JP2010196163A - Thick, high-tension, hot-rolled steel sheet excellent in low temperature toughness, and manufacturing method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for manufacturing a thick, high-tension, hot-rolled steel sheet, which has high strength of ≥510 MPa TS and high ductility in combination and is well balanced between the strength and the durability and further, exhibits excellent low-temperature toughness. <P>SOLUTION: The steel material having a composition containing 0.02-0.08% C, 0.01-0.10% Nb, 0.001-0.05% Ti, so as to satisfy äTi+(Nb/2)}/C<4, is heated and subjected to the hot-rolling, and then subjected to a primary and a secondary accelerative coolings at ≥10°C/s cooling speed in the center part of the sheet thickness, and then, winding is performed at a prescribed winding temperature. In the primary accelerative cooling, the cooling is performed up to a primary cooling-stop temperature of ≥500°C so that the cooling speed difference between the surface layer and the center part of the sheet thickness becomes <80°C/s. In the secondary accelerative cooling, the cooling wherein the cooling speed difference between the surface layer and the center part of the sheet thickness becomes ≥80°C/s, is performed up to a specific cooling-stop temperature or lower depending on the alloy element contents and the cooling speed. Further, the winding temperature is set to a specific temperature or lower, depending on the alloy element contents. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention is a thick-walled, high-tensile hot-rolled steel sheet suitable for use as a material for high-strength ERW steel pipes or high-strength spiral steel pipes that require high toughness for line pipes that transport crude oil, natural gas, and the like, and production thereof In particular, it relates to the improvement of low temperature toughness.

In recent years, oil and natural gas mining and pipeline construction have been actively carried out in extremely cold regions such as the North Sea, Canada and Alaska due to soaring crude oil since the oil crisis and the demand for diversified energy supply sources. It has come to be. Also, once the development has been abandoned, the development of a corrosive sour gas field, etc., has become active.
Furthermore, in the pipeline, in order to improve the transportation efficiency of natural gas and oil, there is a tendency to perform high-pressure operation with a large diameter. In order to withstand the high-pressure operation of the pipeline, the transport pipe (line pipe) needs to be a thick steel pipe, and a UOE steel pipe made of a thick steel plate has been used. However, recently, due to the strong demand for further reduction of pipeline construction costs and the lack of supply capacity of UOE steel pipes, there is a strong demand for reducing the material cost of steel pipes. Instead of UOE steel pipes, high-strength ERW steel pipes or high-strength spiral steel pipes made of coil-shaped hot-rolled steel sheets (hot-rolled steel strips), which are more productive and cheaper, have come to be used.

  These high-strength steel pipes are required to maintain excellent low-temperature toughness from the viewpoint of preventing line pipe breakage. In order to produce a steel pipe having both such high strength and high toughness, in steel sheets that are steel pipe materials, transformation strengthening using accelerated cooling after hot rolling and alloying elements such as Nb, V, Ti, etc. Strengthening by precipitation strengthening using precipitates and toughness by microstructure refinement using controlled rolling have been attempted.

In addition, in line pipes used for transporting crude oil and natural gas containing hydrogen sulfide, in addition to characteristics such as high strength and high toughness, so-called hydrogen-induced crack resistance (HIC resistance), stress corrosion crack resistance, and so on It is required to have excellent sour resistance.
In response to such a request, for example, Patent Document 1 includes C: 0.005 to less than 0.030%, B: 0.0002 to 0.0100%, Ti: 0.20% or less, and Nb: 0.25% or less. Two kinds are included so as to satisfy (Ti + Nb / 2) / C: 4 or more, and further steel containing an appropriate amount of Si, Mn, P, S, Al, N is hot-rolled, and then 5 to 20 ° C./s. Toughness, cooled at a cooling rate of 550 ° C and wound up in a temperature range of more than 550 ° C to 700 ° C, the structure is composed of ferrite and / or bainitic ferrite, and the amount of solid solution C in the grain is 1.0 to 4.0 ppm A method for producing a high-strength hot-rolled steel sheet having a low yield ratio and an excellent strength has been proposed. With the technique described in Patent Document 1, a high-strength hot-rolled steel sheet having excellent toughness, weldability, and sour resistance and having a low yield ratio is obtained without causing material unevenness in the thickness direction and the length direction. You can do that. However, in the technique described in Patent Document 1, since the amount of solid solution C in the grains is 1.0 to 4.0 ppm, crystal grain growth is likely to occur due to heat input during circumferential welding, and the weld heat affected zone is coarse. There exists a problem that it becomes a grain and the toughness fall of the welding heat affected zone of a circumferential welded part tends to occur.

In Patent Document 2, C: 0.01 to 0.12%, Si: 0.5% or less, Mn: 0.5 to 1.8%, Ti: 0.010 to 0.030%, Nb: 0.01 to 0.05%, Ca: 0.0005 to 0.0050%, carbon equivalent: 0.40, Ca / O: 1.5 to 2.0 so as to satisfy, the steel slab comprising, exit hot rolled at Ar _{3 +} 100 ° C. or higher, 20 seconds After cooling, the above Ar ₃ point There has been proposed a method for producing a high-strength steel sheet excellent in hydrogen-induced cracking resistance, which is cooled from temperature, cooled to 550 to 650 ° C. within 20 seconds, and then wound up at 450 to 500 ° C. According to the technology described in Patent Document 2, API standard X60 to X70 grade steel plates for line pipe having hydrogen-induced crack resistance can be manufactured. However, with the technique described in Patent Document 2, it is impossible to secure a desired cooling time with a thick steel plate, and there is a problem that further improvement of the cooling capacity is required to secure desired characteristics. there were.

Moreover, although it is a thick steel plate, in patent document 3, C: 0.03-0.06%, Si: 0.01-0.5%, Mn: 0.8-1.5%, S: 0.0015% or less, Al: 0.08% or less, Ca: 0.001 The steel containing up to 0.005%, O: 0.0030% or less, and containing Ca, S, O so as to satisfy a specific relationship is heated to a cooling rate of 5 ° C./s or more from the temperature above the Ar ₃ transformation point. Accelerated cooling to 400 to 600 ° C, and then immediately reheat to a steel plate surface temperature of 600 ° C or higher and a plate thickness center temperature of 550 to 700 ° C at a heating rate of 0.5 ° C / s or higher. There has been proposed a method for producing a steel sheet for high-strength line pipe excellent in hydrogen-induced cracking resistance, in which the temperature difference between the center of the plate thickness is 20 ° C. or more. In the technique described in Patent Document 3, a steel sheet is obtained in which the fraction of the second phase in the metal structure is 3% or less, and the hardness difference between the surface layer and the thickness center is within 40 points in terms of Vickers hardness. The thick steel plate is excellent in hydrogen-induced crack resistance. However, the technique described in Patent Document 3 has a problem that a reheating process is required, the manufacturing process becomes complicated, and further arrangement of a reheating facility or the like is required.

Moreover, although it is a thick steel plate, in patent document 4, C: 0.01-0.3%, Si: 0.6% or less, Mn: 0.2-2.0%, P, S, Al: 0.06% or less, Ti: 0.005-0.035% , N: at 0.001 to 0.006% temperature slab the cooling process after the hot rolling Ac ₁ -50 ° C. or less of including, performs rolling over 2% cumulative, then, Ac ₁ super Ac ₃ less than A method for producing a steel material having a coarse ferrite layer on the front and back surfaces, which is heated to a temperature and allowed to cool, has been proposed. In the technique described in Patent Document 4, it is said that it contributes to the improvement of SCC sensitivity, weather resistance, and corrosion resistance of steel materials, and further suppression of material deterioration after cold working. However, the technique described in Patent Document 4 has a problem that a reheating process is required, the manufacturing process becomes complicated, and further arrangement of reheating equipment and the like is required.

Furthermore, recently, steel pipes for extremely cold regions are often required to have excellent fracture toughness, particularly CTOD characteristics and DWTT characteristics, from the viewpoint of preventing burst fracture of pipelines.
In response to such a requirement, for example, Patent Document 5 contains appropriate amounts of C, Si, Mn, and N, and further contains Si and Mn in a range where Mn / Si satisfies 5 to 8, and further includes Nb. : Rolling ratio of the first rolling performed at 1100 ° C or higher after heating the steel slab containing 0.01 to 0.1%: 15-30%, Total rolling ratio at 1000 ° C or higher: 60% or higher, Rolling ratio of final rolling : After rough rolling under the condition of 15-30%, once the surface layer is cooled to a temperature of ₁ point or less at a cooling rate of 5 ° C / s or more, the surface layer is reheated or forced overheated. Finishing rolling is started when the temperature reaches (Ac ₃ −40 ° C.) to (Ac ₃ + 40 ° C.), and the total rolling reduction at 950 ° C. or less is 60% or more, and the rolling end temperature is Ar ₃ points or more. Finish the finish rolling, start cooling within 2 s after finishing the finish rolling, cool to 600 ° C. or less at a rate of 10 ° C./s or more, and wind up in the temperature range of 600 to 350 ° C. Method of manufacturing an electric resistance welded steel pipe for hot rolled steel sheet are described. The steel sheet manufactured by the technique described in Patent Document 5 has a refined structure of the steel sheet surface layer without adding an expensive alloy element or heat-treating the entire steel pipe, resulting in low temperature toughness, particularly DWTT characteristics. An excellent high-strength ERW steel pipe can be manufactured. However, the technique described in Patent Document 5 has a problem that a steel plate with a large thickness cannot secure a desired cooling rate, and further cooling capacity needs to be improved in order to secure desired characteristics. there were.

Patent Document 6 contains appropriate amounts of C, Si, Mn, Al, and N, and further includes Nb: 0.001 to 0.1%, V: 0.001 to 0.1%, Ti: 0.001 to 0.1%, Cu, Ni , contain one or two or more of Mo, the steel slab Pcm value is 0.17 or less, after heating, the surface temperature is terminated finish rolling at (Ar ₃ -50 ℃) above conditions, A method for producing a hot-rolled steel strip for a high-strength ERW pipe excellent in low-temperature toughness and weldability that is cooled immediately after rolling and wound up and cooled at a temperature of 700 ° C. or lower is described.

Japanese Unexamined Patent Publication No. 08-319538 Japanese Unexamined Patent Publication No. 09-296216 JP 2008-056962 JP Japanese Patent Laid-Open No. 2001-240936 Japanese Patent Laid-Open No. 2001-207220 JP 2004-315957 A

Recently, however, steel sheets for high-strength ERW steel pipes are required to further improve low-temperature toughness, particularly CTOD characteristics and DWTT characteristics. The technique described in Patent Document 6 has a problem that the low-temperature toughness is not sufficient, and the excellent low-temperature toughness cannot be provided to the extent that the required CTOD characteristics and DWTT characteristics are sufficiently satisfied.
The present invention solves the above-mentioned problems of the prior art, and does not require a large amount of alloying element addition, has both high strength and excellent ductility, excellent strength and ductility balance, and excellent low temperature An object of the present invention is to provide a thick, high-tensile hot-rolled steel sheet suitable for high-strength ERW steel pipe or high-strength spiral steel pipe having toughness, particularly excellent CTOD characteristics, and DWTT characteristics, and a method for producing the same. .

As used herein, “high-tensile hot-rolled steel sheet” refers to a hot-rolled steel sheet having a high strength of tensile strength TS: 510 MPa or more, and “thick-walled” steel sheet refers to a steel sheet having a thickness of 11 mm or more. It shall be said.
In addition, “excellent CTOD characteristics” here refers to the case where the critical opening displacement CTOD value in a CTOD test conducted at a test temperature of −10 ° C. is 0.30 mm or more in accordance with the provisions of ASTM E 1290. It shall be said.

The “excellent DWTT property” here is a DWTT test conducted in accordance with the provisions of ASTM E 436, and the minimum temperature (DWTT temperature) at which the ductile fracture surface ratio is 85% is −35 ° C. or less. This shall be the case.
Further, “excellent in strength / ductility balance” here means a case where TS × El is 18000 MPa% or more. The elongation El (%) is a value obtained by testing using a plate-like test piece (parallel part width: 12.5 mm, distance between gauge points: 50 mm) in accordance with ASTM E8.

  In order to achieve the above-described object, the present inventors have intensively studied various factors affecting low temperature toughness, particularly DWTT characteristics and CTOD characteristics. As a result, it came to mind that the DWTT characteristic and the CTOD characteristic, which are toughness tests at the full thickness, are greatly influenced by the structure uniformity in the thickness direction. And it discovered that the influence of the structure nonuniformity of the sheet thickness direction on the DWTT characteristic and the CTOD characteristic, which are toughness tests at the full thickness, was manifested by a thick material having a sheet thickness of 11 mm or more.

  Further, according to further studies by the present inventors, “excellent DWTT characteristics” and “excellent CTOD characteristics” are the average grain size of ferrite at a position 1 mm (surface layer portion) in the plate thickness direction from the surface. Difference from the average crystal grain size of ferrite at the plate thickness center position (plate thickness center portion), ΔD is 2 μm or less, and the fraction of the second phase at the position (surface layer portion) 1 mm from the surface in the plate thickness direction (surface layer portion) It was found that the difference between the volume fraction) and the second phase structure fraction (volume fraction) at the plate thickness center position (plate thickness center portion), that is, ΔV, is 2% or less, can be secured.

First, the experimental results on which the present invention is based will be described.
A slab composed of 0.037% C-0.20% Si-1.59% Mn-0.016% P-0.0023% S-0.041% Al-0.061% Nb-0.013% Ti-balance Fe in mass% was used as a steel material. Note that (Ti + Nb / 2) / C is 1.18.
The steel material having the above composition is heated to 1230 ° C and hot rolled to a finish rolling start temperature of 980 ° C and a finish rolling finish temperature of 800 ° C to obtain a hot rolled sheet having a thickness of 12.7 mm. After rolling, cooling at a cooling rate of 18 ° C / s in the temperature region where the temperature at the center of the plate thickness is 750 ° C or lower is applied to various cooling stop temperatures, and then accelerated cooling is performed at various winding temperatures. Winding and hot rolled steel sheet (steel strip) were used.

  Test pieces were collected from the obtained hot-rolled steel sheet, and DWTT characteristics and structure were investigated. The structure is 1 mm from the surface in the plate thickness direction (surface layer part), the plate thickness center position (plate thickness center part), the average crystal grain size of ferrite (μm), the second phase structure fraction (volume%) Asked. From the measured values, the average crystal grain size difference ΔD of ferrite and the structure fraction of the second phase at a position (surface layer part) 1 mm from the surface in the sheet thickness direction and the sheet thickness center position (sheet thickness center part). The difference ΔV was calculated respectively. Here, “ferrite” refers to bainitic ferrite, bainite, or a mixed phase thereof. The second phase is either pearlite, martensite, MA (also called island martensite), upper bainite, or a mixed phase composed of two or more of these.

The obtained results are shown in FIG. 1 as the relationship between ΔD and ΔV exerted on DWTT.
From FIG. 1, it was found that “excellent DWTT characteristics” in which DWTT is −35 ° C. or less can be reliably maintained when ΔD is 2 μm or less and ΔV is 2% or less.
Next, FIG. 2 shows the relationship between ΔD and ΔV and the cooling stop temperature, and FIG. 3 shows the relationship between ΔD and ΔV and the coiling temperature.

2 and 3, it is necessary to adjust the cooling stop temperature to 620 ° C. or lower and the coiling temperature to 647 ° C. or lower in order to make ΔD 2 μm or less and ΔV 2% or less. I understand that.
According to further studies by the present inventors, the cooling stop temperature and the coiling temperature necessary for ΔD to be 2 μm or less and ΔV to be 2% or less include the inclusion of alloy elements mainly affecting the bainite transformation start temperature. It was found that it was determined depending on the amount and the cooling rate from the end of hot rolling. That is, in order to set ΔD to 2 μm or less and ΔV to 2% or less, the cooling stop temperature is the temperature at the center position of the plate thickness of the steel plate,
BFS (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni−1.5CR
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%), CR: cooling rate (° C./s))
The temperature is equal to or lower than the BFS defined in, and the coiling temperature is the temperature at the center of the plate thickness of the steel sheet.
BFS0 (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%))
It is important to set the temperature below BFS0 as defined in.

  Next, the inventors further examined the influence of cooling conditions on the improvement of ductility. The result is shown in FIG. Figure 4 shows cooling in a temperature range of 500 ° C or higher, changing the difference in average cooling rate between the surface layer and the center of the plate thickness, and cooling in the temperature range of less than 500 ° C. The balance between strength and ductility was investigated by increasing the water density at the time of primary cooling so that the difference in the average cooling rate of the part became 80 ° C / s or more, and further changing the cooling stop temperature and coiling temperature. It is. As shown in FIG. 4, when cooling after hot rolling, cooling is performed so that the difference in average cooling rate between the surface layer and the central part of the plate thickness is within a specific range (less than 80 ° C./s) in the temperature range up to 500 ° C. It was found that by adjusting the conditions, ductility was remarkably improved in addition to low-temperature toughness, and the strength / ductility balance TS × El was stabilized to 18000 MPa% or more. FIG. 4 shows that when the difference between the cooling stop temperature and the coiling temperature is less than 300 ° C., the strength / ductility balance TS × El is further stabilized to 18000 MPa% or more.

The present invention has been completed based on the above findings and further studies. That is, the gist of the present invention is as follows.
(1) By mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.50%, Mn: 0.5 to 1.8%, P: 0.025% or less, S: 0.005% or less, Al: 0.005 to 0.10%, Nb: 0.01 ˜0.10%, Ti: 0.001˜0.05%, and C, Ti, Nb is expressed by the following formula (1) (Ti + (Nb / 2)) / C <4 (1)
(Here, Ti, Nb, C: content of each element (mass%))
The steel material having the composition comprising the balance Fe and inevitable impurities is heated, hot rolling consisting of rough rolling and finish rolling is performed, and then accelerated cooling is performed after the hot rolling is completed. In forming a rolled steel sheet, the accelerated cooling is made of primary accelerated cooling and secondary accelerated cooling, and the primary accelerated cooling has an average cooling rate of 10 ° C./s or more at the thickness center position and the thickness center. Cooling with a difference between the average cooling rate of the position and the average cooling rate at the position of 1 mm from the surface in the thickness direction is less than 80 ° C / s, and the temperature at the position of 1 mm from the surface in the thickness direction. Cooling performed to a primary cooling stop temperature that is a temperature range of 650 ° C. or lower and 500 ° C. or higher, and the secondary accelerated cooling is performed at an average cooling rate of 10 ° C./s or higher at the thickness center position, Average cooling rate and average cooling rate at the position of 1mm from the surface in the plate thickness direction Cooling speed difference, the cooling is 80 ° C. / s or higher, the temperature of the plate thickness center position following (2)
BFS (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni−1.5CR (2)
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%), CR: cooling rate (° C./s))
The cooling is performed to the secondary cooling stop temperature below BFS defined in Fig. 3. After the secondary accelerated cooling, the following equation (3)
BFS0 (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni (3)
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%))
A method for producing a thick, high-tensile hot-rolled steel sheet having an excellent balance between strength and ductility, characterized by winding at a winding temperature of BFS0 or less as defined in 1.

(2) In (1), a method for producing a thick, high-tensile hot-rolled steel sheet, wherein air cooling is performed for 10 seconds or less between the primary accelerated cooling and the secondary accelerated cooling.
(3) The thick wall characterized in that, in (1) or (2), the accelerated cooling is 10 ° C./s or more at an average cooling rate in a temperature range of 750 to 650 ° C. at a plate thickness center position. Manufacturing method of high-tensile hot-rolled steel sheet.

(4) In any one of (1) to (3), the difference between the cooling stop temperature at a position of 1 mm in the plate thickness direction from the surface and the coiling temperature in the secondary accelerated cooling is within 300 ° C. A method for producing a thick, high-tensile hot-rolled steel sheet.
(5) In any one of (1) to (4), in addition to the above composition, in addition to mass, V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: 0.01 A method for producing a thick, high-tensile hot-rolled steel sheet, characterized in that the composition contains one or more of ˜0.50% and Ni: 0.01 to 0.50%.

(6) In any one of (1) to (5), in addition to the above-described composition, the thick high-tensile hot-rolled steel sheet further comprises a composition containing Ca: 0.0005 to 0.005% by mass%. Manufacturing method.
(7) By mass%, C: 0.02 to 0.08%, Si: 0.01 to 0.50%, Mn: 0.5 to 1.8%, P: 0.025% or less, S: 0.005% or less, Al: 0.005 to 0.10%, Nb: 0.01 ˜0.10%, Ti: 0.001˜0.05%, and C, Ti, Nb is expressed by the following formula (1) (Ti + (Nb / 2)) / C <4 (1)
(Here, Ti, Nb, C: content of each element (mass%))
The composition consisting of the balance Fe and unavoidable impurities and the structure 1mm from the surface in the plate thickness direction is the structure having the ferrite phase as the main phase, and the position 1mm from the surface in the plate thickness direction. The difference ΔD between the average crystal grain size of the ferrite phase and the average crystal grain size of the ferrite phase at the center position of the plate thickness is 2 μm or less, and the fraction of the second phase at the position of 1 mm from the surface in the plate thickness direction (volume% ) And the second phase structure fraction (volume%) at the center position of the sheet thickness has a structure in which the difference ΔV is 2% or less. .

(8) In (7), in addition to the above composition, in terms of mass%, V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: 0.01 to 0.50%, Ni: 0.01 A thick-walled, high-tensile hot-rolled steel sheet characterized by comprising one or more of ˜0.50%.
(9) A thick-walled, high-tensile hot-rolled steel sheet according to (7) or (8), characterized in that, in addition to the above composition, the composition further contains Ca: 0.0005 to 0.005% by mass%.

  According to the present invention, it is possible to easily and inexpensively produce a thick high-tensile hot-rolled steel sheet with less structural fluctuation in the thickness direction, excellent balance between strength and ductility, and excellent low-temperature toughness, especially DWTT and CTOD characteristics. It has a remarkable industrial effect. In addition, according to the present invention, it is possible to easily manufacture an ERW steel pipe for a line pipe and a spiral steel pipe for a line pipe, which have an excellent balance between strength and ductility, low temperature toughness, and excellent circumferential weldability when laying a pipeline. There is also an effect.

It is a graph which shows the relationship between DWTT, (DELTA) D, and (DELTA) V. It is a graph which shows the relationship between (DELTA) D and (DELTA) V and the cooling stop temperature of accelerated cooling. It is a graph which shows the relationship between (DELTA) D and (DELTA) V and coiling temperature. It is a graph which shows the relationship between the strength / ductility balance TS × El and the difference (cooling rate difference) between the cooling rate at the position of 1 mm from the surface in the plate thickness direction and the cooling rate at the plate thickness central position.

The manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the method for producing a hot-rolled steel sheet of the present invention, a steel material having a predetermined composition is heated, and hot rolling comprising rough rolling and finish rolling is performed to obtain a hot-rolled steel sheet.
First, the reasons for limiting the composition of the steel material used in the present invention will be described. Unless otherwise specified, mass% is simply expressed as%.

C: 0.02 to 0.08%
C is an element having an action of increasing the strength of steel, and in the present invention, it is necessary to contain 0.02% or more in order to ensure a desired high strength. On the other hand, an excessive content exceeding 0.08% increases the structural fraction of the second phase such as pearlite, and lowers the base metal toughness and the weld heat affected zone toughness. For this reason, C was limited to the range of 0.02 to 0.08%. In addition, Preferably it is 0.02 to 0.05%.

Si: 0.01-0.50%
Si has an action of increasing the strength of steel through solid solution strengthening and improvement of hardenability. Such an effect is recognized when the content is 0.01% or more. On the other hand, Si has the effect of concentrating C in the γ phase during the γ → α transformation and promoting the formation of the martensite phase as the second phase, resulting in a decrease in the toughness of the steel sheet. Moreover, Si forms an oxide containing Si during ERW welding, lowers the weld zone quality, and lowers the weld heat affected zone toughness. From this point of view, it is desirable to reduce Si as much as possible, but up to 0.50% is acceptable. For these reasons, Si was limited to 0.01 to 0.50%. Preferably it is 0.40% or less.

In addition, since hot rolled steel sheets for electric resistance welded steel pipes contain Mn, Si forms a low melting point Mn silicate and facilitates oxide discharge from the weld zone, so Si is contained in an amount of 0.10 to 0.30%. May be.
Mn: 0.5-1.8%
Mn has the effect of improving hardenability, and increases the strength of the steel sheet through the improvement of hardenability. Further, Mn forms MnS and fixes S, thereby preventing segregation of S grain boundaries and suppressing slab (steel material) cracking. In order to acquire such an effect, 0.5% or more of content is required. On the other hand, if the content exceeds 1.8%, solidification segregation during slab casting is promoted, Mn-concentrated portions remain in the steel sheet, and the occurrence of separation increases. In order to eliminate this Mn enriched part, it is necessary to heat to a temperature exceeding 1300 ° C., and it is not practical to carry out such a heat treatment on an industrial scale. For this reason, Mn was limited to the range of 0.5 to 1.8%. In addition, Preferably it is 0.9 to 1.7%.

P: 0.025% or less P is inevitably contained as an impurity in steel, but has an effect of increasing the strength of steel. However, when it exceeds 0.025% and it contains excessively, weldability will fall. For this reason, P was limited to 0.025% or less. In addition, Preferably it is 0.015% or less.
S: 0.005% or less S is inevitably contained as an impurity in steel like P, but if it exceeds 0.005% and excessively contained, slab cracking occurs and coarse MnS is contained in the hot-rolled steel sheet. Forming and causing a reduction in ductility. For this reason, S was limited to 0.005% or less. In addition, Preferably it is 0.004% or less.

Al: 0.005-0.10%
Al is an element that acts as a deoxidizer, and in order to obtain such an effect, it is desirable to contain 0.005% or more. On the other hand, the content exceeding 0.10% significantly impairs the cleanliness of the welded part during ERW welding. For this reason, Al was limited to 0.005 to 0.10%. In addition, Preferably it is 0.08% or less.

Nb: 0.01-0.10%
Nb is an element that has the effect of suppressing the coarsening and recrystallization of austenite grains, enabling the austenite non-recrystallization temperature range rolling in hot finish rolling, and by precipitating finely as carbonitride, It has the effect | action which makes a hot-rolled steel plate high intensity | strength with little content, without impairing property. In order to acquire such an effect, 0.01% or more of content is required. On the other hand, an excessive content exceeding 0.10% may cause an increase in rolling load during hot finish rolling, which may make hot rolling difficult. For this reason, Nb was limited to the range of 0.01 to 0.10%. In addition, Preferably it is 0.03-0.09%.

Ti: 0.001 to 0.05%
Ti has the effect of forming nitrides and fixing N to prevent cracking of slabs (steel material), and finely precipitates as carbides, thereby increasing the strength of the steel sheet. Such an effect becomes remarkable when the content is 0.001% or more. However, when the content exceeds 0.05%, the yield point is remarkably increased by precipitation strengthening. For this reason, Ti was limited to the range of 0.001 to 0.05%. In addition, Preferably it is 0.005-0.035%.

In the present invention, Nb, Ti, and C in the above ranges are included, and the following formula (1) (Ti + (Nb / 2)) / C <4 (1)
Nb, Ti, and C content are adjusted so as to satisfy the above.
Nb and Ti are elements that have a strong tendency to form carbides. When the C content is low, most of the C becomes carbides, and the amount of solid solution C in the ferrite grains is assumed to decrease drastically. The drastic decrease in the amount of C dissolved in ferrite grains adversely affects the circumferential weldability during pipeline construction. When circumferential welding is performed using a steel pipe manufactured using a steel plate with extremely reduced solid solution C in the ferrite grains as a line pipe, grain growth in the heat-affected zone of the circumferential weld becomes significant. The heat affected zone toughness of the circumferential weld may be reduced. For this reason, in this invention, Nb, Ti, and C are adjusted and contained so that Formula (1) may be satisfied. Thereby, it becomes possible to make solid solution C amount in a ferrite grain 10 ppm or more, and can prevent the fall of the heat affected zone toughness of a circumference welded part.

In the present invention, the above components are basic components. In addition to this basic composition, V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: One or more of 0.01 to 0.50%, Ni: 0.01 to 0.50%, and / or Ca: 0.0005 to 0.005% can be selected and contained as necessary.
One or more of V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: 0.01 to 0.50%, Ni: 0.01 to 0.50% V, Mo, Cr, Cu , Ni is an element that improves the hardenability and increases the strength of the steel sheet, and can be selected from one or more as required.

V is an element that has an effect of improving hardenability and forming carbonitride to increase the strength of the steel sheet, and such an effect becomes remarkable when the content is 0.01% or more. On the other hand, excessive content exceeding 0.10% deteriorates weldability. For this reason, V is preferably 0.01 to 0.10%. Further, it is more preferably 0.03 to 0.08%.
Mo is an element that has an effect of improving hardenability and forming carbonitride to increase the strength of the steel sheet. Such an effect becomes remarkable when the content is 0.01% or more. On the other hand, a large content exceeding 0.50% reduces weldability. For this reason, it is preferable to limit Mo to 0.01 to 0.50%. In addition, More preferably, it is 0.05 to 0.30%.

Cr is an element that has the effect of improving hardenability and increasing the strength of the steel sheet. Such an effect becomes remarkable when the content is 0.01% or more. On the other hand, an excessive content exceeding 1.0% tends to cause frequent welding defects during ERW welding. For this reason, it is preferable to limit Cr to 0.01 to 1.0%. In addition, More preferably, it is 0.01 to 0.80%.
Cu is an element that has the effect of improving the hardenability and increasing the strength of the steel sheet by solid solution strengthening or precipitation strengthening. In order to acquire such an effect, it is desirable to contain 0.01% or more, but inclusion exceeding 0.50% reduces hot workability. For this reason, it is preferable to limit Cu to 0.01 to 0.50%. In addition, More preferably, it is 0.10 to 0.40%.

  Ni is an element that has the effect of improving hardenability, increasing the strength of the steel, and improving the toughness of the steel sheet. In order to acquire such an effect, it is desirable to contain 0.01% or more. On the other hand, if the content exceeds 0.50%, the effect is saturated and an effect commensurate with the content cannot be expected, which is economically disadvantageous. For this reason, it is preferable to limit Ni to 0.01 to 0.50%. In addition, More preferably, it is 0.10 to 0.40%.

Ca: 0.0005 to 0.005%
Ca has the action of fixing S as CaS, spheroidizing sulfide inclusions, and controlling the form of inclusions, reducing the lattice strain of the matrix surrounding inclusions, and reducing the hydrogen trapping ability It is an element which has the effect | action to make. In order to acquire such an effect, it is desirable to make it contain 0.0005% or more, but if it contains more than 0.005%, CaO will increase and corrosion resistance and toughness will be reduced. For this reason, when it contains Ca, it is preferable to limit to 0.0005 to 0.005%. In addition, More preferably, it is 0.0009 to 0.003%.

The balance other than the components described above consists of Fe and inevitable impurities. Inevitable impurities include N: 0.005% or less, O: 0.005% or less, Mg: 0.003% or less, and Sn: 0.005% or less.
N: 0.005% or less N is inevitably contained in steel, but excessive inclusion frequently causes cracking during casting of a steel material (slab). For this reason, it is desirable to limit N to 0.005% or less. More preferably, it is 0.004% or less.

O: 0.005% or less O exists as various oxides in steel, and causes hot workability, corrosion resistance, toughness and the like to decrease. For this reason, it is desirable to reduce as much as possible in the present invention, but it is acceptable up to 0.005%. Since extreme reduction leads to an increase in refining costs, it is desirable to limit O to 0.005% or less.

Mg: 0.003% or less
Mg, like Ca, forms oxides and sulfides and has the effect of suppressing the formation of coarse MnS, but if it exceeds 0.003%, Mg oxide and Mg sulfide clusters occur frequently, and toughness Cause a decline. For this reason, it is desirable to limit Mg to 0.003% or less.
Sn: 0.005% or less
Sn is mixed from scraps used as steelmaking raw materials. Sn is an element that easily segregates at grain boundaries and the like, and if it is contained in a large amount exceeding 0.005%, the grain boundary strength is lowered and the toughness is lowered. For this reason, it is desirable to limit Sn to 0.005% or less.

As a manufacturing method of the steel material, it is preferable to melt the molten steel having the above composition by a conventional melting method such as a converter, and to make a steel material such as a slab by a conventional casting method such as a continuous casting method, The present invention is not limited to this.
The steel material having the above composition is heated and hot-rolled. Hot rolling consists of rough rolling using a steel material as a sheet bar and finish rolling using the sheet bar as a hot-rolled sheet.

  The heating temperature of the steel material is not particularly limited as long as it can be rolled into a hot-rolled sheet, but it is preferably a temperature in the range of 1100 to 1300 ° C. When the heating temperature is less than 1100 ° C., the deformation resistance is high, the rolling load increases, and the load on the rolling mill becomes excessive. On the other hand, when the heating temperature is higher than 1300 ° C., the crystal grains are coarsened and the low-temperature toughness is reduced, the amount of scale generation is increased, and the yield is lowered. For this reason, it is preferable that the heating temperature in hot rolling shall be 1100-1300 degreeC.

The heated steel material is roughly rolled into a sheet bar. The rough rolling conditions are not particularly limited as long as a sheet bar having a desired size and shape can be obtained. From the viewpoint of securing toughness, the rolling end temperature of rough rolling is preferably 1050 ° C. or lower.
The obtained sheet bar is further subjected to finish rolling. In addition, it is preferable to adjust the finish rolling start temperature by performing accelerated cooling on the sheet bar before finish rolling or by performing oscillation on the table. Thereby, the reduction rate in the temperature range effective for high toughness in the finish rolling mill can be increased.

In finish rolling, it is preferable that the effective rolling reduction is 20% or more from the viewpoint of increasing toughness. Here, the “effective reduction ratio” refers to the total reduction amount (%) in the temperature range of 950 ° C. or lower. In order to achieve the desired high toughness over the entire plate thickness, it is preferable that the effective rolling reduction at the central portion of the plate thickness satisfies 20% or more, more preferably 40% or more.
After completion of hot rolling (finish rolling), the hot-rolled sheet is subjected to accelerated cooling on a hot run table. It is desirable to start the accelerated cooling while the temperature at the center of the plate thickness is 750 ° C. or higher. When the temperature in the central portion of the plate thickness is less than 750 ° C., high-temperature transformation ferrite (polygonal ferrite) is formed, and a second phase is formed around the polygonal ferrite due to C discharged during the γ → α transformation. For this reason, the precipitation fraction of the second phase becomes high at the center of the plate thickness, and the above-described desired structure cannot be formed.

  Accelerated cooling consists of primary accelerated cooling and secondary accelerated cooling. The primary accelerated cooling and the secondary accelerated cooling may be performed continuously, or an air cooling process within 10 s may be provided between the primary accelerated cooling and the secondary accelerated cooling. By performing air cooling between the primary accelerated cooling and the secondary accelerated cooling, overcooling of the surface layer is prevented. Thereby, the formation of martensite is prevented. The air cooling time is preferably 10 s or less from the viewpoint of preventing the inside of the plate thickness from staying in a high temperature region.

The accelerated cooling in the present invention is performed at a cooling rate of 10 ° C./s or more at the average cooling rate at the center position of the plate thickness. In addition, let the average cooling rate of the plate | board thickness center position in primary accelerated cooling be the average in the temperature range from 750 degreeC-a primary cooling stop. Moreover, let the average cooling rate of the plate | board thickness center position in secondary accelerated cooling be the average in the temperature range from a primary cooling stop to a secondary cooling stop.
If the average cooling rate at the center position of the plate thickness is less than 10 ° C./s, high-temperature transformation ferrite (polygonal ferrite) is likely to be formed, and the precipitation fraction of the second phase is increased at the center portion of the plate thickness. An organization cannot be formed. For this reason, the accelerated cooling after the end of hot rolling is performed at a cooling rate of 10 ° C./s or more at the average cooling rate at the center position of the plate thickness. In addition, Preferably it is 20 degrees C / s or more. In order to avoid the formation of polygonal ferrite, it is particularly preferable to carry out at a cooling rate of 10 ° C./s or more in the temperature range of 750 to 650 ° C.

  In the primary accelerated cooling according to the present invention, the cooling rate within the above range, the average cooling rate at the plate thickness center position (plate thickness center portion), and the average cooling rate at the position 1 mm (surface layer) in the plate thickness direction from the surface, Accelerated cooling adjusted so that the difference in cooling rate is less than 80 ° C./s. In addition, let an average cooling rate be the average between the rolling completion temperature of finish rolling, and primary cooling stop temperature. By adopting primary accelerated cooling that is adjusted so that the difference in cooling rate between the surface layer and the center of the plate thickness is less than 80 ° C / s, bainite or bainitic ferrite is formed even in the vicinity of the surface layer, resulting in ductility. The desired strength / ductility balance can be secured. On the other hand, in accelerated cooling where the cooling rate difference between the center of the plate thickness and the surface layer exceeds 80 ° C / s, the structure in the vicinity of the surface layer, and further in the region up to 5 mm in the plate thickness direction, the martensite phase It becomes easy to become a structure containing, and ductility falls. For this reason, in the present invention, the primary accelerated cooling is performed at a cooling rate of 10 ° C./s or more at the average cooling rate at the plate thickness center position and from the surface to the plate thickness direction. The cooling was limited to accelerated cooling adjusted so that the cooling rate difference from the average cooling rate at a position of 1 mm was less than 80 ° C./s. Such primary accelerated cooling can be achieved by adjusting the water density of the cooling water.

Further, in the present invention, the secondary accelerated cooling performed after the above-described primary accelerated cooling is performed at a cooling rate in the above-described range (an average cooling rate at the plate thickness center position of 10 ° C./s or more), and Cooling in which the difference in cooling rate between the average cooling rate at the plate thickness center position and the average cooling rate at a position 1 mm from the surface in the plate thickness direction is 80 ° C / s or higher, the temperature at the plate thickness center position is the following (2 )formula
BFS (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni−1.5CR (2)
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%), CR: cooling rate (° C./s))
Cooling performed to the secondary cooling stop temperature below the BFS defined in. If the difference in cooling rate between the average cooling rate at the center position of the plate thickness in secondary accelerated cooling and the average cooling rate at the position of 1 mm from the surface in the plate thickness direction is less than 80 ° C / s, the structure at the center of the plate thickness is desired. (A structure composed of a bainitic ferrite phase rich in ductility, a bainite phase, or a mixed structure thereof). On the other hand, when the secondary cooling stop temperature exceeds BFS, polygonal ferrite is formed, the second phase structure fraction increases, and desired properties cannot be secured. For this reason, secondary accelerated cooling is performed when the difference in cooling rate between the average cooling rate at the center of the plate thickness and the average cooling rate at the position of 1 mm from the surface in the plate thickness direction is 80 ° C / s or more. It was assumed that the secondary cooling stop temperature was below the BFS at the temperature at the center position. The secondary cooling stop temperature is more preferably (BFS-20 ° C) or lower.

After the secondary accelerated cooling is stopped at the secondary cooling stop temperature or lower, the hot rolled sheet is wound in a coil shape at a winding temperature of BFS0 or lower. More preferably, it is (BFS 0-20 ° C.) or less. BFS0 is the following formula (3)
BFS0 (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni (3)
(Here, C, Mn, Cr, Mo, Cu, Ni: content of each element (mass%))
Defined by

  By setting the cooling stop temperature of secondary accelerated cooling to a temperature below BFS and the coiling temperature to a temperature below BFS0, as shown in FIGS. 2 and 3, ΔD is 2 μm or less and ΔV is 2% for the first time. The uniformity of the structure in the plate thickness direction becomes remarkable. Thereby, the excellent DWTT characteristic and the outstanding CTOD characteristic can be ensured, and it can be set as the thick-walled high-tensile-strength hot-rolled steel plate which markedly improved low-temperature toughness.

  In the secondary accelerated cooling according to the present invention, the difference between the cooling stop temperature at a position of 1 mm from the surface in the plate thickness direction and the coiling temperature (temperature at the plate thickness central position) when the secondary cooling is stopped is as follows. It is preferable to apply so that it may be within 300 degreeC. When the difference between the cooling stop temperature and the coiling temperature at a position of 1 mm from the surface in the plate thickness direction exceeds 300 ° C, a composite structure containing a martensite phase is formed in the surface layer depending on the steel composition, and ductility decreases. However, the desired strength / ductility balance may not be ensured. For this reason, in the secondary accelerated cooling in the present invention, the difference between the cooling stop temperature at a position of 1 mm from the surface in the thickness direction and the coiling temperature (temperature at the thickness center position) is within 300 ° C. It is said that it is preferable to apply. Such secondary acceleration cooling adjustment can be achieved by adjusting the water density or selecting a cooling bank.

  In addition, although the upper limit of a cooling rate is determined depending on the capability of the cooling device to be used, it is preferable that it is slower than the martensite production cooling rate which is a cooling rate without the deterioration of steel plate shapes, such as curvature. Such a cooling rate can be achieved by cooling using a flat nozzle, a rod-like nozzle, a circular tube nozzle, or the like. In the present invention, the temperature at the center of the plate thickness, the cooling rate, and the like calculated by heat transfer calculation are used.

In addition, it is preferable that the hot-rolled sheet wound up in a coil shape is cooled to room temperature at a rate of 20 to 60 ° C./hr at a cooling rate at the center of the coil. If the cooling rate is less than 20 ° C./hr, the growth of crystal grains proceeds, so that the toughness may decrease. Further, at a cooling rate exceeding 60 ° C./hr, the temperature difference between the coil central portion and the coil outer peripheral portion or inner peripheral portion becomes large, and the coil shape tends to deteriorate.
The thick high-tensile hot-rolled steel sheet of the present invention obtained by the above-described manufacturing method has the above-described composition, and further has a structure having a ferrite phase as a main phase at least 1 mm from the surface in the sheet thickness direction. . The “ferrite phase” here means a hard low-temperature transformation ferrite, and means any one of a bainite phase, a bainitic ferrite phase, or a mixed phase thereof. Soft high temperature transformation ferrite (granular polygonal ferrite) is not included. The second phase is any one of a pearlite phase, a martensite phase, MA (also called island martensite), upper bainite, or a mixed phase composed of two or more thereof. In the thick high-tensile hot-rolled steel sheet of the present invention, it goes without saying that the structure at the center position of the sheet thickness is a structure having a similar ferrite phase as the main phase. Here, the “main phase” refers to a case where the structure fraction (volume%) is 90% or more, more preferably 98% or more.

The difference ΔD between the average crystal grain size of the ferrite phase at a position 1 mm from the steel sheet surface and the average crystal grain size (μm) of the ferrite phase at the center position of the plate thickness is 2 μm or less, and the plate thickness from the surface. It has a structure in which the difference ΔV between the structure fraction (volume%) of the second phase at a position of 1 mm in the direction and the structure fraction (volume%) of the second phase at the center position of the plate thickness is 2% or less.
Only when ΔD is 2 μm or less and ΔV is 2% or less, the low-temperature toughness of the thick, high-tensile hot-rolled steel sheet, in particular, the DWTT characteristics and CTOD characteristics using a full-thickness test piece are significantly improved. When either one of ΔD or ΔV falls outside the desired range, as is apparent from FIG. 1, DWTT is higher than −35 ° C., DWTT characteristics are lowered, and low-temperature toughness is deteriorated. For this reason, in the present invention, the microstructure is the difference between the average crystal grain size of the ferrite phase at a position 1 mm from the steel sheet surface in the thickness direction and the average crystal grain size (μm) of the ferrite phase at the center position of the thickness. ΔV is 2 μm or less, and the difference ΔV between the structure fraction (volume%) of the second phase at the position 1 mm from the surface in the sheet thickness direction and the structure fraction (volume%) of the second phase at the sheet thickness center position is 2 It was limited to the organization which is less than%. By having such a composition and structure, it is possible to obtain a steel sheet having an excellent balance between strength and ductility.

  Note that a hot-rolled steel sheet having a structure in which ΔD is 2 μm or less and ΔV is 2% or less is the average grain size of the ferrite phase at a position of 1 mm and a thickness of 1/4 in the thickness direction from the steel sheet surface ( μm) difference ΔD * is 2 μm or less, second phase structure fraction (%) difference ΔV * is 2% or less, and the position of 1 mm from the steel sheet surface in the thickness direction and the thickness of 3/4 position It has been confirmed that the difference ΔD ** in the average crystal grain size (μm) of the ferrite phase satisfies 2 μm or less and the difference ΔV ** in the structure fraction (%) of the second phase also satisfies 2% or less.

  Hereinafter, the present invention will be described in detail based on examples.

  Using a slab (steel material) (thickness: 215 mm) having the composition shown in Table 1, hot rolling is performed under the hot rolling conditions shown in Table 2, and after completion of the hot rolling, cooling is performed under the cooling conditions shown in Table 2. And it wound up in coil shape at the coiling temperature shown in Table 2, and it was set as the hot-rolled steel plate (steel strip) of the board thickness shown in Table 2. Using these hot-rolled steel sheets as the raw material, open pipes were formed by continuous roll forming in the cold, and the end faces of the open pipes were electro-welded to form electric-welded steel pipes (outer diameter 660 mmφ).

Test pieces were collected from the obtained hot-rolled steel sheet and subjected to structure observation, tensile test, impact test, DWTT test, and CTOD test. The DWTT test and CTOD test were also conducted on ERW steel pipes. The test method was as follows.
(1) Microstructure observation A specimen for microstructural observation is collected from the obtained hot-rolled steel sheet, the cross section in the rolling direction is polished and corroded, and the optical microscope (magnification: 1000 times) or scanning electron microscope (magnification: 2000 times) is used. Observe at least 2 fields of view, identify the type of tissue by imaging, and use an image analyzer to determine the average crystal grain size of the ferrite phase and the fraction of the second phase other than the ferrite phase (volume%) It was measured. The observation position was a position of 1 mm in the thickness direction from the surface of the steel plate and the central portion of the thickness. The average crystal grain size of the ferrite phase is obtained by measuring the area of each ferrite grain, calculating the equivalent circle diameter from the area, arithmetically averaging the equivalent circle diameter of each obtained ferrite grain, and calculating the average crystal at the position. The particle size was taken.

(2) Tensile test From the obtained hot-rolled steel sheet, a plate-shaped specimen (parallel part width: 12.5 mm, distance between gauge points: 50 mm) so that the direction perpendicular to the rolling direction (C direction) is the longitudinal direction. ), And a tensile test was performed at room temperature in accordance with ASTM E 8 to determine the tensile strength TS and elongation El, and the strength / ductility balance TS × El was calculated.

(3) Impact test V-notch test specimens were taken from the center of the thickness of the obtained hot-rolled steel sheet so that the direction perpendicular to the rolling direction (C direction) was the longitudinal direction, and conformed to the provisions of JIS Z 2242 Then, a Charpy impact test was performed, and the absorbed energy (J) at a test temperature of −80 ° C. was obtained. The number of test pieces was three, and the arithmetic average of the obtained absorbed energy values was obtained to obtain the absorbed energy value vE- ₈₀ (J) of the steel sheet. The case where vE- ₈₀ was 300 J or more was evaluated as “good toughness”.

(4) DWTT test From the obtained hot-rolled steel sheet, a DWTT test piece (size: plate thickness x width 3 in. X length 12 in.) Was set so that the direction perpendicular to the rolling direction (C direction) was the longitudinal direction. The sample was collected and subjected to a DWTT test in accordance with ASTM E 436, and the lowest temperature (DWTT) at which the ductile fracture surface ratio was 85% was determined. The case where DWTT was −35 ° C. or less was evaluated as having “excellent DWTT characteristics”.

In the DWTT test, a DWTT test piece was sampled from the base material portion of the ERW steel pipe so that the longitudinal direction of the test piece became the pipe circumferential direction, and tested in the same manner as the steel plate.
(5) CTOD test From the obtained hot-rolled steel sheet, a CTOD specimen (size: plate thickness x width (2 x plate thickness) x length so that the direction orthogonal to the rolling direction (C direction) is the longitudinal direction. (10 × plate thickness)) was collected, and a CTOD test was conducted at a test temperature of −10 ° C. in accordance with the provisions of ASTM E 1290 to obtain a critical opening displacement (CTOD value) at −10 ° C. The test load was applied by a three-point bending method, a displacement meter was attached to the notch, and the critical opening displacement CTOD value was obtained. The case where the CTOD value was 0.30 mm or more was evaluated as having “excellent CTOD characteristics”.

In addition, CTOD test is also taken from the ERW steel pipe, so that the direction orthogonal to the tube axis direction is the longitudinal direction of the test piece, CTOD test piece is taken, and the notch is introduced into the base metal part and the seam part, Tested in the same manner as the steel sheet.
The obtained results are shown in Table 3.

Each of the inventive examples has an appropriate structure, TS: high strength of 510 MPa or more, DWTT of vE- ₈₀ of 300 J or more, CTOD value of 0.30 mm or more, −35 ° C. or less, and excellent low temperature toughness. Furthermore, it is a hot rolled steel sheet having an excellent balance of strength and ductility of TS × El: 18000 MPa% or more. In addition, the ERW steel pipe using the hot-rolled steel sheet of the example of the present invention has a CTOD value of 0.30 mm or more, a DWTT of −20 ° C. or less in both the base material portion and the seam portion, and a steel pipe having excellent low temperature toughness It has become.

On the other hand, comparative examples that are out of the scope of the present invention show that the low temperature toughness is reduced as vE- ₈₀ is less than 300 J, the CTOD value is less than 0.30 mm, or the DWTT exceeds -35 ° C. Or the elongation is low and the desired balance between strength and ductility cannot be ensured.

Claims (9)

% By mass
C: 0.02 to 0.08%, Si: 0.01 to 0.50%,
Mn: 0.5 to 1.8%, P: 0.025% or less,
S: 0.005% or less, Al: 0.005-0.10%,
Nb: 0.01-0.10%, Ti: 0.001-0.05%
And containing C, Ti, Nb so as to satisfy the following formula (1), heating a steel material having a composition composed of the balance Fe and inevitable impurities, and performing hot rolling consisting of rough rolling and finish rolling To make a hot-rolled steel sheet,
The accelerated cooling is a cooling consisting of primary accelerated cooling and secondary accelerated cooling, and the primary accelerated cooling is performed at an average cooling rate at a thickness center position of 10 ° C./s or more and an average cooling rate at a thickness center position. Cooling with a cooling rate difference of less than 80 ° C / s from the average cooling rate at the position of 1mm from the surface to the plate thickness direction, the temperature at the position of 1mm from the surface to the plate thickness direction is 650 ° C or less and 500 ° C or more The cooling to be performed up to the primary cooling stop temperature, which is the temperature range of the temperature range, and the secondary accelerated cooling is performed at an average cooling rate at the plate thickness center position of 10 ° C./s or more, from the average cooling rate at the plate thickness center position and the surface Cooling in which the difference in cooling rate from the average cooling rate at the position of 1 mm in the plate thickness direction is 80 ° C / s or more, the temperature at the plate thickness center position is the BFS or less that is defined by the following formula (2) The cooling is performed up to the cooling stop temperature. After the secondary accelerated cooling, the temperature at the center position of the plate thickness is expressed by the following equation (3). BFS0 following winding method for manufacturing a thick-walled high-strength hot-rolled steel sheet with excellent strength-ductility balance, characterized in that winding at the temperature defined.
(Ti + (Nb / 2)) / C <4 (1)
BFS (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni−1.5CR (2)
BFS0 (℃) = 770−300C−70Mn−70Cr−170Mo−40Cu−40Ni (3)
Here, C, Mn, Cr, Mo, Cu, Ni: Content of each element (mass%)
CR: Cooling rate (° C / s)   The method for producing a thick, high-tensile hot-rolled steel sheet according to claim 1, wherein air cooling is performed for 10 seconds or less between the primary accelerated cooling and the secondary accelerated cooling.   3. The high-thickness high-tensile hot rolling according to claim 1, wherein the accelerated cooling is 10 ° C./s or more at an average cooling rate in a temperature range of 750 to 650 ° C. at a center position of the plate thickness. A method of manufacturing a steel sheet.   4. The difference between the cooling stop temperature at a position of 1 mm from the surface in the plate thickness direction and the coiling temperature in the secondary accelerated cooling is within 300 ° C. 4. A method for producing a thick, high-tensile hot-rolled steel sheet.   In addition to the above composition, in addition to mass, V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: 0.01 to 0.50%, Ni: 0.01 to 0.50% Or it is set as the composition containing 2 or more types, The manufacturing method of the thick-wall high tension hot-rolled steel plate in any one of Claim 1 thru | or 4 characterized by the above-mentioned.   The method for producing a thick high-tensile hot-rolled steel sheet according to any one of claims 1 to 5, wherein in addition to the composition, the composition further contains Ca: 0.0005 to 0.005% by mass. % By mass
C: 0.02 to 0.08%, Si: 0.01 to 0.50%,
Mn: 0.5 to 1.8%, P: 0.025% or less,
S: 0.005% or less, Al: 0.005-0.10%,
Nb: 0.01-0.10%, Ti: 0.001-0.05%
And containing C, Ti, Nb so as to satisfy the following formula (1), the composition consisting of the balance Fe and inevitable impurities, and the structure at a position of 1 mm from the surface in the plate thickness direction is the main phase. The difference ΔD between the average crystal grain size of the ferrite phase at a position 1 mm from the surface in the plate thickness direction and the average crystal grain size of the ferrite phase at the plate thickness center position is 2 μm or less, and the plate thickness direction from the surface The difference ΔV between the structure fraction (volume%) of the second phase at the position of 1 mm and the structure fraction (volume%) of the second phase at the center position of the plate thickness is 2% or less. Thick, high-tensile hot-rolled steel sheet with excellent strength and ductility balance.
(Ti + (Nb / 2)) / C <4 (1)
Here, Ti, Nb, C: Content of each element (mass%)   In addition to the above composition, in addition to mass, V: 0.01 to 0.10%, Mo: 0.01 to 0.50%, Cr: 0.01 to 1.0%, Cu: 0.01 to 0.50%, Ni: 0.01 to 0.50% Or it is set as the composition containing 2 or more types, The thick-walled high tension hot-rolled steel plate of Claim 7 characterized by the above-mentioned.   The thick-walled high-tensile hot-rolled steel sheet according to claim 7 or 8, further comprising, in addition to the composition, Ca: 0.0005 to 0.005% by mass.

Images