JP2010196165A - Thick high-tensile-strength hot-rolled steel sheet having excellent low-temperature toughness and process for producing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a process for the production of a thick high-tensile-strength hot-rolled steel sheet which combines high strength and high toughness. <P>SOLUTION: When a steel material having a composition containing 0.02 to 0.08% C, 0.01 to 0.10% Nb and 0.001 to 0.05% Ti, where C, Ti and Nb satisfy (Ti+(Nb/2))/C<4, is heated and hot-rolled, and is then subjected to accelerated cooling; the residence time from the sheet thickness central temperature T(°C) to a temperature (T-20°C) is controlled to ≤20s upon the start of the accelerated cooling, and then, accelerated cooling at a cooling rate of ≥10°C/s in the center of the sheet thickness is performed for ≤30s from the T(°C) to below a specified cooling end temperature dependent on the cooling rate, and taking up is performed at below a specified take up temperature dependent on the amounts of the alloy elements. In this way, the thick high-tensile-strength hot-rolled steel sheet, where the average crystal grain size of a ferrite phase in the center of the sheet thickness is ≤5 μm, while the structural fraction of the second phase is ≤2% and the structural uniformity in the sheet thickness direction is excellent, and low temperature toughness and, particularly, DWTT characteristics are remarkably improved, is obtained. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention is 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, etc. The present invention relates to a hot-rolled steel sheet and a manufacturing method thereof, and particularly relates to 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 technique described in Patent Document 1, 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%, Al: 0.06% or less, Ti: 0.005-0.035%, N: 0.001 in Ac ₁ -50 ° C. below the temperature of the slab course the after hot rolling cooling containing 0.006%, performs rolling over 2% cumulative, then heated to a temperature of less than Ac ₁ super Ac ₃ A method of manufacturing a steel material having a coarse ferrite layer on the front and back surfaces is 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 After heating a steel slab containing one or more of Mo and having a Pcm value of 0.17 or less, finish rolling is completed under conditions where the surface temperature is (Ar3-50 ° C) or more, and rolling A method for producing a hot-rolled steel strip for a high-strength electric-welded pipe excellent in low-temperature toughness and weldability that is immediately cooled and wound up and slowly cooled at a temperature of 700 ° C. or less 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, 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. In particular, in the case of a hot-rolled steel sheet with an extremely thick thickness exceeding 22 mm, the center of the plate thickness tends to be slower to cool than the surface layer, and the crystal grain size at the center of the plate thickness tends to become coarser. There is a problem that further improvement is difficult.

  The present invention solves the above-mentioned problems of the prior art, has a thickness exceeding 22 mm without requiring a large amount of alloying elements, and has a high tensile strength of 530 MPa or more and excellent low temperature toughness. An ultra-thick high-tensile hot-rolled steel sheet suitable for X70 to X80 grade high-strength ERW steel pipes or high-strength spiral steel pipes, which has particularly excellent CTOD characteristics and DWTT characteristics, and a method for producing the same For the purpose. “Excellent low temperature toughness” as used herein is a DWTT test conducted in accordance with the provisions of ASTM E 436. The minimum temperature (DWTT) at which the ductile fracture surface ratio is 85% is −30 ° C. or lower, and In accordance with ASTM E 1290, this refers to a 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 order to achieve the above-mentioned object, the present inventors diligently studied various factors affecting low temperature toughness, particularly DWTT characteristics. As a result, it was found that the DWTT characteristic, which is a toughness test at the full thickness, is greatly influenced by the structure uniformity in the thickness direction. The present inventors have also found that the influence of the structure non-uniformity in the thickness direction on the DWTT characteristics, which is a toughness test at the full thickness, is manifested by a thick material having a thickness of 11 mm or more.

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

Further, 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 mainly depend on the amount of alloying elements contained. It is determined. That is, in order to set ΔD to 2 μm or less and ΔV to 2% or less, the cooling stop temperature is expressed by the following equation:
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 below the BFS defined by
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.

However, in an extremely thick hot-rolled steel sheet having a thickness exceeding 22 mm, even if ΔD and ΔV are within the above-described ranges, the DWTT characteristics are lowered and the desired “excellent DWTT characteristics” cannot be ensured. Therefore, the present inventors, in the ultra-thick hot-rolled steel sheet having a plate thickness of more than 22 mm, delays the cooling of the center portion of the plate thickness compared to the surface layer portion, and tends to coarsen the crystal grains. Considering that the diameter increases and the second phase increases, we have further studied diligently on the adjustment method of the thickness center structure of the extra-thick hot-rolled steel sheet. As a result, the time during which the steel sheet stays in the high temperature range is set to 20 s or less after the finish rolling, after the finish rolling, until the temperature T (° C.) decreases from the temperature T (° C.) at the start of accelerated cooling to 20 ° C. In addition, the temperature at the center position of the steel sheet thickness after finishing finish rolling is calculated from the temperature T (° C.) at the start of accelerated cooling by the following formula (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))
It was found that the cooling time to the BFS temperature defined in (3) was 30 s or less. Thereby, the structure of the plate thickness central part can be made a structure in which the average crystal grain size of the ferrite phase is 5 μm or less and the structure fraction (volume%) of the second phase is 2% or less.

First, the experimental results on which the present invention is based will be described.
A slab composed of 0.039% C-0.24% Si-1.61% Mn-0.019% P-0.0023% S-0.038% Al-0.059% Nb-0.010% Ti-balance Fe in mass% was used as a steel material. Note that (Ti + Nb / 2) / C is 1.0.
The steel material having the composition described above is heated to 1200 ° C, hot rolled to a finish rolling start temperature of 1000 ° C and a finish rolling finish temperature of 800 ° C to obtain a hot rolled sheet having a thickness of 23.8 mm. After the end of rolling, accelerated cooling was performed under various conditions, followed by winding at various winding temperatures to obtain a hot rolled steel sheet (steel strip).

  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 (μm) of the ferrite phase, the structure fraction (volume%) of the second phase ) From the measured values obtained, the average crystal grain size difference ΔD of the ferrite phase and the structure of the second phase at a position (surface layer portion) 1 mm from the surface (surface layer portion) and the plate thickness central position (plate thickness center portion). The rate difference ΔV was calculated respectively. The “ferrite phase” here refers to either a bainitic ferrite phase, a bainite phase, or a mixed phase thereof.

The obtained results are shown in FIG. 1 in relation to the average crystal grain size of the ferrite phase at the central portion of the plate thickness on the DWTT and the structure fraction of the second phase. FIG. 1 shows a case where ΔD is 2 μm or less and ΔV is 2% or less.
From FIG. 1, when the average crystal grain size of the ferrite phase at the center of the plate thickness is 5 μm or less and the structure fraction of the second phase is 2% or less, DWTT is − It can be seen that the steel sheet has “excellent DWTT characteristics” at 30 ° C. or lower.

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 composition comprising the balance Fe and inevitable impurities, the average crystal grain size of the ferrite phase at the center of the plate thickness is 5 μm or less, and the second phase structure fraction (volume%) is 2% or less. 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 position 1 mm from the surface in the plate thickness direction. The low-temperature toughness is characterized by having a structure in which the difference ΔV between the second phase structure fraction (volume%) and the second phase structure percentage (volume%) at the center of the plate thickness is 2% or less. Excellent thickness: Extra-thick high-tensile hot-rolled steel sheet exceeding 22 mm.

(2) In (1), 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 An ultra-thick high-tensile hot-rolled steel sheet having a composition containing one or more of ˜0.50%.
(3) The ultra-thick high-tensile hot-rolled steel sheet according to (1) or (2), further comprising Ca: 0.0005 to 0.005% by mass% in addition to the above composition.

(4) 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 of the composition comprising the balance Fe and unavoidable impurities is heated so as to satisfy, and hot rolling comprising rough rolling and finish rolling is performed to form a hot rolled steel sheet, and then after the finish rolling is completed. The above-mentioned hot-rolled steel sheet is subjected to accelerated cooling at an average cooling rate of 10 ° C./s or more at the center position of the thickness in the following 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))
To the cooling stop temperature below the BFS defined in, and then 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%))
When the coil is wound at a coiling temperature of BFS0 or less as defined by the above, the temperature at the thickness center position of the hot-rolled steel sheet is changed from the temperature at the start of the accelerated cooling: T (° C) to the temperature: (T-20 ° C) The plate thickness with excellent low-temperature toughness is characterized in that the residence time until reaching is within 20 s and the cooling time from the temperature T at the center of the plate thickness to the temperature of the BFS is adjusted to 30 s or less. : A method for producing ultra-thick high-tensile hot-rolled steel sheets exceeding 22 mm.

(5) In (4), 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 method for producing an ultra-thick high-tensile hot-rolled steel sheet, characterized in that the composition contains one or more of ˜0.50%.
(6) In (4) or (5), in addition to the said composition, it is set as the composition which contains further Ca: 0.0005-0.005% by the mass%, The manufacturing method of the ultra-thick high-tensile hot-rolled steel sheet characterized by the above-mentioned. .

  According to the present invention, the structure at the center of the plate thickness is miniaturized, the variation in the structure in the plate thickness direction is small, the plate thickness: an extreme thickness exceeding 22 mm, the tensile strength TS: high strength of 530 MPa or more, and excellent In addition, an ultra-thick high-tensile hot-rolled steel sheet having both excellent low-temperature toughness, particularly excellent DWTT characteristics and CTOD characteristics, can be produced easily and inexpensively, and the industrially remarkable effect is achieved. In addition, according to the present invention, there is also an effect that an ERW steel pipe for line pipe and a spiral steel pipe for line pipe, which are excellent in low temperature toughness and circumferential weldability when laying a pipeline, can be easily manufactured.

It is a graph which shows the relationship between the average crystal grain diameter of the ferrite phase in plate | board thickness center position, and the structure fraction of a 2nd phase which affects DWTT.

First, the reasons for limiting the composition of the ultra-thick high-tensile hot-rolled steel sheet of 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. In order to acquire such an effect, it is desirable to contain 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 0.45%.

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. Such an effect becomes remarkable when the content is 0.0005% or more. However, if the content exceeds 0.005%, CaO is increased and corrosion resistance and toughness are lowered. 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.

  The ultra-thick high-tensile hot-rolled steel sheet of the present invention has the above-described composition, and further, the average crystal grain size of the ferrite phase at the center of the sheet thickness is 5 μm or less, and the second phase structure fraction (volume%) is 2. And the difference ΔD between the average crystal grain size of the ferrite phase at a position 1 mm from the steel sheet surface in the plate thickness direction 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 A structure in which the difference ΔV between the structure fraction (volume%) of the second phase at a 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% or less. Have. The term “ferrite phase” used herein means hard low-temperature transformation ferrite, and refers to any of bainitic ferrite phase, bainite phase, or a mixed phase thereof. Soft high temperature transformation ferrite (granular polygonal ferrite) is not included. The structure at the center of the plate thickness is that the main phase is either bainitic ferrite phase, bainite phase, or a mixed phase thereof, and the second phase is pearlite, martensite, MA (martensite-austenite constituent, island martensite). ), Upper bainite, or a mixed phase composed of two or more thereof. Here, the “main phase” refers to a case where the structure fraction (volume%) is 90% or more, more preferably 98% or more.

  When ΔD is 2 μm or less and ΔV is 2% or less, low-temperature toughness, particularly DWTT characteristics and CTOD characteristics using a full-thickness test piece are remarkably improved. When either one of ΔD or ΔV falls outside the desired range, the DWTT characteristic is lowered and the low temperature toughness is deteriorated. If the plate thickness is more than 22 mm, the average grain size of the ferrite phase at the center of the plate thickness should be 5 μm or less, and the second phase structure fraction (volume%) should be 2% or less. I need. When the average crystal grain size of the ferrite phase exceeds 5 μm, or when the structure fraction (volume%) of the second phase exceeds 2%, the DWTT characteristics are lowered and the low temperature toughness is deteriorated.

  For this reason, in the present invention, the structure is such that the average crystal grain size of the ferrite phase at the center position of the plate thickness is 5 μm or less, the structure fraction (volume%) of the second phase is 2% or less, and the steel sheet surface The difference ΔD between the average crystal grain size of the ferrite phase at a position of 1 mm in the plate thickness direction and the average crystal grain size (μm) of the ferrite phase at the plate thickness center position is 2 μm or less, and the position of 1 mm from the surface in the plate thickness direction The difference ΔV between the structure fraction (volume%) of the second phase and the second phase structure volume (volume%) at the center position of the plate thickness was limited to a structure having 2% or less.

  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.

Below, the preferable manufacturing method of this invention hot-rolled steel plate is demonstrated.
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 in the entire plate thickness, it is preferable that the effective rolling reduction at the central portion of the plate thickness satisfies 20% or more. The finish rolling finish temperature FDT is preferably set to 750 ° C. or higher for accelerated cooling thereafter.

  After completion of hot rolling (finish rolling), the hot-rolled sheet is subjected to accelerated cooling on a hot run table. In the present invention, in order to set the crystal grain size of the ferrite phase at the center position of the plate thickness to a predetermined value or less and the structure fraction of the second phase to 2% or less by volume, accelerated cooling after finishing rolling is finished. The residence time from the temperature T (° C) at the center of the steel plate thickness at the start (hereinafter also referred to as the cooling start point) to the temperature of (T-20 ° C) is set within 20 s, and the residence time at high temperature is shortened. To do. If the residence time from T (° C.) to (T-20 ° C.) is longer than 20 s, the crystal grain size at the time of transformation tends to become coarse, and high temperature transformation ferrite (polygonal ferrite) is generated. It becomes difficult to avoid. In addition, in order to make the residence time from T (° C.) to (T-20 ° C.) within 20 s, in the plate thickness range of the steel plate of the present invention, the plate passing speed on the hot run table is 120 mpm or more. It is preferable to do.

  Moreover, it is desirable to start the accelerated cooling while the temperature at the central portion of the plate thickness is 750 ° C. or higher. When the temperature at the center of the plate thickness is less than 750 ° C., high-temperature transformation ferrite (polygonal ferrite) is formed, and C discharged during γ → α transformation is concentrated to untransformed γ. The second phase is formed around polygonal ferrite. For this reason, the structure fraction of the second phase increases at the center of the plate thickness, and the above-described desired structure cannot be formed.

The accelerated cooling is preferably performed at a cooling rate of 10 ° C./s or higher, preferably 20 ° C./s or higher, at an average cooling rate at the center of the plate thickness, to a cooling stop temperature of BFS or lower.
When the cooling rate is less than 10 ° C./s, high-temperature transformation ferrite (polygonal ferrite) is likely to be formed, the second phase structure fraction becomes high at the center of the plate thickness, and the above-described desired structure cannot be formed. For this reason, it is preferable to perform accelerated cooling after completion | finish of hot rolling with the cooling rate of 10 degrees C / s or more by the average cooling rate of a plate | board thickness center part. 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. In addition, such a cooling rate can be achieved by a water cooling device 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.

Further, the cooling stop temperature of the above-described accelerated cooling is preferably set to a temperature equal to or lower than BFS at the center thickness position. In addition, More preferably, it is (BFS-20 degreeC) or less. BFS is the following equation (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))
Defined by

  In the present invention, in order to set the crystal grain size of the ferrite phase at the center position of the plate thickness to a predetermined value or less and the structure fraction of the second phase to 2% or less by volume, the above-described cooling start point T Adjust the cooling time from (℃) to BFS temperature to 30s or less. When the cooling time from T (° C.) to the BFS temperature is longer than 30 s, high temperature transformation ferrite (polygonal ferrite) is likely to be formed, and C discharged during γ → α transformation is concentrated to untransformed γ, A second phase such as a pearlite phase or upper bainite is formed around polygonal ferrite. For this reason, the structure fraction of the second phase increases at the center of the plate thickness, and the above-described desired structure cannot be formed. For this reason, the cooling time from the cooling start point T (° C.) to the BFS temperature is limited to 30 s or less. The adjustment of the cooling time from the cooling start point T (° C.) to the BFS temperature can be performed by adjusting the plate passing speed and the cooling water amount.

Further, in the present invention, after the accelerated cooling is stopped below the above-described cooling stop temperature, the hot-rolled sheet is wound in a coil shape at a winding temperature of BFS 0 or less at the temperature at the center position of the plate thickness. 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 accelerated cooling to BFS or less and the coiling temperature to BFS0 or less, ΔD is 2 μm or less and ΔV is 2% or less, and the uniformity of the structure in the thickness direction is remarkable. It becomes. Thereby, it is possible to ensure excellent DWTT characteristics and excellent CTOD characteristics.
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.

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

  Using a slab (steel material) (thickness: 230 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 3 fields of view, image, identify the structure, and use an image analyzer to determine the average crystal grain size of the ferrite phase and the structure fraction (volume%) of the second phase other than the ferrite phase. It was measured. The observation position was a position of 1 mm in the thickness direction from the surface of the steel plate and a central position of the thickness. The average crystal grain size of the ferrite phase was determined by a cutting method, and the nominal grain size was defined as the average crystal grain size at the position.

(2) Tensile test From the obtained hot-rolled steel sheet, a plate-shaped specimen (parallel part width: 25 mm, distance between gauge points: 50 mm) so that the direction (C direction) perpendicular to the rolling direction is the tensile test direction. ), And a tensile test was performed at room temperature in accordance with ASTM E8M-04 to determine the tensile strength TS.
(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 200 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 −30 ° 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 sampled and subjected to a CTOD test at a test temperature of −10 ° C. in accordance with ASTM E 1290, and a critical opening displacement (CTOD value) at −10 ° C. was obtained. The test load was applied by a three-point bending method, a displacement gauge was attached to the notch, and the critical opening displacement CTOD value was obtained.If the CTOD value is 0.30 mm or more, it has “excellent CTOD characteristics”. Then I evaluated.

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 examples of the present invention has an appropriate structure, TS: high strength of 530 MPa or more, DWTT of vE- ₈₀ of 200 J or more, CTOD value of 0.30 mm or more, −30 ° C. or less, and excellent low temperature toughness And has excellent CTOD characteristics and excellent DWTT characteristics. The ERW steel pipe using the hot-rolled steel sheet of the example of the present invention also has a CTOD value of 0.30 mm or more and a DWTT of −5 ° C. or less for both the base metal part and the seam part, and has excellent low temperature toughness. Yes.

On the other hand, the comparative example out of the scope of the present invention shows that the low temperature toughness is lowered because the vE- ₈₀ is less than 200 J, the CTOD value is less than 0.30 mm, or the DWTT exceeds -20 ° C. Yes.

Claims (6)

% 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 C, Ti, Nb so as to satisfy the following formula (1), the composition comprising the balance Fe and inevitable impurities, and the average crystal grain size of the ferrite phase at the center position of the plate thickness is 5 μm or less, The average crystal grain size of the ferrite phase at a position of 1 mm from the surface in the plate thickness direction and the average crystal grain size of the ferrite phase at the plate thickness central position is 2% or less in volume fraction of the two phases. The difference ΔV is 2 μm or less, and the difference ΔV between the structure fraction (volume%) of the second phase at the position of 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. A sheet thickness excellent in low-temperature toughness characterized by having a structure of not more than 10%: extra-thick high-tensile hot-rolled steel sheet exceeding 22 mm.
(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 ultra-thick high tension hot-rolled steel plate of Claim 1 characterized by the above-mentioned.   The extra-thick high-tensile hot-rolled steel sheet according to claim 1 or 2, 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), 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 Then, the hot rolled steel sheet after the finish rolling is finished, and the accelerated cooling of 10 ° C./s or more at the average cooling rate at the center position of the sheet thickness is defined by the following formula (2). When the coil is wound up to a cooling stop temperature of BFS or less and then wound at a coiling temperature of BFS 0 or less defined by the following equation (3), the temperature at the center of the thickness of the hot-rolled steel sheet is The residence time from temperature: T (° C.) to temperature: (T-20 ° C.) should be within 20 s, and the cooling time from the temperature T at the plate thickness center position to the BFS temperature should be 30 s or less. Thickness excellent in low temperature toughness characterized by adjusting to: Production method.
(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)   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 ultra-thick high tension hot-rolled steel plate of Claim 4 characterized by the above-mentioned.   The method for producing an ultra-thick high-tensile hot-rolled steel sheet according to claim 4 or 5, wherein, in addition to the composition, the composition further contains Ca: 0.0005 to 0.005% by mass.

Images