JP5833964B2 - Steel sheet excellent in bending workability, impact property and tensile property, and method for producing the same - Google Patents

Description

  The present invention relates to a steel plate excellent in bending workability, impact properties (base material toughness and toughness after bending) and tensile properties, and a method for producing the same, and when the plate thickness is large (for example, about 100 mm). In addition to excellent tensile properties (yield strength, tensile strength) and toughness of the base material, the present invention also relates to a steel sheet having excellent bending workability and toughness after bending, and also excellent in HAZ toughness and weldability, and a method for producing the same. Is.

  High tensile strength steel plates with a yield strength of 500 MPa or more used as welded structural materials such as bridges, ships, offshore structures, pressure vessels, line pipes, etc. require toughness and weldability in addition to strength. It is also required to secure weldability at In addition to excellent cold bending workability, it is also required to ensure excellent toughness after bending and secure good low temperature toughness for use in cold regions of about -20 to -50 ° C. There is a case. In particular, for cold bending, extremely severe cold bending such as a bending inner radius of 2.5 t like a square steel pipe may be performed. Even in such a case, it is required to ensure toughness after cold bending. Many studies have been made to improve these characteristics, and many proposals have been made on the component composition and production conditions of a steel sheet for improving the characteristics.

  In the old days, there are a method in which reheating and quenching is performed off-line and further reheating and tempering, and a method in which quenching is performed immediately after rolling a steel sheet, so-called direct quenching, followed by a tempering process offline. However, these require an offline tempering process, and there are problems such as a decrease in productivity and a longer work period. In recent years, the tempering process is omitted and no offline heat treatment is required, so-called non-adjustment. Various quality manufacturing methods have been proposed.

  As a non-refining manufacturing method, for example, in Patent Document 1, it is added to obtain strength in a conventional non-tempering process by utilizing precipitation strengthening by Nb carbonitride and Ti carbide as components. As a non-refining process, cooling is performed at a cooling rate of 2 to 30 ° C / second from a temperature range of 800 ° C or higher, and then 550 to 700 ° C. By cooling at a cooling rate of 0.4 ° C./second or less from the temperature range, a high-tensile steel sheet having high tensile strength with a yield strength of 450 MPa or more, small acoustic anisotropy and excellent weldability is obtained. It has been proposed.

  Similarly, in Patent Document 2, the yield stress is 460 MPa or more by increasing the amount of Mn added and applying a non-refining process including pre-cooling and post-cooling in addition to optimization of chemical components. In addition, a technique relating to a high-strength steel that is excellent in the strength and toughness of the base metal and also in the toughness of the welded portion and a manufacturing method thereof has been proposed.

On the other hand, Patent Document 3 proposes a technique related to a non-heat treated 60 kg grade structural steel having excellent low temperature toughness after cold bending and excellent toughness after strain aging. This technique is applied with a process of cooling to a temperature range of 300 to 600 ° C. at a cooling rate of 2 ° C./second or more from Ar ₃ or higher after rolling (a process in which accelerated cooling is stopped halfway to room temperature) after rolling. The size of cementite is reduced by rolling down in the slab heating temperature and the recrystallization region to refine the prior austenite crystal grain size and bainite packet size and to promote ferrite precipitation by ensuring the cumulative reduction ratio in the non-recrystallization temperature region. In addition, the amount of precipitation is reduced. As a result, it has been shown that excellent toughness can be secured even after strain aging.

JP 2006-241556 A JP 2009-263777 A JP 2001-64723 A

  As described above, various non-tempered manufacturing methods have been proposed from the viewpoint of omitting the tempering treatment. However, it was not made from the viewpoint of securing excellent toughness after cold bending and securing good low temperature toughness for use in cold regions of about -20 to -50 ° C. In Patent Document 1, since rolling is performed at a high temperature to reduce acoustic anisotropy in the manufacturing process, the base material toughness that can be achieved particularly with a thick material having a plate thickness of 80 mm or more is -50 in terms of vTrs. It is about -60 degreeC, and when the toughness ensuring after cold bending process and the use in a cold region are considered, it is thought that the further examination is required.

  Moreover, in patent document 2, since the latter stage cooling is implemented after the former stage cooling, and the stop temperature of the latter stage cooling is not more than 450 ° C. and about 300 ° C., it is tempered in the slow cooling process after the completion of the cooling. It seems that the effect is small and the hardness of the steel plate surface portion is large. As a result, for example, when a square steel pipe having a severe bending inner radius of 2.5 t is manufactured, even if bending can be performed, it seems difficult to prevent cracking of the surface layer portion of the bending.

  On the other hand, Patent Document 3 is intended to improve the toughness after strain aging assuming cold bending or the like, but the assumed strain amount is about 5% (about 10 t in bending inner radius). In a severe cold bending process such as the above-described inner bending radius of 2.5 t, the strain amount on the outer surface of the bending is about 20%, so it is difficult to ensure low temperature toughness after bending. It is done. Actually, in the example of the patent document 3 (example of the present invention), there is an example in which vTs (vTrs) before strain aging is about −70 ° C. When the base material toughness is at this level, It seems difficult to ensure toughness after bending at 2.5 t.

  The present invention has been made in view of the above circumstances, and its purpose is to provide high yield strength and high tensile strength even when the plate thickness is large, and toughness of the base material, bending workability, and after bending work. Establish a technology that provides high-tensile steel sheets with excellent toughness and high HAZ toughness and weldability (weld crack resistance) with high productivity and low cost without the need for off-line heat treatment There is.

A steel sheet excellent in bending workability, impact properties and tensile properties of the present invention that has solved the above problems is
C: 0.02 to 0.05% (meaning “mass%”; the same applies to chemical components),
Si: 0.10 to 0.40%,
Mn: 1.85 to 2.50%,
P: 0.012% or less (excluding 0%),
S: 0.005% or less (excluding 0%),
Nb: 0.020 to 0.050%,
Ti: 0.005-0.020%,
N: 0.0020 to 0.0060%, and Al: 0.010 to 0.060%
And the balance consists of iron and inevitable impurities, and
Weld crack susceptibility composition P _CM, as defined by the following formula (1) is not more than 0.20%, and,
Acicular ferrite fraction in the entire structure of steel: 70% by area or more,
The average crystal grain size (equivalent circle diameter) of the entire structure is 7 μm or less, and the MA (Martensite-Austenite Constituent) fraction: 0.5 area% or less is satisfied. It is characterized by being 220 or less.

P _CM = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B (1)
[In Formula (1), C, Si, Mn, Cu, Ni, Cr, Mo, V, and B show the content (mass%) in steel of each element. ]
The steel sheet further includes at least one element selected from the group consisting of Cu: 0.50% or less (not including 0%) and Ni: 0.50% or less (not including 0%) as other elements. An element may be contained.

  Moreover, the said steel plate may contain Ca: 0.0005 to 0.0050% as another element.

This invention also includes the manufacturing method of the said steel plate, Comprising: This manufacturing method heats the steel piece which has the said component composition to 1050-1200 degreeC, and then the temperature range whose surface temperature is 900-1050 degreeC. And after the hot rolling so that the cumulative rolling reduction is 30% or more in the temperature range of 750 to 850 ° C., the surface temperature is from the temperature of Ar ₃ or more. It is characterized in that it is cooled to a temperature range of 450 to 600 ° C. at an average cooling rate of 4 to 100 ° C./s and then air-cooled.

  According to the present invention, even when the plate thickness is as thick as 80 mm or more, it exhibits high tensile properties (yield strength (YS) is 500 MPa or more and tensile strength (TS) is 570 MPa or more), and base material toughness, High-tensile steel sheet with excellent bending workability and toughness after bending work, and also excellent HAZ toughness and weldability (weld crack resistance), with high productivity and low cost without the need for offline heat treatment Can be provided. The steel plate of the present invention having the above characteristics can be used as a welded structural member such as a bridge, a ship, an offshore structure, a pressure vessel, or a line pipe.

FIG. 1 is a graph showing the relationship between the fraction of acicular ferrite and the yield strength. FIG. 2 is a graph showing the relationship between the fraction of acicular ferrite and the tensile strength. FIG. 3 is a graph showing the relationship between the average crystal grain size of all structures and the yield strength. FIG. 4 is a graph showing the relationship between the MA fraction and the yield strength. FIG. 5 is a graph showing the relationship between the average crystal grain size of all structures and vTrs (impact characteristics).

  In view of the above circumstances, the inventors of the present invention have high yield strength and tensile strength of the base material even when the plate is thick (thick), and excellent base material toughness, bending workability, bending work. The present inventors have earnestly studied methods for obtaining a steel sheet excellent in later toughness, HAZ toughness and weldability (weld crack resistance).

  As a result, in order to secure the above characteristics by applying a process that stops accelerated cooling halfway to room temperature for thick materials that cannot increase the cooling rate inside the steel plate, low carbon is used as the chemical component. In addition, after the ferrite nose is made longer by adding Nb, an austenite stabilizing element (Mn, and further Ni, if necessary) is added to lower the transformation temperature. Appropriately subjected to recrystallization rolling, and in the above process, the cooling start temperature, cooling rate and cooling stop temperature of accelerated cooling are controlled within a predetermined range, and the structure is made to be a fine acicular ferrite-based structure, In addition, an MA (Martensite-Austenite Constituent) organization (hereinafter referred to as “MA”), which is a hard phase in which C is locally concentrated. ) That a very small and found to be important (In the following, the present invention, a process to stop in the middle of the room temperature accelerated cooling is sometimes referred to as "accelerated cooling process of the present invention").

  In addition, even when severe cold bending is performed such that the bending inner radius becomes 2.5 t, in order to ensure good bending workability without causing surface cracks, it is necessary to reduce the surface hardness of the steel sheet. Since it was effective, the method was examined, and it was confirmed that the chemical composition should be low carbon and the maximum hardness should be kept low, and that the accelerated cooling stop temperature should be set to a relatively high temperature to effectively utilize the tempering effect. I found it.

In addition, by suppressing the weld cracking susceptibility composition (P _CM ) to 0.20% or less in the component composition of the steel material, weld cracking is also suppressed and the weldability is excellent, and even with a large heat input such as 15 kJ / mm. A steel plate having high toughness (HAZ toughness) of the weld heat affected zone can be obtained.

  Hereinafter, the steel sheet of the present invention will be described in detail. First, the reason for defining the component composition of the steel sheet of the present invention will be described.

[C: 0.02 to 0.05%]
C has an effect of increasing the strength of the steel sheet. When the C content is less than 0.02%, an acicular ferrite structure is not sufficiently obtained, and it is difficult to secure a necessary base material strength. Therefore, the C content is set to 0.02% or more in the present invention. Preferably it is 0.03% or more.

  On the other hand, C is an element that deteriorates the HAZ toughness and is also an element that easily deteriorates the weld crack resistance. On the other hand, when the C content exceeds 0.05%, the strength of the base material is easily secured, but the hardness sensitivity to the cooling rate increases. As a result, when the cooling rate is increased in the accelerated cooling process of the present invention, the hardness of the steel plate surface portion is increased and bending workability is deteriorated. Furthermore, if the C content is excessive, MA tends to remain after the accelerated cooling process of the present invention, and it becomes difficult to obtain a yield strength of 500 MPa or more. Therefore, in the present invention, the upper limit of the C amount is set to 0.05%. The C content is preferably 0.04% or less.

[Si: 0.10 to 0.40%]
Si is an element effective as a deoxidizing material. It is also an element that is effective in improving the strength of the base metal by securing an acicular ferrite structure. In order to exhibit such a strengthening mechanism, it is necessary to contain 0.10% or more of Si. Preferably it is 0.15% or more. However, when the Si content is excessive, the base material toughness and the toughness (impact characteristics) after bending are likely to deteriorate. Further, if the Si content is excessive, the HAZ toughness and weldability are liable to be deteriorated, so the content is made 0.40% or less. A preferable upper limit is 0.35%.

[Mn: 1.85 to 2.50%]
Mn is an element that is effective in improving the hardenability and improving the strength by stabilizing the austenite and lowering the transformation temperature, and also effective in securing the impact characteristics due to the effect of grain refinement due to the low-temperature transformation. is there. In addition, the acicular ferrite structure in the present invention can be secured at a lower cost than the addition of elements such as Cu and Ni. In order to exert such effects, it is necessary to contain 1.85% or more of Mn. Preferably it is 1.90% or more. However, if Mn is contained excessively, the HAZ toughness deteriorates, so the upper limit of the Mn content is 2.50%. A preferable upper limit is 2.40%.

[P: 0.012% or less (excluding 0%)]
P, an unavoidable impurity, is an element that adversely affects impact properties and HAZ toughness, so the P content needs to be suppressed to 0.012% or less. Preferably it is 0.010% or less.

[S: 0.005% or less (excluding 0%)]
S forms MnS and deteriorates impact characteristics (base material toughness, toughness after bending) and HAZ toughness, so it is preferably as small as possible. From such a viewpoint, the S content needs to be 0.005% or less. Preferably it is 0.003% or less.

[Nb: 0.020 to 0.050%]
Nb is an element effective for forming a non-recrystallized region in the low temperature region of austenite, and by rolling in this low-temperature non-recrystallized region, the microstructure of the base material is made finer and the toughness is increased. Can do. Further, it is an element effective for increasing the strength of the base material by realizing precipitation strengthening after the accelerated cooling process of the present invention described later. Further, in the present invention, Nb is “low C and high Mn as described above, and a predetermined amount of Nb is added, and in the production process, Nb is heated to a temperature at which Nb is solid-solved. It is an indispensable element for obtaining an acicular ferrite structure by applying a reduction. In addition, the solute Nb has the effect of delaying the ferrite transformation in the continuous cooling transformation of steel (making the ferrite nose longer), which suppresses the formation of polygonal ferrite and contributes to increasing the strength of the base material. To do. In order to exhibit these effects, it is necessary to contain 0.020% or more of Nb. Preferably it is 0.030% or more. However, if the Nb content becomes excessive, the HAZ toughness deteriorates, so it is necessary to make it 0.050% or less. A preferable upper limit is 0.040%.

[Ti: 0.005 to 0.020%]
Ti forms N and nitride (TiN) to prevent coarsening of austenite grains (γ grains) during heating before hot rolling, and ensures high yield strength by refining the resulting structure , And an element that contributes to improvement of impact characteristics and HAZ toughness. Furthermore, by securing N and securing solid solution Nb, it is effective to secure the austenite non-recrystallized region and to enhance the yield strength by strengthening precipitation after the accelerated cooling process of the present invention in the manufacturing process. Element. In order to exhibit these effects, it is necessary to contain Ti 0.005% or more. Preferably it is 0.010% or more. However, when the Ti content is excessive, TiC is precipitated in addition to TiN, and impact characteristics and HAZ toughness are deteriorated. Therefore, the Ti content is 0.020% or less. Preferably it is 0.018% or less.

[N: 0.0020 to 0.0060%]
N is an element effective for producing TiN together with Ti, preventing coarsening of γ grains during heating and hot welding before hot rolling, and improving impact properties and HAZ toughness. When the N content is less than 0.0020%, TiN becomes insufficient, the γ grains become coarse, and impact characteristics and HAZ toughness deteriorate. Therefore, in this invention, it is necessary to make N amount 0.0020% or more. Preferably it is 0.0025% or more. On the other hand, if the N content becomes excessive and exceeds 0.0060%, impact characteristics and HAZ toughness are deteriorated. Therefore, in the present invention, the upper limit of the N amount is set to 0.0060%. A preferred upper limit is 0.0055%.

[Al: 0.010 to 0.060%]
Since Al is an element necessary for deoxidation, 0.010% or more is contained. Preferably it is 0.020% or more, More preferably, it is 0.030% or more. On the other hand, if Al is excessively contained, alumina-based coarse inclusions are formed and the impact characteristics are deteriorated, so the content is made 0.060% or less. Preferably it is 0.050% or less.

  The components of the steel sheet of the present invention are as described above, and the balance is composed of iron and inevitable impurities. Moreover, in addition to the above elements, the following elements can be further contained, and the strength, toughness, and the like can be further increased by containing appropriate amounts of these elements. Hereinafter, these elements will be described in detail.

[One or more elements selected from the group consisting of Cu: 0.50% or less (excluding 0%) and Ni: 0.50% or less (not including 0%)]
Cu and Ni are both effective elements for improving the strength and toughness of the base material without significantly affecting the weldability and HAZ toughness. In order to exert these effects, it is preferable to contain 0.10% or more (more preferably 0.15% or more) of Cu and Ni. An object of the present invention is to reduce the addition amount of expensive Cu and Ni as much as possible by securing an inexpensive addition amount of Mn. Therefore, the upper limit of the content of these elements is not limited metallurgically, but is preferably 0.50% or less from the viewpoint of reducing raw material costs. More preferably, it is 0.45% or less.

[Ca: 0.0005 to 0.0050%]
Ca is an element that effectively acts to detoxify the weld crack resistance by spheroidizing MnS. In order to exert such an effect, it is preferable to contain 0.0005% or more of Ca. More preferably, it is 0.0010% or more. However, when the Ca content is excessive, inclusions are coarsened and the base material toughness is deteriorated. Therefore, it is preferable that the upper limit of the Ca content is 0.0050%. A more preferred upper limit is 0.0040%.

In the present invention, a weld cracking susceptibility composition (P _CM ) defined by the following formula (1) is specified.

[P _{CM represented by the} following formula (1): 0.20% or less]
P _CM = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B (1)
[In Formula (1), C, Si, Mn, Cu, Ni, Cr, Mo, V, and B show the content (mass%) in steel of each element. ]
P _CM is referred to as weld crack susceptibility composition, even in the steel sheet degree of restraint is large thickness, for example, 100mm and the thick, in order to suppress weld cracking stably, it is necessary to 0.20% or less. P _CM is preferably less 0.19%.

Incidentally, preferably as the value of P _CM is small, particularly lower limit is not, the chemical component composition of the present invention, the lower limit of the P _CM becomes approximately about 0.14%. In the present invention, elements not included in the above formula (1) were calculated with the content as zero.

  Next, the reason why the steel structure (microstructure) is limited in the present invention will be described.

  In the present invention, in order to ensure desired properties (particularly high yield strength and tensile strength, and excellent impact properties), the fraction of acicular ferrite in the entire structure of steel is 70 area% or more, and the entire structure The average crystal grain size (equivalent circle diameter) must be 7 μm or less, and the MA fraction should be 0.5 area% or less.

  In addition, as shown in the Example mentioned later, the site | part which prescribes | regulates a structure | tissue is 1/4 thickness part. The said site | part is a site | part generally used for evaluating the mechanical property of a base material, and the structure | tissue in the site | part was prescribed | regulated.

  The reason for the above definition will be described below.

[Fraction of acicular ferrite in the entire structure of steel: 70 area% or more]
In the present invention, in order to ensure the tensile properties and the base material toughness of the base material, in addition to the optimization of the component composition (low C + high Mn + Nb, Ti addition) and the optimization of heating and rolling conditions, the accelerated cooling of the present invention By adopting the process, steel transformation strengthening and Nb precipitation strengthening are utilized. However, when transformation starts at the highest temperature among the transformation structures of steel, mainly diffusion transformation and soft polygonal ferrite increases, it becomes difficult to satisfy tensile properties, particularly yield strength of 500 MPa or more. Therefore, it is necessary to make the main structure an acicular ferrite that is a structure that transforms at a lower temperature than polygonal ferrite and that is effective in securing tensile properties and impact properties.

  Specifically, the acicular ferrite needs to be 70 area% or more with respect to the entire structure of the steel. When the fraction of the acicular ferrite is less than 70 area%, that is, when the polygonal ferrite structure is increased as a structure other than the acicular ferrite, it becomes difficult to ensure the strength of the base material as described above. Further, an increase in the fraction of bainite structure or martensite structure is not preferable because it becomes difficult to ensure the base material toughness.

  The fraction of acicular ferrite is more preferably 80 area% or more. The higher the fraction of acicular ferrite, the better and there is no particular upper limit.

  Although the definition of the acicular ferrite is unclear, the acicular ferrite in the present invention includes pseudopolygonal ferrite and granular bainitic ferrite. On the other hand, a structure in which no pseudopolygonal ferrite was present and the old γ grain boundary was clearly preserved was distinguished as a bainite structure.

  Examples of the structure other than the acicular ferrite include polygonal ferrite, bainite, and martensite that are inevitably formed in the manufacturing process. From the viewpoint of obtaining more excellent characteristics, the polygonal ferrite is preferably suppressed to 20 area% or less, more preferably 10 area% or less.

[Average crystal grain size of all structures (equivalent circle diameter): 7 μm or less]
In the present invention, in order to ensure excellent toughness after bending, it is necessary to increase the toughness (particularly, low temperature toughness) of the base material (vTrs ≦ −85 ° C.). For this purpose, as described above, the steel structure must be mainly composed of acicular ferrite, and the average crystal grain size of the entire structure must be 7 μm or less in terms of the equivalent circle diameter. When the average crystal grain size exceeds 7 μm, it is difficult to ensure the low temperature toughness of the base material even in a structure mainly composed of acicular ferrite. Further, when the structure becomes coarse, the effect of increasing the yield strength due to the refinement of the structure becomes small, and it becomes difficult to satisfy the yield strength of 500 MPa or more. The average crystal grain size of the entire structure is preferably 6 μm or less.

[MA fraction: 0.5 area% or less]
The present invention is characterized by ensuring a high tensile strength and achieving a high yield strength. For this purpose, the MA fraction must be 0.5 area% or less. When the fraction of MA exceeds 0.5 area%, the yield strength decreases due to the yield ratio reduction effect of hard MA, and high yield strength cannot be achieved. The fraction of MA is preferably 0.3 area% or less.

[Maximum Vickers hardness of steel plate surface: 220 or less]
Furthermore, in the present invention, it is necessary to reduce the hardness of the surface portion of the steel sheet in order to obtain a high-tensile steel sheet having excellent bending workability. Specifically, it is necessary to suppress the maximum value of the Vickers hardness of the steel plate surface portion (a position at a depth of 1 mm from the steel plate surface as shown in Examples described later) to 220 or less. If the maximum value of the Vickers hardness exceeds 220, cracks may occur in the steel plate surface when a severe cold bending process such as a bending inner radius of 2.5 t is performed. The maximum value of the Vickers hardness is more preferably 215 or less.

  Next, the manufacturing method of the steel plate of this invention is demonstrated.

In the present invention, the steel slab having the chemical composition described above is heated to 1050 to 1200 ° C., and then the cumulative rolling reduction is 30% or more in the temperature range of the surface temperature of 900 to 1050 ° C., and the surface temperature is After hot rolling so that the cumulative rolling reduction is 30% or more in the temperature range of 750 to 850 ° C., the surface temperature is 450 ° C./s at an average cooling rate of 4 to 100 ° C./s from the temperature of Ar ₃ or more. Cool to a temperature range of 600 ° C., then air cool. The reason for the above definition will be described below.

[Heating temperature of steel slab during hot rolling: 1050 to 1200 ° C.]
This heating temperature greatly affects the structure control before hot rolling. Even if a prescribed amount of Nb is contained, if the heating temperature is less than 1050 ° C., the solid solution of Nb becomes insufficient, the recrystallization suppression effect by the solid solution Nb is reduced, and the effect of refining the structure is reduced. In addition, if the amount of dissolved Nb is small, the effect of suppressing the formation of polygonal ferrite by delaying the ferrite transformation during the continuous cooling transformation during the accelerated cooling, or after the acceleration cooling midway stop in the accelerated cooling process of the present invention The effect of precipitation strengthening is reduced, making it difficult to ensure excellent tensile properties. Therefore, in this invention, heating temperature was 1050 degreeC or more. Preferably it is 1080 degreeC or more. On the other hand, when the heating temperature exceeds 1200 ° C., the austenite (γ) grain size becomes coarse, so that the impact characteristics are deteriorated and a desired yield strength cannot be ensured. Therefore, in this invention, the upper limit of heating temperature shall be 1200 degreeC. More preferably, it is 1180 degrees C or less.

[Cumulative rolling reduction at a surface temperature of 900 to 1050 ° C .: 30% or more]
The temperature range where the surface temperature is 900 to 1050 ° C. is a temperature range where austenite is recrystallized during hot rolling even when the solid solution Nb amount is sufficiently secured. In order to ensure excellent base metal toughness (and toughness after bending) and desired yield strength, the cumulative rolling reduction in this temperature range must be 30% or more, and austenite grains must be recrystallized repeatedly to refine them. There is. If the cumulative rolling reduction is less than 30%, the austenite grains immediately after the heating cannot be refined, and as a result, the final structure becomes coarse and it is difficult to ensure the above characteristics. A preferable cumulative rolling reduction in this temperature range is 40% or more.

  Further, the upper limit of the cumulative rolling reduction is not particularly limited from the viewpoint of the above-mentioned miniaturization, but is about 80% from the viewpoint of the productivity of the rolling process and the total rolling reduction ratio.

[Cumulative rolling reduction at a surface temperature of 750 to 850 ° C .: 30% or more]
The temperature range where the surface temperature is 750 to 850 ° C. is a so-called non-recrystallized region in which austenite does not recrystallize during hot rolling as long as the solid solution Nb amount is sufficiently secured. In order to ensure excellent impact properties and desired yield strength, the austenite grains are repeatedly refined by recrystallization by hot rolling in the above recrystallization temperature region, and the cumulative reduction ratio is further increased in this non-recrystallization region. It is necessary to secure 30% or more. Thereby, strain can be accumulated in austenite, transformation nuclei in the accelerated cooling step after hot rolling can be increased, and the final structure after transformation can be refined. If the cumulative rolling reduction in this temperature range is less than 30%, the transformation nuclei are insufficient, the final structure becomes coarse, and it becomes difficult to ensure the above characteristics. A preferable cumulative rolling reduction in this temperature range is 40% or more.

  Further, the upper limit of the cumulative rolling reduction is not particularly limited from the viewpoint of the above-mentioned miniaturization, but is about 80% from the viewpoint of the productivity of the rolling process and the total rolling reduction ratio.

[Accelerated cooling start temperature (cooling start temperature): Ar ₃ or higher temperature]
When the surface temperature is lower than Ar ₃ point, soft polygonal ferrite is generated, which causes a decrease in the strength of the base material. Therefore, accelerated cooling needs to start from a temperature of Ar ₃ or higher. The cooling start temperature of accelerated cooling is preferably a temperature of (Ar ₃ points + 20 ° C.) or higher. The upper limit of the cooling start temperature for accelerated cooling is about 800 ° C.

The Ar ₃ was calculated by the following formula (2). In the following formula (2), elements not contained in steel were calculated as zero.
Ar ₃ (° C.) = 910-310 × C-80 × Mn-20 × Cu-15 × Cr-55 × Ni-80 × Mo (2)
[In Formula (2), C, Mn, Cu, Cr, Ni, and Mo show content (mass%) in steel of each element. ]

[Average cooling rate of accelerated cooling: 4 to 100 ° C./s]
In order to secure sufficient acicular ferrite and ensure high tensile properties, the average cooling rate of accelerated cooling needs to be 4 ° C./s or more. When the average cooling rate is less than 4 ° C./s, for example, at a slow cooling rate such as air cooling, the acicular ferrite fraction decreases and polygonal ferrite increases, so that the strength of the base material cannot be secured. . The average cooling rate of accelerated cooling is preferably 8 ° C./s or more. On the other hand, when the average cooling rate exceeds 100 ° C./s, the surface portion is mainly composed of martensite due to shear transformation, and the surface hardness is increased. Therefore, the upper limit of the average cooling rate is set to 100 ° C./s. Preferably it is 80 degrees C / s or less.

[Accelerated cooling stop temperature (cooling stop temperature): 450 to 600 ° C temperature range]
In the accelerated cooling process, in order to increase the strength by transformation strengthening and precipitation strengthening, it is necessary to stop the accelerated cooling in a relatively high temperature range of 450 to 600 ° C. If the temperature is lower than 450 ° C., transformation strengthening can be obtained, but the tempering effect due to the intermediate stoppage is reduced, and the surface hardness is increased, and MA remains to cause a decrease in yield strength. Therefore, the stop temperature of accelerated cooling was set to 450 ° C. or higher. Preferably it is 470 degreeC or more. On the other hand, when the temperature exceeds 600 ° C., sufficient transformation strengthening cannot be obtained, and a polygonal ferrite main structure is obtained, making it difficult to obtain sufficient base material strength. Therefore, the stop temperature of accelerated cooling is set to 600 ° C. or lower. Preferably it is 570 degrees C or less.

  After the accelerated cooling, the steel sheet of the present invention can be obtained by air cooling to room temperature. In the present invention, as described above, the cooling is stopped in the temperature range of 450 to 600 ° C. by accelerated cooling, and then air cooling is performed to enhance precipitation strengthening by Nb carbonitride.

  EXAMPLES Hereinafter, the present invention will be described more specifically with reference to examples. However, the present invention is not limited by the following examples, but may be appropriately modified within a range that can meet the purpose described above and below. Of course, it is possible to implement them, and they are all included in the technical scope of the present invention.

  (Chemical) component composition shown in Table 1 (the balance is iron and inevitable impurities, and in Table 1, the blank indicates that no element is added). The obtained slab is heated to the temperature shown in Table 2 or 3 (slab heating temperature) and then hot-rolled, and then subjected to accelerated cooling to obtain a steel plate having the thickness shown in Table 2 or Table 3. It was. In some cases, this accelerated cooling was not performed and air cooling was performed.

  The slab heating temperature is an average temperature calculated in the thickness direction at the center of the slab, and is calculated from the in-furnace atmosphere temperature of the heating furnace and the in-furnace time. Moreover, the temperature in hot rolling, the accelerated cooling start temperature, and the accelerated cooling stop temperature are all temperatures measured by a radiation thermometer installed in the line. Here, the accelerated cooling stop temperature is the surface temperature after recuperation after completion of accelerated cooling. The average cooling rate during accelerated cooling is calculated from the steel sheet surface temperature at the start of accelerated cooling, the steel sheet surface temperature at the time of stopping, and the cooling time.

  Using the steel sheet obtained as described above, the observation of the structure and the evaluation of the characteristics were carried out as follows.

<Observation of steel structure>
[Measurement of acicular ferrite fraction]
The acicular ferrite fraction was measured as follows.
(1) A sample is taken from the steel plate so that a plate thickness cross section including the steel plate front and back surfaces parallel to the rolling direction and perpendicular to the steel plate surface can be observed.
(2) Mirror finish of the observation surface is performed by polishing with wet emery polishing paper (# 150 to # 1000) or a polishing method having the same function (polishing using an abrasive such as diamond slurry).
(3) The polished sample is corroded using a 3% nital solution to reveal grain boundaries.
(4) At t (plate thickness) / 4 site, the exposed tissue is photographed at a magnification of 400 times (in this example, photographed as a 6 cm × 8 cm photograph). Next, in the photograph taken, it is discriminated whether polygonal ferrite is generated at the prior austenite grain boundary and is painted black.

  On the other hand, when polygonal ferrite is not formed and the prior austenite grain boundaries clearly remain, the structure of the region surrounded by the austenite grain boundaries is the bainite structure or martensite structure mainly composed of shear transformation. Judge and paint black.

  As the case where the bainite structure or the martensite structure is generated, when the specified component composition is satisfied in the present invention, in the manufacturing process, when the speed during accelerated cooling is extremely large, or processing such as hot rolling is not added. In the case of accelerated cooling, or when the processing rate in hot rolling is small. Moreover, when not satisfy | filling the component composition prescribed | regulated by this invention, when containing B, the addition amount of C increases, and the hardenability is high, ie, the case where a transformation temperature falls further, etc. are mentioned.

  In addition, since MA cannot be discriminated by the above corrosion, it is separately measured by a method described later.

Next, the photograph is taken into an image analysis apparatus (the area of the photograph corresponds to 150 μm × 200 μm when the magnification is 400 times). The image analysis apparatus captures the image so that the total area is 1 mm × 1 mm or more at any magnification (that is, at least 35 images are captured when the magnification is 400 times).
(5) In the image analysis apparatus, the area ratio other than black was calculated for each photograph, and the fraction obtained by subtracting the MA fraction described later was defined as the acicular ferrite fraction. In Tables 4 and 5, the black-colored polygonal ferrite and the fraction of bainite and / or martensite are also shown for reference.

[Measurement of average crystal grain size (equivalent circle diameter) of all structures]
The average crystal grain size (equivalent circle diameter) of the entire structure was measured as follows.
(1) A sample including front and back portions having a plate thickness cut in a direction parallel to the rolling direction is prepared.
(2) Using a wet emery polishing paper of # 150 to # 1000 or a polishing method having a function equivalent to that, a mirror finish is applied using an abrasive such as abrasive paper or diamond slurry.
(3) Using an EBSP (Electron Back Scattering Pattern) device manufactured by TexSEM Laboratories, measuring range: 200 × 200 μm, 0.5 μm pitch, crystal orientation difference of 15 ° or more at t / 4 part in the plate thickness direction The size of the large-angle grain is measured with the boundary as the grain boundary. At this time, measurement points whose confidence index indicating the reliability of the measurement direction is smaller than 0.1 are excluded from the analysis target.
(4) The average value of the sizes surrounded by the large-angle grain boundaries obtained in this way is calculated and set as the “average crystal grain size of all structures” in the present invention. In addition, about the thing surrounded by a large-angle grain boundary and whose size is 1.0 micrometer or less, it is judged as measurement noise and is excluded from the object of average value calculation.

[Method for observing MA and measuring fraction]
The MA fraction was measured as follows.
(1) A sample is taken from the steel plate so that a plate thickness cross section including the steel plate front and back surfaces parallel to the rolling direction and perpendicular to the steel plate surface can be observed.
(2) Mirror finish of the observation surface is performed by polishing with wet emery polishing paper (# 150 to # 1000) or a polishing method having the same function (polishing using an abrasive such as diamond slurry).
(3) The polished sample is corroded with a repeller solution to reveal MA. The MA appearing portion is colored white on the optical micrograph. In addition, since martensite does not become white by this corrosion, martensite and MA can be distinguished.
(4) At t (plate thickness) / 4 site, the exposed tissue is photographed at a magnification of 1000 times (in this example, photographed as a 6 cm × 8 cm photograph). Next, the photograph is taken into the image analysis apparatus (the area of the photograph corresponds to 60 μm × 80 μm when the magnification is 1000 times). In the image analysis apparatus, the total area is 0.4 mm × 0.4 mm or more (that is, in the case of 1000 times, at least 35 of the above photographs are captured).
(5) In the image analysis apparatus, the area ratio of MA is calculated for each photograph, and the average value of all photographs is set as the area ratio of MA.

<Evaluation of bending workability (measurement of maximum value of surface Vickers hardness)>
The maximum value (surface hardness) of the surface Vickers hardness was measured as follows.
(1) A sample is taken from the steel plate so that a plate thickness cross section including the steel plate front and back surfaces parallel to the rolling direction and perpendicular to the steel plate surface can be observed.
(2) Mirror finish of the observation surface is performed by polishing with wet emery polishing paper (# 150 to # 1000) or a polishing method having the same function (polishing using an abrasive such as diamond slurry).
(3) In the polished sample, the Vickers hardness is measured with a load of 98 N at 10 points at a 1 mm pitch in the horizontal direction 1 mm below the surface, and the highest of the 10 Vickers hardnesses. Was the maximum value of the Vickers hardness of the steel plate surface. And when this maximum value was 220 or less, it was evaluated that surface hardness was low and it was excellent in bending workability.

<Evaluation of tensile properties>
A No. 4 test piece of JISZ 2201 was taken from the portion of t (plate thickness) / 4 in the direction perpendicular to the rolling direction, and a tensile test was performed according to the procedure of JISZ 2241 to measure the yield strength and the tensile strength. And those having a yield strength of 500 MPa or more and a tensile strength of 570 MPa or more were evaluated as having excellent tensile properties.

<Evaluation of impact characteristics (Charpy impact test)>
A V-nots test piece of JISZ 2242 was taken from the site of t (plate thickness) / 4 in the direction perpendicular to the rolling direction, and a Charpy impact test was performed in accordance with the procedure of JISZ 2242 to obtain vTrs. In addition, when calculating | requiring vTrs, it implemented 3 each at each test temperature. Then, those having a vTrs of −85 ° C. or less were evaluated as having excellent impact characteristics, specifically excellent in base material toughness, and excellent in toughness of a bent portion after bending.

<Evaluation of HAZ toughness>
Using a reproducible thermal cycle tester, a thermal cycle assuming a welding heat input of 15 kJ / mm was applied, a V-notch test piece of JISZ 2242 was taken, and a Charpy impact test was performed in the manner of JISZ 2242 to evaluate HAZ toughness. did. The test temperature was −20 ° C., and the average value of the three samples was determined. And the case where this average value was 100 J or more was evaluated as having excellent HAZ toughness.

<Evaluation of weldability (measurement of crack prevention temperature)>
For the evaluation of the cracking prevention temperature, welding was carried out at 5 ° C, 25 ° C, 50 ° C and 75 ° C in the manner of JISZ 3158 by covering arc welding, and the cracking prevention temperature was measured. Those having a crack prevention temperature of 5 ° C. were evaluated as having excellent weldability.

  These results are shown in Tables 4 and 5.

  It can consider as follows from Tables 1-5 (the following No. shows experiment No. of Tables 2-5).

  No. A1-1, A1-3, A1-4, A2 to A5, A6-3, A6-4, A6-7, A6-8, A7-3, A7-4, A7-6 to A7-8, A8- 4, A9 to A13 satisfy the component composition specified in the present invention and are obtained by manufacturing under the specified conditions, and thus exhibit high tensile properties (yield strength / tensile strength) and toughness of the base material. And bendability, toughness after bending, weldability, and HAZ toughness are all excellent.

  On the other hand, the above-mentioned No. Examples other than the above did not satisfy at least one of the component composition and the production conditions defined in the present invention, and as a result, any of the above characteristics was inferior.

  Specifically, no. In A1-2, the (slab) heating temperature was too low, so Nb was not completely dissolved, the hardenability was insufficient, and an acicular ferrite structure was not obtained. As a result, the tensile properties were inferior.

  No. In A1-5, since the (slab) heating temperature was too high, the austenite (γ) crystal grains were coarsened, and as a result, the average crystal grain size of the entire structure was increased. As a result, the impact characteristics were inferior and the desired yield strength could not be ensured.

  No. No. A6-1 is not subjected to reduction in a temperature range (γ recrystallization temperature range) where the surface temperature of the steel sheet is 900 to 1050 ° C. In A6-2, since the rolling reduction in the above temperature range was insufficient, the average crystal grain size of all the structures became large. As a result, the impact characteristics were inferior and the desired yield strength could not be ensured.

  No. A6-5 is not subjected to reduction in the temperature range (γ non-recrystallization temperature range) where the surface temperature of the steel sheet is 750 to 850 ° C. In A6-6, since the rolling reduction in the above temperature range was insufficient, the average crystal grain size of all the structures became large. As a result, the impact characteristics were inferior and the desired yield strength could not be ensured.

  No. Since A7-1 did not carry out accelerated cooling, an acicular ferrite structure was not sufficiently obtained, became a polygonal ferrite structure main body, and was inferior in tensile properties.

No. In A7-2, since the cooling start temperature in accelerated cooling was lower than Ar ₃ , polygonal ferrite was precipitated, and the acicular ferrite structure was not sufficiently obtained. As a result, the tensile properties were inferior.

  No. Since A7-5 had a slow cooling rate in accelerated cooling, an acicular ferrite structure was not sufficiently obtained, and it became a polygonal ferrite structure main body, resulting in inferior tensile properties.

  No. In A7-9, since the cooling rate in the accelerated cooling was too fast, the surface hardness was too large and the bending workability was poor.

  No. In A8-1 and A8-2, since the cooling stop temperature in the accelerated cooling was too low, MA was generated, and the desired yield strength could not be secured. Moreover, the surface hardness was too large, and the bending workability was poor.

  No. In A8-5, since the cooling stop temperature in accelerated cooling was too high, an acicular ferrite structure could not be obtained sufficiently, and it became a polygonal ferrite structure mainly, resulting in inferior tensile properties.

No. B1, since the P _CM exceeds the upper limit of the prescribed resistance weld cracking resistance is deteriorated.

  No. As for B2, since the amount of C is insufficient, an acicular ferrite structure cannot be obtained sufficiently, resulting in inferior tensile properties.

  No. In B3, since the amount of C is excessive, MA is excessively generated, and a desired yield strength cannot be ensured. Moreover, the surface hardness became too large and the bending workability was poor. Furthermore, a fine structure was not obtained, resulting in poor impact characteristics. The HAZ toughness was also deteriorated.

  No. In B4, since the amount of Si was insufficient, an acicular ferrite structure was not sufficiently obtained, resulting in inferior tensile properties. No. Since B5 has an excessive amount of Si, impact characteristics and HAZ toughness deteriorated.

  No. In B6, since the amount of Mn was insufficient, an acicular ferrite structure was not sufficiently obtained, and the tensile properties were deteriorated. In addition, the average grain size of the entire structure was increased, resulting in poor impact characteristics. No. B7 resulted in deterioration of HAZ toughness because of the excessive amount of Mn.

  No. B8 resulted in inferior impact properties and HAZ toughness due to an excessive amount of P.

  No. B9 resulted in inferior hitting characteristics and HAZ toughness because the amount of S was excessive.

  No. B10 resulted in inferior impact characteristics and HAZ toughness due to the excessive amount of Al.

  No. In B11, since the amount of Nb is insufficient, a sufficient acicular ferrite structure is not obtained, the average crystal grain size of the entire structure is increased, the tensile properties are deteriorated, and the impact properties are also deteriorated. No. Since B12 has an excessive amount of Nb, HAZ toughness deteriorated.

  No. Since B13 has insufficient Ti amount, TiN is not sufficiently formed, austenite grains are coarsened by heating before hot rolling, the desired yield strength cannot be obtained, and impact properties and HAZ toughness are also deteriorated. As a result. No. Since B14 has an excessive amount of Ti, TiC was precipitated, and impact characteristics (base material toughness, toughness after bending) and HAZ toughness deteriorated.

  No. In B15, since the amount of N is insufficient, TiN is not sufficiently formed, austenite grains are coarsened by heating before hot rolling, the average crystal grain size of the entire structure is increased, and a desired yield strength is obtained. In addition, impact characteristics and HAZ toughness deteriorated. No. In B16, since the N amount is excessive, the impact characteristics and the HAZ toughness deteriorated.

  Moreover, the figure which arranged the relationship between the structure | tissue and the characteristic using the Example is shown in FIGS. FIG. 1 is a graph showing the relationship between the fraction of acicular ferrite and the yield strength, and FIG. 2 is a graph showing the relationship between the fraction of acicular ferrite and tensile strength. From FIG. 1 and FIG. 2, it can be seen that in order to obtain a yield strength of 500 MPa or more and a tensile strength of 570 MPa or more, the acicular ferrite fraction must be 70 area% or more. FIG. 3 is a graph showing the relationship between the average crystal grain size of all the structures and the yield strength, and FIG. 4 is a graph showing the relationship between the MA fraction and the yield strength. From these FIG. 3 and FIG. 4, in order to make the yield strength 500 MPa or more, it is necessary to make the average crystal grain size of the entire structure 7 μm or less and to suppress the MA fraction to 0.5 area% or less. Recognize.

  Further, FIG. 5 is a graph showing the relationship between the average crystal grain size of all structures and vTrs (impact characteristics). From FIG. 5, it can be seen that in order to achieve vTrs: −85 ° C. or lower, the average crystal grain size of the entire structure must be 7 μm or less.

Claims (4)

C: 0.02 to 0.05% (meaning “mass%”; the same applies to chemical components),
Si: 0.10 to 0.40%,
Mn: 1.85 to 2.50%,
P: 0.012% or less (excluding 0%),
S: 0.005% or less (excluding 0%),
Nb: 0.020 to 0.050%,
Ti: 0.005-0.020%,
N: 0.0020 to 0.0060%, and Al: 0.010 to 0.060%
And the balance consists of iron and inevitable impurities, and
Weld crack susceptibility composition P _CM, as defined by the following formula (1) is not more than 0.20%, and,
Acicular ferrite fraction in the entire structure of steel: 70% by area or more,
The average crystal grain size (equivalent circle diameter) of the whole structure: 7 μm or less, and the fraction of MA (Martensite-Austenite Constituent): 0.5 area% or less,
Furthermore, a steel sheet excellent in bending workability, impact characteristics and tensile characteristics, characterized in that the maximum value of Vickers hardness of the steel sheet surface portion is 220 or less.
P _CM = C + Si / 30 + Mn / 20 + Cu / 20 + Ni / 60 + Cr / 20 + Mo / 15 + V / 10 + 5 × B (1)
[In Formula (1), C, Si, Mn, Cu, Ni, Cr, Mo, V, and B show the content (mass%) in steel of each element. ]   Further, as other elements, one or more elements selected from the group consisting of Cu: 0.50% or less (not including 0%) and Ni: 0.50% or less (not including 0%) are contained. The steel plate according to claim 1.   Furthermore, the steel plate of Claim 1 or 2 containing Ca: 0.0005 to 0.0050% as another element. A method for producing the steel sheet according to claim 1,
The steel slab having the component composition according to any one of claims 1 to 3 is heated to 1050 to 1200 ° C, and then the cumulative rolling reduction is 30% or more in a temperature range where the surface temperature is 900 to 1050 ° C, and After performing hot rolling so that the cumulative rolling reduction is 30% or more in the temperature range of 750 to 850 ° C., the average cooling rate of 4 to 100 ° C./s from the temperature of the surface temperature of Ar ₃ or more A method for producing a steel sheet excellent in bending workability, impact properties and tensile properties, characterized by cooling to a temperature range of 450 to 600 ° C. and then air cooling.

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