtain precipitation Cu of an amount necessary for precipitation strengthening to a low carbon bainite structure to make the tensile strength exceed 1150 MPa, it is necessary to contain 2.0% or more of Cu, but it exceeds 4.0%. If contained, the slab cracking and the like become prominent and the load in the refining process becomes high. Therefore, the Cu amount is set to a range of 2.0 to 4.0%. Preferably, it is 2.0 to 3.0% of range.

Nb: 0.05-0.07%
Nb forms alloy carbide to exert precipitation strengthening action, and Nb also has an effect of expanding the austenite non-recrystallized region during hot rolling, and in order to obtain these effects together, 0.05 % Or more must be contained. On the other hand, if the content exceeds 0.07%, the toughness is remarkably impaired, so the Nb content is in the range of 0.05 to 0.07%.

Ti: 0.015-0.025%
Ti is added to form carbides and strengthen precipitation. Ti also forms nitrides, suppresses austenite grain growth during slab heating, and contributes to refinement of the bainite structure. In order to achieve both carbide formation for precipitation strengthening and nitride formation for pinning, a content of 0.015% or more is necessary. On the other hand, if the content exceeds 0.025%, the toughness is remarkably impaired, so the Ti content is in the range of 0.015 to 0.025%.

  In the present invention, in order to obtain a bainite-based structure, it is necessary to further contain one or more selected from Cr, Mo, V, and B, which are hardenability improving elements.

Cr: 0.05-0.6%
Cr acts as a hardenability-enhancing element and can replace a large amount of Mn addition. In order to obtain this effect, it is necessary to contain 0.05% or more, but if added over 0.6%, the weld heat affected zone (hereinafter also referred to as HAZ) toughness deteriorates significantly, so Cr When it contains, let the quantity be 0.05 to 0.6% of range.

Mo: 0.05-0.6%
Mo also acts as a hardenability-enhancing element and can replace a large amount of Mn addition. In order to obtain this effect, it is necessary to contain 0.05% or more, but it is an expensive element, and even if it exceeds 0.6%, the strength increase is saturated, so it contains Mo. In that case, the amount is in the range of 0.05 to 0.6%.

V: 0.01 to 0.1%
V is precipitation-hardened during multiple welding thermal cycles due to the combined addition with Nb, contributing to the prevention of HAZ softening. This effect is manifested by the inclusion of 0.01% or more, but if it exceeds 0.1%, precipitation hardening remarkably deteriorates the HAZ toughness. The range is 0.1%.

B: 0.0005 to 0.003%
B segregates at austenite grain boundaries and suppresses ferrite transformation, thereby contributing particularly to prevention of HAZ strength reduction. In order to acquire this effect, it is necessary to contain 0.0005% or more, but even if it contains more than 0.003%, the effect is saturated. The range is 0005 to 0.003%.

  Although the basic component composition of the present invention is as described above, when the purpose is to improve the toughness of the base metal or the welded portion, one or more of Ca, REM, Zr, and Mg can be contained as a selective element.

Ca: 0.0005 to 0.01%
Ca is an element effective for controlling the form of sulfide in steel, and has an action of suppressing the generation of MnS harmful to toughness. In order to acquire this effect, it is preferable to contain 0.0005% or more, but if it contains more than 0.01%, a CaO-CaS cluster is formed and the toughness is deteriorated. The amount is preferably in the range of 0.0005 to 0.01%.

REM: 0.0005 to 0.02%
REM is also an element effective for controlling the form of sulfide in steel, and has an action of suppressing the generation of MnS harmful to toughness. In order to obtain this effect, it is preferable to contain 0.0005% or more, but since it is an expensive element and the effect is saturated even if it contains more than 0.02%, when REM is contained, The amount is preferably in the range of 0.0005 to 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, it is preferable to contain 0.0005% or more, but if it exceeds 0.03%, the cleanliness in the steel is remarkably lowered and the toughness is lowered. Therefore, when it contains Zr, it is preferable to make the quantity into the range of 0.0005 to 0.03%.

Mg: 0.0005 to 0.02%
Mg produces fine oxides in the steel during the steelmaking process, and in particular, has a pinning effect that suppresses the austenite grain coarsening in the weld heat affected zone. In order to obtain a sufficient pinning effect, it is preferable to contain 0.0005% or more, but if it exceeds 0.02%, the cleanliness in the steel is lowered and the toughness is lowered. When it contains, it is preferable to make the quantity into 0.0005 to 0.02% of range.

  In the present invention, Ni can further be contained for the purpose of suppressing so-called Cu cracking during hot rolling.

Ni: 1.0-4.0%
When hot-rolling Cu-containing steel, when the slab heating temperature exceeds the melting point of Cu, a liquid phase is generated, and Cu cracks starting from this are generated, which are so-called Cu cracks. In order to suppress this Cu cracking, it is effective to contain Ni in an amount more than half the Cu content. Is preferably in the range of 1.0 to 4.0%. Furthermore, the Ni content is more preferably ½ or more of the Cu content and less than or equal to the Cu content.

  The balance other than the above components is composed of Fe and inevitable impurities.

2. About metal structure The metal structure (micro structure) of a steel plate or a steel pipe base material is bainite. In the case of a ferrite single-phase structure having a soft microstructure, a double-phase structure of ferrite and bainite, or a double-phase structure of ferrite and martensite, it is difficult to achieve a target tensile strength exceeding 1150 MPa. On the other hand, when the microstructure is mainly martensite, the strength can be sufficiently secured, but the ductility is remarkably lowered. For this reason, in order to achieve both high strength and high ductility, the microstructure needs to be bainite. The smaller the area fraction of the microstructure other than bainite, the better.

  However, when the area fraction of the microstructure other than bainite is small, the influence is small, so other metal structures of 5% or less in total area fraction, that is, ferrite, pearlite, cementite, martensite, islands One or more kinds of martensite (also referred to as MA) may be contained. In addition, when residual austenite is present as a microstructure other than bainite, it can be expected to improve the elongation accompanying processing-induced transformation, but after plastic working, it becomes hard martensite, which may cause a decrease in ductility. Therefore, the area fraction is preferably less than 2%, more preferably less than 1%.

Furthermore, Cu is precipitated in the bainite described above, the average particle size of the precipitate is 40 nm or less, and the number of Cu precipitate particles per unit area of the steel plate or steel pipe cross section is 1.0 × 10 ³ particles / μm ^2. That's it. Since Cu precipitates are finely dispersed to inhibit dislocation movement and increase the work hardening rate, it is extremely effective for increasing the strength of the bainite structure, and Cu itself has ductility, so the tensile strength increases. Despite this, the uniform elongation is improved.

  Therefore, in order to confirm the effect, the basic chemical composition of the steel is 0.11% C, 0.20% Si, 1.45% Mn, 0.02% Al, 0.05% Nb, 0.020 in mass%. Comparison of 15mm thick steel plates prepared by making small steel ingots with various changes in Cu from 1.5% to 5.0% as% Ti, and applying hot rolling and accelerated cooling to the microstructure of bainite On the other hand, a 15 mm thick steel plate was prepared, which was subjected to rapid rolling and accelerated cooling, followed by rapid heating and tempering that minimizes temper softening of bainite. In rapid heating and tempering, the heating temperature was variously changed to change the Cu precipitation state in bainite.

  From these steel sheets, a full thickness tensile test piece based on API-5L is taken in parallel with the rolling longitudinal direction, a tensile test is performed, and the obtained tensile strength and uniform elongation are subjected to rapid heat treatment for Cu precipitation. The amount of increase in tensile strength and uniform elongation due to Cu precipitation was calculated by comparing the case of performing the above and the case of accelerated cooling.

Next, three thin film samples for a transmission electron microscope each having a thickness of 0.3 mm were taken in parallel with the thickness direction from the rolling longitudinal cross section of the steel plate subjected to the rapid heat treatment for Cu precipitation, and transmitted. The Cu precipitates in the bainite lath were observed with an electron microscope at a magnification of 100,000. Photographs of Cu deposits taken for 3 fields of view for each thin film sample, a total of 9 fields, are image-analyzed, and the number of precipitates per unit observation area (1 μm ² ) and the average value of the diameters of each precipitate are measured with an image analyzer. Quantified. And the relationship with the above-mentioned amount of tensile strength increase and the amount of uniform elongation was derived from the number of Cu precipitates and the precipitate size in the bainite of the steel sheet.

As shown in FIG. 1, the amount of increase in tensile strength when Cu is precipitated is such that the tensile strength is 100 MPa or more when there are 1.0 × 10 ³ particles / μm ² or more of precipitates having an average particle size of 40 nm or less. Since an increase is obtained, the lower limit of the average particle diameter is set to 40 nm or less, and the number of particles is set to 1.0 × 10 ³ particles / μm ² or more. When the average particle diameter exceeds 40 nm, 1.0 × 10 ³ cells / [mu] m can not be obtained so much strength increased even number of ² or more particles, also the number of particles 1.0 × 10 ³ cells / [mu] m ² When the ratio is less than 1, the strength cannot be increased regardless of the average particle size.

  As for uniform elongation, as shown in FIG. 2, the improvement in uniform elongation depends on the average particle diameter rather than the number of particles, and when the average particle diameter is 10 nm or more, significant improvement is observed. The average particle size is preferably 10 nm or more.

  Next, the reason why the product of tensile strength (MPa) and uniform elongation (%) in the steel sheet rolling direction is 8500 or more will be described. In the conventional steel sheet strengthening method, the uniform elongation decreases conversely with the increase in tensile strength. In particular, when the strength is increased such that the tensile strength exceeds 1150 MPa, the uniform elongation is significantly reduced. As a result, after forming into a welded steel pipe, the uniform elongation in the longitudinal direction of the steel pipe also becomes a remarkably low value, and when some tensile deformation occurs after laying the pipe, local breakage easily occurs, which poses a safety problem.

  Therefore, the product of tensile strength (MPa) and uniform elongation (%) is selected as an index indicating the balance between tensile strength and uniform elongation, and the value is used as the highest strength grade X100 grade (tensile). 8500 or more, which exceeds the product of the tensile strength (MPa) and uniform elongation (%) of 8000 (strength of about 800 MPa) that is considered to have the highest ductility (uniform elongation of about 10%).

  The reason why the product of the tensile strength (MPa) in the longitudinal direction of the steel pipe and the uniform elongation (%) is 8500 or more is also as described above. In order to stably set the product of the tensile strength (MPa) and the uniform elongation (%) in the longitudinal direction of the steel pipe to 8500 or more, the tensile strength (MPa) and the uniform elongation (%) in the steel plate rolling direction The product of is desirably 8700 or more.

3. Manufacturing conditions Steel having the above-described composition is melted by a conventional method using a melting means such as a converter or an electric furnace, and a steel material such as a slab is formed by a conventional method such as a continuous casting method or an ingot-bundling method. It is preferable to do. Although the melting method and the casting method are not limited to the above-described methods, it is desirable to cast a steel piece by a steelmaking process by a converter method and a continuous casting process from the viewpoint of economy. Thereafter, the shape is rolled into a desired shape, and after rolling, cooling and reheating are performed. The production conditions of the present invention are shown below.

  The steel temperature condition defined in the present invention refers to the average temperature in the steel slab or steel plate thickness direction. The average temperature in the plate thickness direction is obtained by simulation calculation or the like from the plate thickness, surface temperature, cooling conditions, and the like. For example, the average temperature in the plate thickness direction can be obtained by calculating the temperature distribution in the plate thickness direction using the difference method.

(1) Heating temperature: 1100-1200 ° C
In performing hot rolling, Nb and Ti carbides need to be dissolved once in the slab heating stage in order to precipitate carbide in bainite, and the heating temperature needs to be 1100 ° C. or higher in order to sufficiently dissolve both. is there. On the other hand, when the steel slab is heated to a temperature exceeding 1200 ° C., even if pinning with TiN is performed, austenite grains grow significantly and the toughness deteriorates, so the heating temperature is in the range of 1100 to 1200 ° C.

(2) Cumulative rolling reduction at 950 ° C. or lower: 50% or higher In the present invention, the austenite non-recrystallized region during hot rolling is expanded to 950 ° C. by adding Nb. By performing rolling at a cumulative reduction ratio of 50% or more in this temperature range, austenite grains are expanded, and the bainite that is transformed by the subsequent accelerated cooling is refined and the toughness is improved. Therefore, the cumulative reduction ratio is 50 % Or more.

(3) Rolling end temperature: Ar ₃ transformation point or more When the hot rolling temperature is lower than the ferrite transformation start temperature, ferrite is generated during rolling and the strength is lowered. Therefore, the hot rolling end temperature is at least the Ar ₃ transformation point. That's it.

The Ar ₃ transformation point was determined by the following formula (1) calculated from the components.
Ar ₃ = 910-310C-80Mn-20Cu-15Cr-55Ni-80Mo (1)
Here, each component element means content (mass%).

(4) Cooling rate: 40-80 ° C / sec
In order to transform the bainite, accelerated cooling is subsequently performed after hot rolling. When the cooling rate is less than 40 ° C./sec, Cu partially precipitates during cooling, so that the basic idea of the present invention that precipitates Cu from a supersaturated solid solution state in bainite cannot be realized. In order to obtain a supersaturated solid solution state, the lower limit is made 40 ° C./sec. On the other hand, in the case of a cooling rate exceeding 80 ° C./sec, it is difficult to control to the cooling stop temperature described later, and martensitic transformation occurs in the vicinity of the surface and the toughness is significantly reduced. 80 ° C./sec. Preferably it is 50-65 degreeC / sec.

In addition, since accelerated cooling is preferably started from a state where the entire structure is austenite, the accelerated cooling start temperature is preferably equal to or higher than the Ar ₃ transformation point.

(5) Cooling stop temperature for accelerated cooling: 400 to 500 ° C
The cooling stop temperature is set to 400 to 500 [deg.] C. in order to mainly use the bainite as a microstructure that is transformed by accelerated cooling. When the cooling stop temperature is less than 400 ° C., martensitic transformation occurs and causes a decrease in toughness, so the lower limit is set to 400 ° C. On the other hand, when the cooling stop temperature exceeds 500 ° C., the accelerated cooling is stopped without completing the bainite transformation, and pearlite or the like is generated during the subsequent reheating, which causes a decrease in strength. The range is ˜500 ° C. Preferably, it is 420-480 degreeC.

(6) Reheating start temperature: (cooling stop temperature −50 ° C.) to cooling stop temperature Reheating is performed immediately after accelerated cooling in order to finely precipitate Cu, which is supersaturated in the steel, into the bainite structure. When the temperature at which reheating is started is lower than the accelerated cooling stop temperature by more than 50 ° C., it is difficult to finely disperse Cu, and since sufficient precipitation strengthening is not obtained and the target strength is not reached, the reheating starting temperature is , (Cooling stop temperature −50 ° C.) or higher. Preferably, it is (cooling stop temperature −30 ° C.) or higher.

  Here, reheating immediately after stopping accelerated cooling refers to performing the above-described reheating treatment within two minutes after stopping accelerated cooling.

(7) Temperature rise rate during reheating: 5 ° C./sec or more The temperature rise rate during reheating is 5 ° C./sec or more. When the heating rate at the time of reheating is slow, Cu begins to precipitate during the heating, and when the target heating temperature is reached, it becomes coarse and coarse, making it difficult to finely disperse Cu, and sufficient precipitation strengthening cannot be obtained. Since it is below the target strength, the rate of temperature increase during reheating is set to 5 ° C./sec or more.

(8) Reheating temperature: 550 to 650 ° C
The reheating temperature greatly affects the Cu precipitation state. That is, if the heating temperature is less than 550 ° C., sufficient Cu precipitation cannot be obtained. Conversely, if the heating temperature exceeds 650 ° C., the Cu precipitate becomes coarse due to overaging, and sufficient precipitation strengthening occurs. Since it is less than the target strength without being obtained, the reheating temperature is in the range of 550 to 650 ° C.

  The holding time for reheating is not particularly limited, but is preferably about 1 to 100 seconds.

  In the reheating, it is not necessary to set the heating and holding time, and the precipitated Cu may be coarsened if held for an excessive time in the reheating temperature range. The residence time in is preferably 200 seconds or less.

  Further, in the cooling process after reheating, the precipitated Cu is not rapidly coarsened. Therefore, the cooling conditions after reheating are not particularly defined, but basically it is preferably air cooling.

  As equipment for performing reheating after accelerated cooling, a heating device can be installed downstream of the cooling equipment for performing accelerated cooling. As the heating device, it is preferable to use a gas combustion furnace or induction heating device capable of rapid heating of the steel sheet.

  The method for forming the steel plate produced by the above method into a steel pipe is not particularly limited, and any of conventionally used forming methods such as UOE forming, press bend forming, and roll forming can be used. For example, the steel sheet is cold formed into a cylindrical shape so that the rolling direction of the steel sheet is the longitudinal direction of the steel pipe, and the opposite end faces are butt welded to produce the high ductility ultra high strength welded steel pipe of the present invention. can do. Moreover, the high ductility ultra high strength welded steel pipe of this invention which further improved the dimensional accuracy can be manufactured by giving a pipe expansion forming process or a diameter reduction forming process after welding.

  Using steel having the chemical composition shown in Table 1, steel plates having a thickness of 15 to 20 mm were produced under the conditions of hot rolling, accelerated cooling, and reheating shown in Table 2. And these steel plates were used as a raw material, and the welded steel pipe was manufactured by the UOE process so that the rolling longitudinal direction of the steel sheet became the longitudinal direction of the steel pipe. The welding of the steel pipe was a single-layer inner / outer surface welding by a submerged arc welding machine, and the welding heat input was 50 kJ / cm or less in any welding. Moreover, the low carbon type welding wire whose C content does not exceed 0.11% was used for the welding wire.

From each steel pipe, a full-thickness tensile specimen according to API-5L (the longitudinal direction of the pipe is the longitudinal direction of the specimen), a DWTT specimen, and a V-notch Charpy impact specimen described in JISZ2202 from the center of the pipe thickness. The steel pipe was subjected to a tensile test and a DWTT test (test temperature: 0 ° C.) and a Charpy impact test (test temperature: 0 ° C.) to evaluate strength, uniform elongation, and toughness. As toughness, the absorbed energy vE ₀ at 0 ° C. was evaluated for the Charpy impact test, and the ductile fracture surface ratio SA ₀ at 0 ° C. was evaluated for the DWTT test. Also, a sample for observing the metal structure (micro structure) is taken from the center position of the tube thickness, and the plate thickness cross section parallel to the longitudinal direction of the tube is mirror-polished, and then etched with a 3% nitric acid alcohol etchant, Were observed at a magnification of 400 times to confirm the type of microstructure.

Next, an ultrasonic flaw detection test was manually performed over the entire length of the welded portion of the steel pipe to investigate whether the weld metal was cracked. In addition, about the location where the crack generate | occur | produced, the welding part sample was cut and extract | collected and the crack state was confirmed by observing a cross section.
Table 3 summarizes the survey results of the microstructure, strength, uniform elongation, toughness, etc. of the base metal and the evaluation results of the nondestructive inspection of the pipe weld.

  Inventive example No. In Nos. 1 to 6, the composition, production conditions, and metal structure were within the scope of the invention, and good strength and toughness were obtained.

  No. 7 to 14 are comparative examples. In No. 7, the cooling stop temperature was less than 400 ° C., martensite crystallized, and the tensile strength × uniform elongation was insufficient. No. Since the cooling stop temperature of No. 8 exceeded 500 ° C., pearlite crystallized and the tensile strength was insufficient. As a result, tensile strength × uniform elongation was insufficient. No. No. 9 had a reheating temperature of less than 550 ° C. and Cu was insufficiently precipitated, so that the tensile strength was insufficient, and as a result, tensile strength × uniform elongation was insufficient.

  No. No. 10 had a reheating temperature exceeding 650 ° C., and Cu precipitates were coarsened, so that the tensile strength was insufficient, and as a result, tensile strength × uniform elongation was insufficient.

  No. 11, no. 12, no. 13, no. In No. 14, the component composition was out of the scope of the invention, and none of the predetermined characteristics was obtained. That is, no. No. 11 had insufficient base material strength because the amount of C was less than the range of the present invention. No. No. 12 was determined to be cracked in the ultrasonic flaw detection test of the pipe welded portion. When the welded portion sample was cut and collected from the cracked portion and observed for cross section, it was confirmed that the crack was hot. This is no. This is because the amount of C of 12 is larger than the range of the present invention, and C is diluted in the weld metal. No. In No. 13, since the amount of Cu was small and precipitation strengthening was insufficient, the base material strength was small. No. In No. 14, the amount of Nb was excessive, and the toughness decreased.

  In Example 2, a steel plate having a chemical composition shown in Table 1 of Example 1 was used to produce a steel plate having a thickness of 15 to 20 mm under the hot rolling, accelerated cooling, and reheating conditions shown in Table 4. And these steel plates were used as a raw material, and the welded steel pipe was manufactured by the UOE process so that the rolling longitudinal direction of the steel sheet became the longitudinal direction of the steel pipe. The welding of the steel pipe was a single-layer inner / outer surface welding by a submerged arc welding machine, and the welding heat input was 50 kJ / cm or less in any welding. Moreover, the low carbon type welding wire whose C content does not exceed 0.11% was used for the welding wire.

From each steel pipe, a full-thickness tensile specimen according to API-5L (the longitudinal direction of the pipe is the longitudinal direction of the specimen), a DWTT specimen, and a V-notch Charpy impact specimen described in JISZ2202 from the center of the pipe thickness. The steel pipe was subjected to a tensile test and a DWTT test (test temperature: 0 ° C.) and a Charpy impact test (test temperature: 0 ° C.) to evaluate strength, uniform elongation, and toughness. As toughness, the absorbed energy vE ₀ at 0 ° C. was evaluated for the Charpy impact test, and the ductile fracture surface ratio SA ₀ at 0 ° C. was evaluated for the DWTT test. Also, a sample for observing the metal structure (micro structure) is taken from the center position of the tube thickness, and the plate thickness cross section parallel to the longitudinal direction of the tube is mirror-polished, and then etched with a 3% nitric acid alcohol etchant, Were observed at a magnification of 400 times to confirm the type of microstructure.

  Further, three thin film samples for a transmission electron microscope were collected from the plate thickness section parallel to the longitudinal direction of the tube, and each of the three fields of view at a magnification of 100000 times with a transmission electron microscope, totaling 9 fields of Cu precipitates. A photograph was taken, and the average diameter of each Cu precipitate particle and the number of particles per unit area were calculated by image analysis.

Next, an ultrasonic flaw detection test was manually performed over the entire length of the welded portion of the steel pipe to investigate whether the weld metal was cracked. In addition, about the location where the crack generate | occur | produced, the welding part sample was cut and extract | collected, and the crack state was confirmed by observing a cross section.
Table 5 summarizes the survey results of the microstructure, strength, uniform elongation, toughness, etc. of the base metal and the evaluation results of the nondestructive inspection of the pipe weld.

  Inventive example No. In Nos. 15 to 17, the composition, production conditions, microstructure, number of Cu precipitate particles, and average particle diameter were within the range of the invention, and good strength and toughness were obtained.

No. Nos. 18 to 20 are comparative examples. In No. 18, the cooling rate during accelerated cooling was lower than 40 ° C./sec. Therefore, Cu that began to precipitate during cooling became coarse by subsequent reheating, and the average diameter of the precipitate particles exceeded 40 nm, so the tensile strength was high. As a result, tensile strength × uniform elongation was insufficient. No. No. 19 after the accelerated cooling stop, the reheating start is delayed, and the reheating temperature start temperature is lowered by 50 ° C. or more from the cooling stop temperature. Since it was below 1.0 × 10 ³ / μm ² , the tensile strength was insufficient, and as a result, the tensile strength × uniform elongation was insufficient. No. No. 20 has a slow temperature rise rate of 2 ° C./sec during reheating, Cu precipitated during heating is agglomerated and coarsened, the average diameter of the precipitate particles exceeds 40 nm, and precipitates per unit cross-sectional area are Since it was below 1.0 × 10 ³ / μm ² , the tensile strength was insufficient, and as a result, the tensile strength × uniform elongation was insufficient.

  Using steel having the chemical composition shown in Table 6, steel plates having a thickness of 15 to 18 mm were produced under the hot rolling, accelerated cooling, and reheating conditions shown in Table 7.

From each steel plate, a full-thickness tensile specimen according to API-5L (the steel sheet rolling direction is the longitudinal direction of the specimen), a DWTT specimen, and a V-notch Charpy impact specimen described in JISZ2202 from the center of the thickness. The steel sheet was collected and subjected to a tensile test and a DWTT test (test temperature: 0 ° C.) and a Charpy impact test (test temperature: 0 ° C.) to evaluate strength, uniform elongation, and toughness. As toughness, the absorbed energy vE ₀ at 0 ° C. was evaluated for the Charpy impact test, and the ductile fracture surface ratio SA ₀ at 0 ° C. was evaluated for the DWTT test.

  In addition, a sample for observing the metal structure (micro structure) was collected from the center position of the plate thickness, and the plate thickness cross section parallel to the rolling direction of the steel plate was mirror-polished and then etched with a 3% nitric acid alcohol corrosion solution. Were observed at a magnification of 400 times to confirm the type of microstructure.

  Further, three thin film samples for a transmission electron microscope were taken from the thickness cross section parallel to the rolling direction of the steel sheet, and each of the three fields of view at a magnification of 100000 times with a transmission electron microscope, a total of nine fields of Cu precipitates. A photograph was taken, and the average diameter of each Cu precipitate particle and the number of particles per unit area were calculated by image analysis.

  Table 8 summarizes the survey results of the microstructure, strength, uniform elongation, and toughness of the base material.

  Inventive example No. In 21 to 23, the component composition, production conditions and metal structure were within the scope of the invention, and good strength and toughness were obtained.

No. Nos. 24 and 25 are comparative examples. No. 24 was not subjected to reheating treatment after accelerated cooling, so Cu precipitation was insufficient, and the precipitate per unit cross-sectional area was less than 1.0 × 10 ³ pieces / μm ² , so that the tensile strength was insufficient, As a result, tensile strength × uniform elongation was insufficient. No. In No. 25, the amount of Cu was small, and the number of precipitates per unit cross-sectional area was less than 1.0 × 10 ³ pieces / μm ² , so that the tensile strength was insufficient, and as a result, tensile strength × uniform elongation was insufficient.

Claims (8)

  Steel composition is mass%, C: 0.09 to 0.11%, Si: 0.05 to 0.20%, Mn: 1.0 to 1.5%, Al: 0.01 to 0.08 %, Cu: 2.0 to 4.0%, Nb: 0.05 to 0.07%, Ti: 0.015 to 0.025%, and Cr: 0.05 to 0.6% , Mo: 0.05 to 0.6%, V: 0.01 to 0.1%, B: 0.0005 to 0.003%, one or more selected from the balance, the remainder Fe and inevitable impurities A steel sheet for high ductility ultra high strength welded steel pipe, characterized in that the metal structure is bainite and the product of tensile strength (MPa) in the rolling direction and uniform elongation (%) is 8500 or more. The average particle size of the Cu precipitates precipitated in the bainite of the metal structure is 40 nm or less, the number of Cu precipitate particles per unit area of the steel plate cross section is 1.0 × 10 ³ particles / μm ² or more, and the rolling direction The steel sheet for high ductility ultra-high-strength welded steel pipes according to claim 1, wherein the tensile strength of the steel sheet exceeds 1150 MPa.   Further, in terms of mass%, Ca: 0.0005 to 0.01%, REM: 0.0005 to 0.02%, Zr: 0.0005 to 0.03%, Mg: 0.0005 to 0.02% The steel sheet for high ductility ultra-high strength welded steel pipes according to claim 1, comprising at least one selected from the inside.   Furthermore, Ni: 1.0-4.0% is contained by the mass%, The steel plate for high ductility super high strength welded steel pipes in any one of Claim 1 thru | or 3 characterized by the above-mentioned.   It is manufactured using the steel plate according to any one of claims 1 to 4, and has a product of a pipe longitudinal tensile strength (MPa) and uniform elongation (%) of 8500 or more, and has a high ductility and a very high height. Strength welded steel pipe.   The high ductility ultra high strength welded steel pipe according to claim 5, wherein the pipe longitudinal tensile strength is over 1150 MPa. The steel slab having the steel composition according to any one of claims 1 to 4 is heated to 1100 to 1200 ° C, and then hot rolling is started, and the cumulative rolling reduction at 950 ° C or less is set to 50% or more, Ar ₃ Rolling is completed at the transformation point or higher, followed by accelerated cooling at a cooling rate of 40 to 80 ° C./sec and a cooling stop temperature of 400 to 500 ° C. A method for producing a steel sheet for a high ductility ultra-high strength welded steel pipe, characterized in that the steel sheet is air-cooled after rapidly reheating to 550 to 650 ° C at a temperature rising rate of 5 ° C / sec or more from a temperature range.   A steel sheet according to any one of claims 1 to 4, wherein the steel sheet is formed into a tubular shape in a cold state, the end faces facing each other are butt welded, and then expanded or reduced in diameter. Production method.

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