JP2009263685A - High strength steel sheet having reduced deterioration in characteristic after cutting, and method for producing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high strength steel sheet having excellent ductility equal to that of a DP steel and excellent hole expansibility equal to that of the one with a single structure, and in which the damage of edge faces after cutting is extremely slight, and to provide a method for producing the same. <P>SOLUTION: The high strength steel sheet having reduced deterioration in characteristics after cutting is composed of a steel, which comprises, by mass, 0.05 to 0.20% C, 0.3 to 2.00% Si, 1.3 to 2.6% Mn, 0.001 to 0.03% P, 0.0001 to 0.01% S, <0.10% Al, 0.0005 to 0.0100% N and 0.0005 to 0.007% O, and the balance iron with inevitable impurities, has a steel sheet structure mainly composed of ferrite and bainite, and in which the Mn segregation degree in the sheet thickness direction (= the central part Mn peak concentration/the average Mn concentration) is ≤1.20, and the maximum tensile strength is ≥540 MPa. <P>COPYRIGHT: (C)2010,JPO&INPIT

Description

  The present invention relates to a high-strength steel sheet with little deterioration in characteristics after cutting, which is suitable for automobiles, building materials, home appliances, and the like, and a method for manufacturing the same.

  In recent years, high-strength steel sheets have been applied in the automobile field in order to achieve both a function for protecting passengers in the event of a collision and weight reduction for the purpose of improving fuel efficiency. In particular, with regard to ensuring collision safety, in addition to increasing safety awareness, high-strength steel sheets are applied from strengthening laws and regulations to parts with complex shapes that have only been used with low-strength steel sheets until now. There is a need to try. However, since the formability of the material deteriorates as the strength increases, it is necessary to manufacture a steel sheet that satisfies both formability and high strength when applying a high strength steel sheet to a member having a complex shape. is there. Even if it is said that formability is a bit, when applying it to a member having a complicated shape such as an automobile member, for example, it should have simultaneously formability with different ductility, stretch formability, hole expandability, stretch flangeability, etc. Is required.

In particular, a member for an automobile needs to be spot-welded or the like for joining the members, and the member is often provided with a flange. In many cases, the end face of the flange part is processed as it is cut, and it is particularly easy to break due to the damage caused by cutting. Therefore, it is required that the flange part does not break during processing.
Since the stretch flangeability is obtained by processing the end face as it is cut, it is required that the stretch flangeability is good as a material characteristic, and at the same time, damage to the end face due to mechanical cutting of a shear or the like is slight.
As shown in Non-Patent Document 1, the material properties necessary for improving stretch flangeability are uniform stretch and hole expandability. For this reason, it is required to have both uniform elongation and hole expansibility.
On the other hand, there are many damages that are thought to have dragged inclusions at the time of cutting on the shears and punched end faces, and this is the starting point, and it is known that cracks occur during stretch flange molding and hole expansion tests (non- Patent Document 1). For this reason, it is also extremely important to suppress damage to the end face during cutting.

  It is known that ductility and stretch formability, which are important as formability of a thin steel sheet, are correlated with a work hardening index (n value), and a steel sheet having a high n value is known as a steel sheet having excellent formability. For example, as steel plates excellent in ductility and stretch formability, there are DP (Dual Phase) steel plates whose steel plate structure is composed of ferrite and martensite, and TRIP (Transformation Induced Plasticity) steel plates containing residual austenite in the steel plate structure (Patent Document 1). Patent Document 2). On the other hand, as a steel sheet excellent in hole expansibility, a steel sheet having a ferrite single-phase structure in which the steel sheet structure is precipitation strengthened and a steel sheet having a bainite single-phase structure are known (Patent Documents 3 to 5, Non-Patent Document 2).

  The DP steel sheet has excellent ductility by having ferrite having high ductility as a main phase and dispersing martensite which is a hard structure in the steel sheet structure. Further, soft ferrite is easily deformed, and a large amount of dislocations are introduced and hardened together with the deformation, so that the n value is also high. However, if the steel sheet structure is composed of soft ferrite and hard martensite, the deformability of both structures is different, so in forming with large machining such as hole expansion, there is a minute amount at the interface between both structures. There is a problem that microvoids are formed and the hole expandability is significantly deteriorated. In particular, since the martensite volume fraction contained in the DP steel sheet having a maximum tensile strength of 590 MPa or more is relatively large and there are also many ferrite and martensite interfaces, the microvoids formed at the interface are easily connected, crack formation, It leads to breakage. From this, the hole expansibility of DP steel plate is inferior (for example, nonpatent literature 3).

  Similarly, in the TRIP steel plate whose steel plate structure is composed of ferrite and retained austenite, the hole expandability is low. This is because hole expansion processing and stretch flange processing, which are molding processes for automobile members, are performed after punching or mechanical cutting. The retained austenite contained in the TRIP steel sheet transforms into martensite when subjected to processing. For example, in the case of stretch-stretching or overhanging, it is possible to ensure high formability by increasing the strength of the processed portion and suppressing the concentration of deformation by transforming residual austenite into martensite. However, once punching, cutting, or the like is performed, the vicinity of the end face is subjected to processing, so that residual austenite contained in the steel sheet structure is transformed into martensite. As a result, the structure becomes similar to that of the DP steel sheet, and the hole expandability and stretch flange formability are inferior. Alternatively, since the punching process itself involves a large deformation, after punching, microvoids exist at the interface between ferrite and hard structure (here, martensite transformed with retained austenite), which deteriorates hole expandability. It has been reported that

  Or the steel sheet in which a cementite and a pearlite structure exist in a grain boundary is also inferior in hole expansibility. This is because the boundary between ferrite and cementite is the starting point for microvoid formation.

  As a result, as shown in Patent Documents 3 to 5 and Non-Patent Document 1, the development of a steel sheet excellent in hole expansibility has a single-phase structure of ferrite in which the main phase of the steel sheet is bainite or precipitation strengthened, and grains In order to suppress the formation of cementite phase at the boundary, by adding a large amount of alloy carbide forming elements such as Ti and making C contained in the alloy alloy carbide, a high strength hot-rolled steel sheet with excellent hole expandability can be obtained. Has been developed.

  By the way, the high-strength hot-rolled steel sheet contains a lot of alloying elements. In particular, when utilizing the structure strengthening, a large amount of Mn is added. However, when there is much content of Mn, there exists a tendency for Mn to segregate in the center part of a plate | board thickness direction. This is because when the molten steel is continuously cast into the slab, cooling proceeds from the surface of the slab toward the center, but at this time, Mn does not dissolve to the end, Mn gradually moves to the center of the slab, and finally This is because it cools and hardens. The high-strength hot-rolled steel sheet obtained by rolling such a slab has a high concentration of Mn in the thickness direction center portion, so the hardness in the thickness direction center portion is higher than the hardness of the other portions. It has become.

  When cutting or punching is performed on such a high-strength steel plate, the end surface becomes an end surface composed of a shear surface and a fracture surface, but the machining tends to be particularly large near the boundary between the shear surface and the fracture surface. As a result, cracks and secondary shearing may occur after punching in the center of the plate thickness where the center segregation of Mn exists and the strength is different. Such cracks and secondary shears cause cracks during hole expansion and stretch flange forming, and therefore the characteristics are remarkably deteriorated. In particular, when the clearance at the time of punching is changed, the deterioration becomes remarkable.

  In consideration of actual members, high-strength hot-rolled steel sheets are often used as punched or as-cut, and therefore, it is required to have good characteristics such as elongation and hole expansion as they are cut. Here, the clearance at the time of punching or cutting varies at each position of the member, or the clearance may change with time due to wear of the punch or shear, so the actual member is not necessarily in the condition that the hole expandability is excellent. It is not always processed. As a result, even a steel plate excellent in hole expansibility may break.

As described above, in improving the characteristics of an actual member, it is required to stably exhibit excellent characteristics under a wide range of molding conditions, in addition to improving characteristics under ideal conditions.
CAMP-ISIJ Vol. 13 (2000), p399 CAMP-ISIJ vol. 13 (2000), p411 CAMP-ISIJ vol. 13 (2000), p391 JP-A-53-22812 Japanese Patent Laid-Open No. 1-2230715 JP 2003-321733 A JP 2004-256906 A Japanese Patent Application Laid-Open No. 11-296991 JP-A-63-293121

As described above, in order to improve stretch flangeability, in addition to improving material properties such as ductility and hole expandability, it is essential to improve the properties of the cut end face.
The present invention has been made in consideration of the improvement of material properties such as ductility and hole expansibility, as well as suppression of end face damage after cutting. An object of the present invention is to provide a high-strength steel sheet having a hole expandability similar to that of a structure and at the same time causing extremely little damage to the end face after cutting, and a method for producing the same.

The gist of the present invention aimed at solving the above problems is as follows.
(1) By mass%, C: 0.05% to 0.20%, Si: 0.3 to 2.00%, Mn: 1.3 to 2.6%, P: 0.001 to 0.03 %, S: 0.0001 to 0.01%, Al: less than 0.10%, N: 0.0005 to 0.0100%, O: 0.0005 to 0.007%, the balance being iron and It is a steel composed of inevitable impurities, the steel sheet structure is mainly composed of ferrite and bainite, the Mn segregation degree in the thickness direction (= center Mn peak concentration / average Mn concentration) is 1.20 or less, and the maximum tensile strength Is a high-strength steel sheet with little deterioration in characteristics after cutting, characterized by being 540 MPa or more.
(2) The high-strength steel sheet with less deterioration in characteristics after cutting according to (1), further comprising B: 0.0001 to less than 0.01% by mass.
(3) Furthermore, Cr: 0.01-1.0%, Ni: 0.01-1.0%, Cu: 0.01-1.0%, Mo: 0.01-1. The high-strength steel sheet with little deterioration in properties after cutting according to (1) or (2), characterized by containing one or more of 0%.
(4) Further, any one of (1) to (3), characterized by containing, in mass%, 0.001 to 0.14% of one or more of Nb, Ti, and V in total. A high-strength steel sheet with little property deterioration after cutting as described in the item.
(5) Further, any one of (1) to (4), characterized by containing 0.0001 to 0.5% in total of one or more of Ca, Ce, Mg, and REM in mass% 2. A high-strength steel sheet with little deterioration in properties after cutting according to item 1.
(6) A high-strength steel sheet with little deterioration in characteristics after cutting, characterized by having zinc-based plating on the surface of the high-strength steel sheet according to any one of (1) to (5).
(7) A cast slab characterized by performing casting while applying reduction in the plate thickness direction when casting the cast slab having the chemical component according to any one of (1) to (5).

(8) The cast slab described in (7) is directly or once cooled and then heated to 1050 ° C. or higher and subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature is reduced to a reduction rate of 85% or more. After performing hot rolling to 820 ° C to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is performed at 620 to 720 ° C. Completed, wound up in a temperature range of 400-630 ° C., pickled, cold rolled at a rolling reduction of 40-70%, and annealed at a maximum heating temperature of 760-870 ° C. when passing through a continuous annealing line Thereafter, the steel sheet is cooled between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more, and is maintained at a temperature range of 450 ° C. to 300 ° C. for 30 seconds or more. Manufacturing method.
(9) The cast slab described in (7) is directly or once cooled, then heated to 1050 ° C. or higher, subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature at a reduction rate of 85% or more. After performing hot rolling to 820 ° C to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is performed at 620 to 720 ° C. Completed, wound up in a temperature range of 400-630 ° C., pickled, cold-rolled with a rolling reduction of 40-70%, and passed through a continuous hot dip galvanizing line at a maximum heating temperature of 760-870 ° C. After annealing, after cooling between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more (zinc plating bath temperature−40) ° C. to (zinc plating bath temperature + 50) ° C., before being immersed in the zinc plating bath, Or either after immersion or both The method for producing a high-strength hot-dip galvanized steel sheet with little deterioration in properties after cutting, characterized by holding in a temperature range of (galvanizing bath temperature + 50) ° C. to 300 ° C. for 30 seconds or more.
(10) The casting slab described in (7) is directly or once cooled, then heated to 110001050 ° C. or higher, subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature at a reduction rate of 85% or more. After performing hot rolling to 820 ° C to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is performed at 620 to 720 ° C. Completed, wound up in a temperature range of 400-630 ° C., pickled, cold-rolled with a rolling reduction of 40-70%, and passed through a continuous hot dip galvanizing line at a maximum heating temperature of 760-870 ° C. After annealing, cooling between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more to (zinc plating bath temperature−40) ° C. to (zinc plating bath temperature + 50) ° C., and then, if necessary, 460 to 540 Alloyed at a temperature of ℃, Cutting characterized by holding for at least 30 seconds in a temperature range of (zinc plating bath temperature +50) ° C. to 300 ° C. either before or after immersion in the lead plating bath or after alloying treatment. A method for producing a high-strength galvannealed steel sheet with less characteristic deterioration later.
(11) The method for producing a high-strength steel sheet with less characteristic deterioration after cutting according to (8), wherein a high-strength steel sheet is produced by the method of (8) and then zinc-based electroplating is performed.

  According to the present invention, by controlling the steel plate components, casting conditions, and rolling conditions, it is possible to stably obtain a high-strength steel plate having a maximum tensile strength of 540 MPa or more and less characteristic deterioration after cutting.

As a result of diligent investigations, the present inventors have found that the cause of deterioration in hole expandability and stretch flangeability due to clearance change is due to changes in end face damage during cutting or punching due to center segregation of Mn. Found that there is. Specifically, when the clearance is increased, the stress in the thickness direction acting during punching or cutting is increased. As a result, it has been found that cracks may occur in the central portion (segregation portion) of the plate thickness, and the subsequent workability is deteriorated.
The present invention is described in detail below.

The high-strength steel sheet with less characteristic deterioration after cutting according to the present invention does not indicate a difference in characteristics between a test piece prepared by mechanical grinding and an as-processed test piece such as cutting or punching, but has a large clearance. Even if it changes, it means a steel sheet that can ensure good stretch flangeability.
Further, the machining is not limited to shear cutting, but refers to a processing method in which there is an end face as it is cut (punched) including punching using a punch, and then the machining is performed.

First, the segregation degree of Mn will be described.
As a result of diligent investigations by the inventors, the steel sheet structure is made of a ferrite and a hard structure, and the cracking of the end face is reduced by reducing the segregation degree of manganese while reducing the hardness difference between the hard structure and the ferrite. It has been found that by suppressing it, it is possible to reduce the characteristic deterioration after cutting.

  When the cross section of the high-strength steel plate is subjected to CMA analysis or EPMA analysis, a Mn segregation zone in which Mn is segregated at a concentration higher than the average concentration of Mn is observed at the center in the thickness direction. A quantification of the degree of segregation of Mn is the degree of Mn segregation in the plate thickness direction, which can be measured using CMA or EPMA. The Mn segregation degree in the plate thickness direction is defined by a value obtained by dividing the central Mn peak concentration by the average Mn concentration. In the high-strength steel plate according to the present invention, it is preferably 1.2 or less, more preferably 1.15. Or less, more preferably 1.0 or less. Further, it is desirable that the width of the Mn segregation band in the plate thickness direction is also narrow. The width of the Mn segregation zone is preferably 20 μm or less, more preferably 15 μm or less, and still more preferably 10 μm or less. In some cases, there is a case where there is non-segregation in which the Mn concentration is lower than the average Mn concentration at the center of the plate thickness. Such a steel plate is also included in the high-strength steel plate according to the present invention.

  In order to reduce the degree of segregation of Mn, it is necessary to cast the slab while reducing it. For example, when casting a slab having a thickness of 240 mm, a reduction of 5 mm or more is applied. That is, if the entry side thickness is 245 mm, the exit side thickness is 240 mm. However, the degree of segregation of Mn occurs when Mn is discharged into the molten steel in the process of solidifying the molten steel. Therefore, even if the solidified slab is reduced, Mn segregation does not improve. For this reason, it is necessary to perform the reduction before the molten steel is completely solidified.

  Further, in a normal high-strength steel sheet, S is an element that is easily segregated in the central portion, like Mn. Segregation of Mn and S is undesirable because it promotes the formation of MnS in the central segregation part and degrades the hole expandability. In addition, since MnS present in these central segregation portions extends long along the plate thickness direction, it causes crack formation along the plate thickness direction during punching or cutting, and subsequent elongation, hole expansion, This is undesirable because it causes deterioration of mechanical properties such as stretch flangeability.

  In addition, Mn segregation greatly varies the structure and properties in the plate thickness direction in steel that utilizes structure strengthening. When controlling the structure of a steel sheet in continuous annealing equipment or alloyed hot dip galvanizing equipment, after annealing into a two-phase region consisting of ferrite and austenite or an austenite single-phase region, ferrite is formed during the cooling process. . Or, in many cases, the structure control is performed at the time of subsequent residence in a specific temperature range or alloying treatment. Since Mn is known to delay the ferrite transformation and bainite transformation, the steel sheet structure changes with a partial change in Mn concentration. Specifically, since the ferrite transformation and bainite transformation are delayed at the center segregation part, the thickness center tends to be a structure containing a lot of martensite. As a result, the thickness center becomes extremely hard. Variations in steel plate strength cause variations in characteristics such as elongation within the plate thickness. In cutting and punching with large deformation, the concentration of deformation is caused at the interface between the hard Mn segregation zone and the normal part, resulting in large stress concentration and, in some cases, crack formation (peeling) at the interface. . As a result, it is not desirable because it causes deterioration of mechanical properties such as elongation, hole expansion and stretch flangeability of the steel sheet after cutting and punching.

In particular, the cutting or punching process is particularly likely to cause crack formation because stress tends to occur in the direction of peeling the hard tissue layer due to the Mn segregation zone existing in layers in the plate thickness direction.
Such a problem is likely to be a problem in a high-strength steel sheet for strengthening the structure. That is, high-strength steel sheets utilizing structural strengthening such as DP steel and TRIP steel tend to add a large amount of Mn in order to control the structure in hot rolling, continuous annealing equipment, and hot dip galvanizing equipment. This tendency becomes more prominent as the strength of the steel plate increases. Therefore, it tends to be a problem with a high strength steel plate of 540 MPa or more. In addition, a high-strength steel sheet of less than 540 MPa often uses a technique other than structure strengthening (for example, solid solution strengthening) as a technique for increasing strength, and often has a single-phase structure. As a result, even when a large amount of Mn is contained, the strength fluctuation due to the structure change hardly occurs, and the characteristic deterioration due to cutting or punching hardly occurs.

  The high-strength steel sheet according to the present invention has a Mn segregation degree in the thickness direction of 1.2 or less, thereby preventing deterioration of mechanical properties such as hole expansion and stretch flangeability, and characteristic deterioration accompanying cutting and punching. Can be prevented.

  In addition, when the Mn segregation zone exists, the ferrite transformation and the bainite transformation are delayed, so that a structure containing a large amount of martensite is formed and is easily hardened. As a result, in a steel plate with significant Mn segregation, a difference in strength occurs between the central segregation portion and the other portions. The hardness difference at this time can be measured with a Vickers tester. Therefore, in the high-strength steel plate according to the present invention, it is preferable that the hardness difference between the central segregation portion and the other portion by the Vickers tester at a load of 50 gf is Hv 70 or less. When the hardness difference exceeds Hv70, the characteristic deterioration due to punching or cutting becomes remarkable, which is not preferable.

  The reason why the load during the Vickers test was set to 50 gf is to measure the hardness of the segregation zone. Even if the segregation zone at the center formed at the time of casting is coarse in the slab, the thickness is reduced and narrowed by hot rolling or cold rolling. That is, when a test is performed with a large load of 50 kgf, the indentation size increases, and thus a hardness variation in a narrow region such as center segregation cannot be detected. On the other hand, when the load is extremely low, the measurement result differs depending on whether the indentation is applied to the hard tissue or the soft tissue, and therefore, the hardness variation due to the difference in the structure distribution such as the central segregation zone cannot be detected. Therefore, as a preliminary experiment, we measured the center segregation band width by CMA analysis or EPMA analysis, or Vickers hardness measurement at various loads, and compared the characteristic change due to hardness fluctuation and mechanical cutting. In the present invention, the load during the Vickers test is set to 50 gf.

Regarding the measurement location, in order to investigate the strength fluctuation in the thickness direction, it is preferable to measure the hardness at the thickness 1/4 t position and the center segregation position. And each 10 points | pieces are measured along a plate | board thickness direction, and what is necessary is just to let the average value be each hardness. In the thin steel plate, the center segregation zone is subjected to hot rolling and cold rolling, and thus is stretched and has a small thickness, which may be difficult to distinguish. In the present invention, it is desirable to specify the exact position of the center segregation by polishing the steel plate and then etching with a Nital reagent to reveal the structure. That is, the central segregation zone containing a large amount of Mn is distinguishable because it becomes a structure containing a lot of martensite in the structure strengthened steel. A hardness difference (ΔHv _{50 g} ) as measured by the Vickers test is 70 or less within the preferable range of the present invention.

Next, the reason for limiting the structure of the steel sheet will be described.
The reason why the steel sheet structure is a multiphase structure of ferrite and hard structure is to obtain excellent ductility. Since soft ferrite is rich in ductility, it is essential to obtain excellent ductility. In addition, it is possible to increase the strength while ensuring excellent ductility by dispersing an appropriate amount of hard structure. In order to ensure excellent ductility, it is necessary to use a ferrite main phase. Further, residual austenite may be included. Residual austenite is transformed into martensite at the time of deformation, thereby hardening the processed portion and hindering concentration of deformation. As a result, particularly excellent ductility can be obtained.

  The hard structure is desirably 50% or more of the bainite structure. The bainite structure is known to be softer than martensite. Therefore, by forming the hard structure into a soft bainite structure, it is possible to suppress the formation of microvoids at the interface between the ferrite and the hard structure during the hole expanding process. The reason why the bainite structure is 50% or more when the hard structure is less than 50% of the volume ratio of the hard structure is that martensite and retained austenite are sufficiently separated and dispersed at the site of crack propagation during hole expansion processing. It is because it is thought that it must not be.

  The volume ratio of the hard tissue is desirably 5% or more. This is because it is difficult to ensure the strength of 540 MPa or more when the volume fraction of the hard tissue is less than 5%. The upper limit is not particularly defined, and excellent ductility and hole expandability, which are the effects of the present invention, are provided, but in the TS range of 590 to 1080 MPa, in order to achieve both ductility and hole expandability or stretch flangeability. It is desirable to include ferrite with a volume ratio exceeding 50%.

  The steel sheet structure is basically a composite structure of ferrite and bainite, but may include residual austenite, martensite, cementite, pearlite, and the like as other hard structures.

  Identification of each phase of the above microstructure, ferrite, pearlite, cementite, martensite, bainite, austenite and the remaining structure, observation of the existing position and measurement of the area ratio are disclosed in Nital reagent and JP-A-59-219473. It is possible to corrode the steel plate rolling direction cross section or the rolling direction perpendicular direction cross section with the above-mentioned reagent, and to quantify by 1000 times optical microscope observation and 1000 to 100000 times scanning and transmission electron microscopes. The structure can also be discriminated from crystal orientation analysis using the FESEM-EBSP method and micro region hardness measurement such as micro Vickers hardness measurement.

If TS is 540 MPa or more, if it is less than this strength, it is possible to achieve both high ductility and hole expandability with TS of less than 540 MPa by increasing the strength of the ferrite single-phase steel using solid solution strengthening. It is because it can plan. In particular, when considering securing TS of 540 MPa or more, in order to ensure excellent ductility, it is necessary to perform reinforcement using martensite or retained austenite, and deterioration of hole expansibility becomes remarkable.
The crystal grain size of ferrite is not particularly limited, but it is preferably 7 μm or less in terms of nominal grain size from the viewpoint of balance of strength elongation.

Next, the reasons for limiting the components of the present invention will be described. All units are mass%.
(C: 0.05% to 0.20%)
C is an essential element when strengthening the structure using bainite or martensite. If C is less than 0.05%, it is difficult to ensure a strength of 540 MPa or more, so the lower limit was set to 0.05%. On the other hand, the reason why the C content is 0.20% or less is that when C exceeds 0.20%, it becomes difficult to ensure spot weldability. For this reason, the range of C is set to 0.05% to 0.20%.

(Si: 0.3-2.00%)
In addition to being a strengthening element, Si does not dissolve in cementite, so it suppresses the formation of coarse cementite at grain boundaries. If less than 0.3% is added, strengthening due to solid solution strengthening cannot be expected, or formation of coarse cementite at grain boundaries cannot be suppressed, so 0.3% or more needs to be added. On the other hand, addition exceeding 2.00% excessively increases the retained austenite, and deteriorates the hole expandability and stretch flangeability after punching or cutting. Therefore, the upper limit needs to be 2.00%. In addition, since the oxide of Si has poor wettability with hot dip galvanizing, it causes non-plating. Therefore, in manufacturing a hot-dip galvanized steel sheet, it is necessary to control the oxygen potential in the furnace and suppress the formation of Si oxide on the steel sheet surface.

(Mn: 1.3-2.6%)
Since Mn is an austenite stabilizing element at the same time as a solid solution strengthening element, it suppresses the transformation of austenite to pearlite. If it is less than 1.3%, the rate of pearlite transformation is too high, and the steel sheet structure cannot be made a composite structure of ferrite and bainite, and a TS of 540 MPa or more cannot be secured. Moreover, the hole expansibility is also inferior. For this reason, the lower limit is set to 1.3% or more. On the other hand, when Mn is added in a large amount, co-segregation with P and S is promoted and workability is significantly deteriorated. Therefore, the upper limit is set to 2.6%.

(P: 0.001 to 0.03%)
P tends to segregate in the central part of the plate thickness of the steel sheet, causing the weld to become brittle. If it exceeds 0.03%, embrittlement of the weld becomes significant, so the appropriate range is limited to 0.03% or less. Although the lower limit value of P is not particularly defined, it is preferable to set this value as the lower limit value because it is economically disadvantageous to set it to less than 0.001%.

(S: 0.0001 to 0.01%)
S adversely affects weldability and manufacturability during casting and hot rolling. Therefore, the upper limit is set to 0.01% or less. Since it is economically disadvantageous to set the lower limit value of S to less than 0.0001%, this value is preferably set as the lower limit value. In addition, since S is combined with Mn to form coarse MnS, the hole expandability is lowered. For this reason, it is necessary to reduce as much as possible in order to improve hole expansibility.

(Al: less than 0.10%)
Al may be added because it promotes ferrite formation and improves ductility. It can also be used as a deoxidizer. However, excessive addition increases the number of Al-based coarse inclusions, causing deterioration of hole expansibility and surface scratches. From this, the upper limit of Al addition was made less than 0.1%. The lower limit is not particularly limited, but it is difficult to set the lower limit to 0.0005% or less, which is a practical lower limit.

(N: 0.0005 to 0.0100%)
N (nitrogen) forms coarse nitrides and degrades bendability and hole expansibility, so it is necessary to suppress the addition amount. This is because when N exceeds 0.01%, this tendency becomes remarkable. Therefore, the range of N content is set to 0.01% or less. In addition, it is better to use less because it causes blowholes during welding. Although the lower limit is not particularly defined, the effect of the present invention is exhibited. However, if the N content is less than 0.0005%, the manufacturing cost is significantly increased, and this is a substantial lower limit.

(O: 0.0005 to 0.007%)
O (oxygen) forms an oxide and degrades bendability and hole expansibility, so the addition amount must be suppressed. In particular, oxides often exist as inclusions, and when they are present on the punched end surface or cut surface, they form notched scratches and coarse dimples on the end surface, so when expanding holes or during strong processing, It causes stress concentration and becomes the starting point of crack formation, resulting in a significant deterioration of hole expansibility or bendability. This is because this tendency becomes significant when O exceeds 0.007%, so the upper limit of the O content is set to 0.007% or less. Since setting it as less than 0.0005% invites excessive cost and is not economically preferable, this was made into the minimum. However, even if O is less than 0.0005%, TS of 540 MPa or more, which is the effect of the present invention, and excellent ductility can be secured.

(B: 0.0001% or more and less than 0.01%)
B is effective for strengthening grain boundaries and strengthening steel by addition of 0.0001% by mass or more. However, when the addition amount exceeds 0.010% by mass, the effect is not only saturated but also heat The upper limit is set to 0.010% because the manufactured product at the time of extension is lowered.

(Cr: 0.01-1.0%)
Cr is a strengthening element and is important for improving hardenability. However, if it is less than 0.01%, these effects cannot be obtained, so the lower limit was set to 0.01%. If the content exceeds 1%, a significant increase in cost is caused, so the upper limit was made 1%.

(Ni: 0.01-1.0%)
Ni is a strengthening element and is important for improving hardenability. However, if it is less than 0.01%, these effects cannot be obtained, so the lower limit was set to 0.01%. If the content exceeds 1%, a significant increase in cost is caused, so the upper limit was made 1%.

(Cu: 0.01 to 1.0%)
Cu is a strengthening element and is important for improving hardenability. However, if it is less than 0.01%, these effects cannot be obtained, so the lower limit was set to 0.01%. On the other hand, if the content exceeds 1%, the manufacturability at the time of production and hot rolling is adversely affected, so the upper limit was made 1%.

(Mo: 0.01 to 1.0%)
Mo is a strengthening element and is important for improving hardenability. However, if it is less than 0.01%, these effects cannot be obtained, so the lower limit was set to 0.01%. If the content exceeds 1%, the cost is significantly increased, so the upper limit is 1%, but 0.3% or less is more preferable.

(Nb: 0.001 to 0.14%)
Nb is a strengthening element. It contributes to increasing the strength of steel sheets by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. If the addition amount is less than 0.001%, these effects cannot be obtained, so the lower limit was set to 0.001%. If the content exceeds 0.14%, the precipitation of carbonitride increases and the formability deteriorates, so the upper limit was made 0.14%.

(Ti: 0.001 to 0.14%)
Ti is a strengthening element. It contributes to increasing the strength of steel sheets by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. If the addition amount is less than 0.001%, these effects cannot be obtained, so the lower limit was set to 0.001%. If the content exceeds 0.14%, the precipitation of carbonitride increases and the formability deteriorates, so the upper limit was made 0.14%.

(V: 0.001 to 0.14%)
V is a strengthening element. It contributes to increasing the strength of steel sheets by strengthening precipitates, strengthening fine grains by suppressing the growth of ferrite crystal grains, and strengthening dislocations by suppressing recrystallization. If the addition amount is less than 0.001%, these effects cannot be obtained, so the lower limit was set to 0.001%. If the content exceeds 0.14%, the precipitation of carbonitride increases and the formability deteriorates, so the upper limit was made 0.14%.

(One or more of Ca, Ce, Mg, and REM in total 0.0001 to 0.5%)
One or two or more selected from Ca, Ce, Mg, and REM can be added in a total amount of 0.0001 to 0.5%. Ca, Ce, Mg, and REM are elements used for deoxidation. By containing one or more kinds in total of 0.0001% or more, the oxide size after deoxidation can be reduced, and the hole is expanded. Contributes to improved performance.

  However, when the content exceeds 0.5% in total, it causes deterioration of molding processability. Therefore, the content is made 0.0001 to 0.5% in total. Note that REM is an abbreviation for Rare Earth Metal and refers to an element belonging to the lanthanoid series. In the present invention, REM and Ce are often added by misch metal and may contain a lanthanoid series element in combination with La and Ce. Even if these lanthanoid series elements other than La and Ce are included as inevitable impurities, the effect of the present invention is exhibited. However, the effects of the present invention are exhibited even when metal La or Ce is added.

Next, the reasons for limiting the production conditions of the steel sheet of the present invention will be described.
In order to reduce the degree of segregation of Mn, it is necessary to cast while reducing the slab during casting. Since the Mn segregation degree occurs when Mn is discharged into the molten steel in the process of solidification of the molten steel, even if the solidified slab is reduced, the Mn segregation does not improve. Therefore, the slab must be reduced before the molten steel is completely solidified. By performing the reduction before complete solidification, Mn discharged to the center can be diffused in the entire plate thickness direction of the slab, and the Mn segregation degree can be reduced. Specifically, for example, when casting a slab having a thickness of 240 mm, it is preferable to apply a reduction of 5 mm or more. That is, if the entry side thickness is 245 mm, the exit side thickness may be 240 mm.

In order to reduce the width of the central segregation zone in the thickness direction, it is desirable to increase the thickness reduction ratio. The thickness reduction ratio refers to the product plate thickness relative to the slab thickness. If a cold-rolled steel sheet having a product plate thickness of 3.0 mm is manufactured, the slab thickness is desirably 200 mm or more. Even if the central segregation of Mn or the like is reduced by applying a reduction during casting, it is difficult to completely remove it. From this, by increasing the thickness of the slab, the ratio of the center segregation zone formed in the slab is reduced, and the thickness of the center segregation zone is reduced by subsequent hot rolling or cold rolling, thereby adversely affecting the thickness. To remove it.
For the same reason, it is desirable to increase the rolling reduction in hot rolling and cold rolling.

  In addition, in order to reduce the influence of the structure fluctuation | variation by Mn segregation, the steel-plate structure control at the time of hot rolling works advantageously. That is, by setting the coiling temperature of the hot-rolled steel sheet to 630 ° C. or less, it is possible to make the structure of the hot-rolled sheet a uniform structure mainly composed of bainite structure, which is advantageous for improving the properties after cold rolling and annealing. work. Therefore, it is desirable that the winding temperature is 630 ° C. or lower.

  The hot-rolled slab heating temperature needs to be 1050 ° C. or higher. When the slab heating temperature is excessively low, the finish rolling temperature is lower than 820 ° C., resulting in a two-phase rolling of ferrite and austenite, and the hot rolled sheet structure becomes a heterogeneous mixed grain structure, which has undergone cold rolling and annealing processes. However, the non-uniform structure is not eliminated and the ductility and hole expansibility are poor. Moreover, since this steel plate ensures a tensile maximum strength of 540 MPa or more after annealing, a relatively large amount of alloy elements is added, and thus the strength at the time of finish rolling tends to be high. The decrease in the slab heating temperature causes a decrease in the finish rolling temperature, further increases the rolling load, and there is a concern that rolling may become difficult or the shape of the steel sheet after rolling may be poor. It is necessary to set it to 1050 ° C. or higher. The upper limit of the slab heating temperature is not particularly defined, and the effect of the present invention is exhibited. However, since it is economically undesirable to make the heating temperature too high, the upper limit of the heating temperature should be less than 1300 ° C. Is desirable.

The slab cast as described above is hot-rolled. In the present invention, it is desirable to hot-roll the slab heated to 1050 ° C. at a reduction rate of 70% or more before finish rolling. The reduction ratio is set to 70% or more in order to refine the MnS size formed in the Mn segregation part and improve the property deterioration after punching with the product plate or after cutting. That is, it is known that S is combined with Mn to form MnS at the hot rolling stage. When cold rolling, MnS formed at the time of hot rolling becomes a form expanded in the rolling direction, so that stress concentration at the tip and interface becomes remarkable, resulting in deterioration of characteristics during processing. Or when MnS exists in a cutting part or a punching part, a wrinkle may be formed in the end surface after a process, and the further characteristic deterioration is caused. This tendency is known to become more prominent as the volume ratio of MnS increases, and reduction or refinement of the MnS volume ratio is required. Moreover, since MnS is easily formed when the Mn concentration is high, it is particularly easy to form MnS at the Mn segregation part. Furthermore, if the width of Mn segregation is large, the MnS formed is also large, and the characteristic deterioration becomes remarkable. Therefore, in order to improve the characteristics after cutting including hole expandability and stretch flangeability, it is necessary to refine MnS formed in the hot-rolled sheet. Therefore, it is desirable to reduce the width of Mn segregation and to refine MnS by increasing the rolling reduction in the hot rolling stage. Specifically, it is known that MnS is dissolved in austenite at a high temperature before slab heating or finish rolling, and precipitates at a lower temperature side than this. Therefore, it is important to reduce the width of Mn segregation and reduce MnS before MnS is precipitated.
Since these effects become significant when the rolling reduction is 70% or more, it is desirable that the rolling reduction is as high as possible. Although the upper limit is not particularly defined, it is difficult to make it more than 90% from the viewpoint of productivity and equipment restrictions, so 90% is a practical upper limit.

The finish rolling temperature needs to be in the range of 820 ° C. or higher and 930 ° C. or lower. When the finish rolling temperature is in the two-phase region of austenite + ferrite, the structure non-uniformity in the steel sheet increases and the formability after annealing deteriorates, so 820 ° C. or higher is desirable.
On the other hand, the effect of the present invention is exhibited even if the upper limit of the finishing temperature is not particularly defined, but when the finishing rolling temperature is excessively high and used, in order to secure the temperature, the slab heating temperature must be excessively high. Don't be. For this reason, the upper limit temperature of the finish rolling temperature is desirably 930 ° C. or lower.

The rolling reduction ratio of the subsequent rolling is 85% or more in total. The rolling reduction may be calculated by dividing the sheet thickness before rolling by the sheet thickness before rolling and multiplying by 100. The rolling reduction in finish rolling is determined for the same reason. That is, it is difficult to sufficiently reduce the thickness of the Mn segregation zone by rolling with a rolling reduction of less than 85%. Further, rolling exceeding 98% is an excessive addition to the equipment, so this is the upper limit. 90 to 94% is a more preferable rolling reduction.
However, since MnS can be precipitated even in austenite at low temperatures, it is desirable to reduce as much as possible before finish rolling for the purpose of refining MnS.

Moreover, it is desirable that the rolled steel sheet is water-cooled after completion of hot rolling and before winding. Water cooling starts water cooling within 7 seconds after the final rolling, water cooling is performed at an average cooling rate of 720 to 800 ° C. at a cooling rate of 25 ° C./second or more, and water cooling is completed at 620 to 720 ° C. desirable.
The effect of the present invention is exhibited without particular limitation on the cooling conditions after finish rolling. However, from the viewpoint of improving the product sheet characteristics by homogenizing the hot-rolled sheet structure, the average cooling rate from finish rolling to winding is desirably 10 ° C./second or more, more preferably 25 ° C./second or more. It is. On the other hand, excessively increasing the cooling rate not only saturates the effect but also causes a significant increase in capital investment. Therefore, the cooling rate is desirably 900 ° C./second or less. The cooling method is not particularly limited, and air cooling, gas cooling, water cooling, mist, or a method using any one of them may be used.

  The coiling temperature after water cooling needs to be 630 ° C. or lower. When the temperature exceeds 630 ° C., coarse ferrite and pearlite structures exist in the hot-rolled structure, so that the structure non-uniformity after annealing increases and the ductility of the final product deteriorates. It is more preferable to wind up at 600 ° C. or less from the viewpoint of improving the strength ductility balance by making the microstructure after annealing fine, and further improving the hole expandability by uniformly dispersing the second phase. In addition, winding at a temperature exceeding 630 ° C. is not preferable because the thickness of the oxide formed on the steel sheet surface is excessively increased, and the pickling property is poor. Although the lower limit is not particularly defined, the effect of the present invention is exhibited. However, when winding at a temperature of 400 ° C. or lower, the strength of the hot rolled sheet increases excessively, and cold rolling becomes difficult. Is preferably 400 ° C. or higher. Note that rough rolling sheets may be joined to each other during hot rolling to continuously perform finish rolling. Moreover, you may wind up a rough rolling board once.

  The hot-rolled steel sheet thus manufactured is pickled. Since pickling can remove oxides on the surface of steel sheets, it can improve the chemical conversion properties of cold-rolled high-strength steel sheets as final products, and improve the hot-plating properties of cold-rolled steel sheets for hot-dip galvanized or galvannealed steel sheets It is important for that. Moreover, pickling may be performed once, or pickling may be performed in a plurality of times.

  The pickled hot-rolled steel sheet is cold-rolled at a rolling reduction of 40 to 70% and passed through a continuous annealing line or a continuous hot-dip galvanizing line. If the rolling reduction is less than 40%, it is difficult to keep the shape flat. Moreover, since the ductility of the final product becomes poor, this is the lower limit. On the other hand, cold rolling exceeding 70% makes the cold rolling difficult because the cold rolling load becomes too large. 45 to 65% is a more preferable range. The effect of the present invention is exhibited without particularly specifying the number of rolling passes and the rolling reduction for each pass.

  After cold rolling, annealing is performed at a maximum heating temperature of 760 to 870 ° C. This is because if the temperature is lower than 760 ° C., an excessive time is required for reverse transformation from cementite or pearlite to austenite. In addition, when the maximum attainable temperature is less than 760 ° C., part of cementite and pearlite cannot be transformed into austenite and remain in the steel sheet structure even after annealing. Since this cementite and pearlite are coarse, it is not preferable because it causes deterioration of hole expansibility. Alternatively, bainite and martensite formed by transformation of austenite, or austenite itself can be transformed into martensite during processing, so that a strength of 540 MPa or more can be achieved, so that part of cementite and pearlite is austenite. If it does not transform into, the hard structure becomes too small and it is not possible to secure a strength of 540 MPa or more. Therefore, the lower limit of the maximum heating temperature needs to be 760 ° C. On the other hand, it is economically undesirable to raise the heating temperature excessively. Therefore, it is desirable that the upper limit of the heating temperature is 870 ° C.

  Thereafter, it is necessary to cool between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more. If the cooling rate is too low, austenite transforms into a pearlite structure during the cooling process, and thus a hard structure of an amount necessary for a strength of 540 MPa or more cannot be secured. Even if the cooling rate is increased, there is no problem in terms of the material. However, excessively increasing the cooling rate leads to an increase in manufacturing cost, so the upper limit is preferably set to 200 ° C./second. The cooling method may be roll cooling, air cooling, water cooling, or any combination of these methods.

  Subsequently, it is necessary to hold at a temperature range of 450 ° C. to 300 ° C. for 30 seconds or more. This is because austenite is transformed into bainite. When holding in a temperature range higher than 450 ° C., coarse cementite precipitates at the grain boundaries, so that the hole expandability is greatly deteriorated. Therefore, the upper limit temperature is set to 450 ° C. On the other hand, when the holding temperature is less than 300 ° C., bainite transformation hardly occurs, and austenite is transformed into martensite in the subsequent cooling process. As a result, the structure cannot be made of ferrite and bainite, and the hole expandability is greatly deteriorated. Therefore, 300 ° C. is the lower limit temperature.

  In a temperature range of 450 ° C. to 300 ° C. for less than 30 seconds, even if a bainite structure is formed, the volume ratio is not sufficient, and the remaining austenite is transformed into martensite in the subsequent cooling process, Poor hole expandability. For this reason, the lower limit of the residence time is 30 seconds or more. The upper limit of the residence time is not particularly defined, and the effect of the present invention can be obtained. However, the increase in residence time can be achieved by lowering the plate passing speed when considering heat treatment in a facility having a finite length. This means that it is not preferable because it is not economical.

  In addition, holding does not only mean isothermal holding, but means retaining in a temperature range of 450 to 300 ° C. That is, after cooling to 300 ° C., it may be heated to 450 ° C., or may be cooled to 450 ° C. and then cooled to 300 ° C.

  After the heat treatment, it is desirable to perform skin pass rolling in order to control surface roughness, plate shape control, or suppression of yield point elongation. The reduction ratio of the skin pass rolling at that time is preferably in the range of 0.1 to 1.5%. If the skin pass rolling rate is less than 0.1%, the effect is small and control is difficult, so this is the lower limit. If it exceeds 1.5%, the productivity is remarkably lowered, so this is the upper limit. The skin pass may be performed inline or offline. Further, a skin pass having a desired reduction rate may be performed at once, or may be performed in several steps.

  The maximum heating temperature when passing through the hot dip galvanizing line after cold rolling is set to 760 to 870 ° C. for the same reason as when passing through the continuous annealing line. Regarding the cooling after annealing, it is necessary to cool between 630 ° C. and 570 ° C. at 3 ° C./second or more for the same reason as when the continuous annealing line is passed through.

  The plating bath immersion plate temperature is preferably in a temperature range from a temperature 40 ° C. lower than the hot dip galvanizing bath temperature to a temperature 50 ° C. higher than the hot dip galvanizing bath temperature. If the bath immersion plate temperature is lower than the hot dip galvanizing bath temperature −40) ° C., the heat removal at the time of immersion in the plating bath is large, and part of the molten zinc may solidify and deteriorate the plating appearance. The lower limit is (hot dip galvanizing bath temperature −40) ° C. However, even if the plate temperature before immersion is lower than (hot dip galvanizing bath temperature −40) ° C., reheating is performed before immersion in the plating bath, and the plate temperature is set to (hot dip galvanizing bath temperature −40) ° C. or higher. It may be immersed in. On the other hand, if the plating bath immersion temperature exceeds (hot dip galvanizing bath temperature + 50) ° C., operational problems accompanying the rise of the plating bath temperature are induced. Further, the plating bath may contain Fe, Al, Mg, Mn, Si, Cr, etc. in addition to pure zinc.

  Moreover, when alloying a plating layer, it carries out at 460 degreeC or more. When the alloying treatment temperature is less than 460 ° C., the progress of alloying is slow and the productivity is poor. If the upper limit exceeds 540 ° C., carbides are formed to reduce the volume ratio of the hard structure (martensite, bainite, retained austenite), and it becomes difficult to ensure the strength of 540 MPa or more, so this is the upper limit.

  It is necessary to perform an additional heat treatment for 30 seconds or more in the temperature range of (zinc plating bath temperature + 50) ° C. to 300 ° C. either before or after immersion in the plating bath, or after immersion. The upper limit of the heat treatment temperature is set to (zinc plating bath temperature +50) ° C. Above this temperature, the formation of cementite and pearlite becomes remarkable, and the volume fraction of the hard tissue is reduced, so it is difficult to ensure the strength of 540 MPa or more. It is to become. On the other hand, if it is less than 300 ° C., the progress of bainite transformation is too slow, and the structure cannot be made of ferrite and bainite. Therefore, the lower limit is set to 300 ° C. or higher.

  The holding time needs to be 30 seconds or more. When the holding time is less than 30 seconds, even if a bainite structure is formed, the volume ratio is not sufficient, and the remaining austenite is transformed into martensite in the subsequent cooling process, and therefore the hole expandability is poor. For this reason, the lower limit of the residence time is 30 seconds or more. The upper limit of the residence time is not particularly defined, and the effect of the present invention can be obtained. However, the increase in residence time can be achieved by lowering the plate passing speed when considering heat treatment in a facility having a finite length. This means that it is not preferable because it is not economical. The holding time does not simply mean isothermal holding, but means residence in this temperature range, and includes cooling and heating in this temperature range.

  Further, the additional heat treatment for 30 seconds or more in the temperature range of (zinc plating bath temperature + 50) ° C. to 300 ° C. may be performed either before or after immersion in the plating bath, or both. Absent. As long as a hard structure having a crystal orientation difference of less than 9 ° with respect to ferrite as the main phase can be secured, the strength of 540 MPa or more, which is the effect of the present invention, can be achieved regardless of the additional heat treatment. This is because excellent ductility and hole expansibility can be obtained.

  After the heat treatment, it is desirable to perform skin pass rolling in order to control surface roughness, plate shape control, or suppression of yield point elongation. The reduction ratio of the skin pass rolling at that time is preferably in the range of 0.1 to 1.5%. If the skin pass rolling rate is less than 0.1%, the effect is small and control is difficult, so this is the lower limit. If it exceeds 1.5%, the productivity is remarkably lowered, so this is the upper limit. The skin pass may be performed inline or offline. Further, a skin pass having a desired reduction rate may be performed at once, or may be performed in several steps.

  Further, in order to further improve the plating adhesion, the present invention does not depart from the present invention even if the steel plate is plated with Ni, Cu, Co, or Fe alone or before the annealing.

Further, regarding annealing before plating, “after degreasing pickling, heating in a non-oxidizing atmosphere, annealing in a reducing atmosphere containing H ₂ and N ₂ , cooling to near the plating bath temperature, and invading the plating bath. Zenjimer method called “Kizuke”, an all-reduction furnace method called “immersion in the plating bath after adjusting the atmosphere during annealing, first oxidizing the steel plate surface, and then reducing it before cleaning by plating” Alternatively, there is a flux method such as “after degreasing and pickling a steel plate, and then fluxing it with ammonium chloride and soaking it in a plating bath”, etc. The effect of can be demonstrated. Regardless of the annealing method prior to plating, the dew point during heating is set to −20 ° C. or higher, which advantageously works on the wettability of the plating and the alloying reaction at the time of plating alloying.

  In addition, even if this cold-rolled steel sheet is electroplated, the tensile strength, ductility, and hole expandability of the steel sheet are not impaired at all. That is, the steel sheet of the present invention is also suitable as a material for electroplating. Even if an organic film or upper layer plating is performed, the effect of the present invention can be obtained.

  In addition, the material of the high strength and high ductility hot dip galvanized steel sheet excellent in formability and hole expansibility of the present invention is manufactured through refining, steel making, casting, hot rolling, and cold rolling processes that are normal iron making processes. However, even if it is manufactured by omitting part or all of it, the effects of the present invention can be obtained as long as the conditions according to the present invention are satisfied.

A 240 mm thick slab containing the chemical components shown in Table 1 was cast. During casting of the slab, rolling was performed with the rolling amount shown in Table 2 before the molten steel was completely solidified.
Next, the obtained slab was heated to 1230 ° C., and then hot-rolled while being reduced at the reduction rate shown in Tables 2 and 3, and further subjected to finish hot rolling under the conditions shown in Tables 2 and 3. . Then, cold rolling, annealing, cooling and heat treatment were performed under the conditions shown in Tables 2 and 3. In this way, a steel plate was produced. The steel plates subjected to hot rolling and cold rolling are indicated as CR in Tables 2 and 3.

  Moreover, about some steel plates, it was set as the hot dip galvanized steel plate by letting a continuous hot dip galvanizing line pass before and after said heat processing. In the continuous hot dip galvanizing line, after annealing under the conditions shown in Table 2 or Table 3, cooling between 630 ° C. and 570 ° C. was carried out to the range of 420 ° C. to 500 ° C. above the average cooling rate shown in Table 2 or Table 3. Thereafter, hot dip galvanization was performed, and the heat treatment shown in Tables 2 and 3 was performed either before or after immersion in the galvanizing bath, or both. In Tables 2 and 3, steel sheets subjected to continuous hot dip galvanization are indicated as GI.

  Furthermore, about the other one part steel plate, it was set as the alloyed hot-dip galvanized steel plate by making it pass through a continuous hot-dip galvanizing line and alloying before and after said heat processing. In the continuous hot dip galvanizing line, after annealing under the conditions shown in Table 2 or Table 3, after cooling between 630 ° C and 570 ° C to 430 ° C to 500 ° C at an average cooling rate or more shown in Table 2 or Table 3, Hot dip galvanization was performed, and an alloying treatment was further performed at a temperature of 500 ° C to 550 ° C. And the heat processing shown in Table 2 and Table 3 was performed either before or after immersion in the galvanizing bath, after alloying treatment, or all. The steel plates that were subjected to continuous hot dip galvanizing and alloyed were indicated as GA in Tables 2 and 3.

  In Table 2 and later, for example, steel No. A-1 to A-14 are examples using steel A having the composition shown in Table 1. Hereinafter, the same applies to the other steels B to I.

  About the obtained cold-rolled steel sheet CR, hot-dip galvanized steel sheet GI, and alloyed hot-dip galvanized steel sheet GA, the Mn segregation degree and the thickness of the Mn segregation zone were measured. The Mn segregation degree was obtained by conducting elemental analysis on the cross section of the steel sheet with an EPMA apparatus and calculating the Mn segregation degree in the thickness direction (= center Mn peak concentration / average Mn concentration). Furthermore, the thickness of the Mn segregation zone was also measured. The Mn segregation zone was defined as a region showing a Mn concentration higher than the average Mn concentration. The results are shown in Tables 4 and 5.

  Moreover, in order to investigate the presence or absence of a strength difference between the central segregation part and the other part, the hardness difference of Vickers hardness was measured. In the measurement, the hardness at the central segregation part and other parts was determined by a Vickers tester with a load of 50 gf, and the difference was defined as the hardness difference. The results are shown in Tables 4 and 5.

  Further, identification of each phase of the microstructure, ferrite, pearlite, cementite, martensite, bainite, austenite and the remaining structure, observation of the existing position, and measurement of the area ratio are disclosed in Nital reagent and Japanese Patent Application Laid-Open No. 59-219473. The steel sheet was eroded with the reagent in the rolling direction of the steel plate and quantified by observation with an optical microscope of 1000 times and scanning and transmission electron microscopes of 1000 to 100,000 times. The results are shown in Tables 4 and 5.

In addition, the obtained cold-rolled steel sheet CR, hot-dip galvanized steel sheet GI, and alloyed hot-dip galvanized steel sheet GA are subjected to a tensile test to measure yield stress (YS), maximum tensile stress (TS), and total elongation (El). did. In addition, this steel plate is a composite structure steel plate which consists of a ferrite and a hard structure, and yield point elongation does not appear in many cases. From this, the yield stress was measured by the 0.2% offset method. A steel sheet having a TS × El of 16000 (MPa ×%) or more was defined as a high-strength steel sheet having a good strength-ductility balance.
The hole expansion rate (λ) was evaluated by punching a circular hole having a diameter of 10 mm under the condition that the clearance was 12.5%, forming the burr on the die side, and molding with a 60 ° conical punch. Under each condition, five hole expansion tests were performed, and the average value was defined as the hole expansion ratio. A steel sheet having a TS × λ of 40000 (MPa ×%) or more was defined as a high-strength steel sheet having a good strength-hole expansibility balance.
The results are shown in Tables 6 and 7.

Further, when the center segregation of Mn is large and the clearance is large, cracks (peeling) may be formed on the punched end face or the cut end face.
In addition, since the characteristics deteriorated when a punched test piece was processed under the conditions that cause cracks, damage to the punched end was also evaluated as an item for evaluating the deterioration of characteristics after cutting. End face damage after cutting varies depending on the cutting direction. From this, in order to simultaneously evaluate the damage on the entire cross section, the end face was examined for damage by punching with a 10 mmφ punch. The total length of the cracks after punching and the secondary shear surface is expressed as the total length of the individual cracks with respect to the entire circumference of 360 °. Crack formation was prominent in steel sheets with large center segregation and a large hardness difference in the sheet thickness direction, and cracks were easily formed along a plane parallel to the rolling direction, although the detailed reason was unknown.
Although the same tendency can be obtained by cutting, the ease of cracking changes depending on the cutting direction. Therefore, in the present invention, end face damage was evaluated by a punching test. The following scoring was performed based on the total number of cracks formed in the punched end face.

○: 120 ° or less.
Δ: More than 120 to 270 ° or less.
X: Over 270 °.

  In order to compare the punching clearance with the characteristics after cutting, the cutting according to the present invention is such that the crack of the punched end face is 120 ° or less under any conditions of clearance 12.5%, 25.0%, 37.5%. A high-strength steel sheet with less characteristic deterioration later was obtained. The results are shown in Tables 6 and 7.

  In the present invention, in order to evaluate the change in stretch flangeability after cutting or punching, the following molding technique was used. Since the stretch flange portion of the actual member is subjected to non-uniform deformation, it is required that the stretch (uniform stretch) is excellent for suppressing breakage at the stretch flange portion. Further, there are many cases where deformation is concentrated on a specific part, and it is necessary to alleviate strain concentration by improving n value (uniform elongation). On the other hand, since the processed end portion is easy to form a crack, it is required that the end portion typified by a hole expansion test using a conical punch is difficult to form a crack. However, since the hole expansion test uniformly deforms around the punched hole, there is a problem that it is difficult to evaluate the stretch flangeability in the actual member, such as non-uniform deformation that occurs in the actual member and crack formation in the strain concentration part. Was. Therefore, in the present invention, the following test method was used.

  Specifically, a molding experiment was performed as shown in FIGS. 1A to 1C, in which a base plate shape of a hole expansion test using a cylindrical punch that is normally performed was divided to simulate stretch flange molding. A die 1 having a shoulder R of 5 mm and a diameter of 106 φ was used. After being restrained by the crease presser 2, the die 1 was molded using a 100φ cylindrical flat bottom punch 3 having a shoulder R of 10 mm and a corner radius of 50 mm.

  The base plate used for forming was obtained by cutting a steel plate having a thickness of 1.2 mm into 180 mm squares and then punching it with a 40φ punch. At this time, in order to investigate the characteristic deterioration accompanying the clearance change at the time of punching, the clearance was variously changed to 12.5 and 25.0.37.5%, and the influence of the clearance on the stretch flangeability after punching was investigated. . Then, it cut | disconnected to 1/2 and it was set as the test piece 4 for a stretch flange test (FIG. 2).

  In this molding method, when the molding height 5 by the punch shown in FIG. 1C is increased, the hole bottom at the end of the test piece 4 is deformed. As a result, a material with inferior ductility, particularly uniform elongation, causes a large strain concentration at the bottom of the hole even at a low molding height 5, leading to fracture. On the other hand, since the material with poor hole expandability is easily cracked, even if the strain concentration at the hole bottom is small, it breaks. From this, it is possible to evaluate the stretch flangeability of an actual member with non-uniform strain concentration and crack formation at the end. In this test, it is possible to evaluate the forming height 5 at which a crack penetrating the plate thickness is formed at the hole bottom as an index of stretch flangeability. That is, the fact that a high molding height can be obtained without occurrence of cracks means that the flange height is high or stretch flange processing with large deformation is possible.

  In the present invention, in order to see the influence of the punching conditions, the clearance is changed to 12.5%, 25.0%, 37.5%, and each clearance with respect to the molding height (h12.5) of the clearance 12.5%. The molding height ratio was evaluated.

  Molding height ratio at clearance 25.0% = h (25.0%) / h (12.5%)

  Molding height ratio at clearance 37.5% = h (37.5%) / h (12.5%)

  In both of the clearances of 25.0% and 37.5%, those satisfying the following range were defined as high-strength steel plates with little characteristic deterioration after cutting. The results are shown in Tables 6 and 7. The high-strength steel sheet with less characteristic deterioration after cutting according to the present invention is less susceptible to ductility deterioration after cutting and hole expansion value due to a hole expansion test after punching.

A: 0.9 or more.
○: 0.8 or more and less than 0.9.
X: Less than 0.8.

  As shown in Tables 4 to 7, steel numbers A-1 to 3, 6, 7, 9, 11, B-1, 3, 5, 6, C-1, 3, D-1, 3, 5, E -1, 2, F-1, G-1, 2, H-1, I-1, 4, and 6 are within the range specified by the present invention for the chemical composition of the steel sheet, and the manufacturing conditions are also the same. It is within the range specified by the invention. As a result, the stretch flangeability hardly changes even if the ductility and the hole expandability are simultaneously provided and the clearance at the time of punching is changed in a wide range of 12.5 to 37.5%. As a result, it is possible to produce a high-strength steel sheet having a maximum tensile strength of 540 MPa or more and having little characteristic deterioration after cutting or punching.

  On the other hand, as shown in Tables 4-7, steel numbers A-4, 10, 12, B-2, 7, C-2, D-2, 4, 6, E-3, F-2, G-3 4, H-2, I-2, 5, and 7 have a reduction of 0 mm during casting. Therefore, center segregation, in particular, Mn segregation degree and segregation band thickness are large, and punching is performed with a large clearance. In this case, cracks are easily formed on the end face after punching, and the characteristic deterioration after punching with a large clearance is large. That is, the characteristic deterioration after punching is large.

  Steel numbers A-5, 13, B-4, G-5, H-3, and I-3 are cold-rolled steel sheets, and the residence time in the temperature range of 300 to 450 ° C. is less than 30 seconds. Therefore, if it is a hot dip galvanized steel sheet and an alloyed hot dip galvanized steel sheet, the residence time in the temperature range of (zinc plating bath temperature +50) ° C. to 300 ° C. is less than 30 seconds. A lot of martensite appears and the hole expandability is inferior.

  Furthermore, Steel No. A-8 has an annealing temperature as low as 720 ° C., and the pearlite structure formed during hot rolling and the cementite that has been spheroidized remain in the steel sheet structure. Therefore, bainite and martensite which are hard structures are present. Since a sufficient volume ratio cannot be secured, a high strength of 540 MPa or more cannot be secured. In addition, the strength-ductility balance is poor.

  Steel No. A-14 has too low a cooling rate in the temperature range of 630 to 570 ° C., and thus austenite is transformed into pearlite, so that a high strength of 540 MPa or more cannot be secured. In addition, the strength-ductility balance is poor.

  Steel numbers J-1 and 2 are low in Si and Mn at 0.03% and 1.14%, respectively, suppress pearlite transformation in the cooling process after annealing, and have a hard structure such as bainite, martensite, and retained austenite. Therefore, it is impossible to secure a high strength of 540 MPa or more.

Steel number K-1 cannot secure a high strength of 540 MPa or more because the C content is as low as 0.027% and a sufficient amount of hard structure cannot be secured.
Steel numbers L-1 to L-3 have a high Mn content of 3.28%, and once the austenite volume fraction is reduced during annealing, a sufficient amount of ferrite cannot be produced in the cooling process. For this reason, the strength-ductility balance is remarkably inferior.

  The present invention has a maximum tensile strength of 540 MPa or more suitable for structural members, reinforcing members, and suspension members for automobiles, has good ductility and hole expansibility, and has little deterioration in properties after cutting. The steel sheet is provided at a low cost, and since this steel sheet is suitable for use in, for example, structural members for automobiles, reinforcing members, suspension members, etc., it can be expected to greatly contribute to weight reduction of automobiles. Industrial effect is extremely high.

It is a figure which shows the evaluation apparatus for evaluating the change of the stretch flange formability in an Example, (a) is a cross-sectional schematic diagram which shows the state before a test start, (b) is the state before a test start. It is a plane schematic diagram to show, (c) is a cross-sectional schematic diagram which shows the state after a test. It is a plane schematic diagram which shows the test piece used for evaluation of the change of the stretch flangeability in an Example.

Claims (11)

% By mass
C: 0.05% to 0.20%,
Si: 0.3 to 2.00%
Mn: 1.3-2.6%,
P: 0.001 to 0.03%,
S: 0.0001 to 0.01%
Al: less than 0.10%,
N: 0.0005 to 0.0100%,
O: 0.0005 to 0.007%
The balance is steel composed of iron and inevitable impurities, the steel sheet structure is mainly composed of ferrite and bainite, and the Mn segregation degree in the thickness direction (= center Mn peak concentration / average Mn concentration) is 1.20. A high-strength steel sheet having less characteristic deterioration after cutting, wherein the maximum tensile strength is 540 MPa or more. Furthermore, in mass%,
B: The high-strength steel sheet with less property deterioration after cutting according to claim 1, characterized by containing 0.0001% or more and less than 0.01%. Furthermore, in mass%,
Cr: 0.01 to 1.0%,
Ni: 0.01 to 1.0%,
Cu: 0.01 to 1.0%,
Mo: 0.01 to 1.0%
The high-strength steel sheet with little characteristic deterioration after cutting according to claim 1 or 2, characterized by containing one or more of the following.   The cutting according to any one of claims 1 to 3, further comprising 0.001 to 0.14% of Nb, Ti, or V in total by mass, containing one or more of Nb, Ti, and V. A high-strength steel plate with less characteristic deterioration afterwards.   Furthermore, 0.0001-0.5% of 1 type or 2 types in total of Ca, Ce, Mg, and REM is contained by mass%, The any one of Claims 1-4 characterized by the above-mentioned. High-strength steel sheet with little deterioration in properties after cutting.     A high-strength steel sheet with little deterioration in characteristics after cutting, characterized by having zinc-based plating on the surface of the high-strength steel sheet according to any one of claims 1 to 5.   6. A cast slab, wherein the casting slab having the chemical component according to claim 1 is cast while being reduced in the thickness direction.   The cast slab according to claim 7 is directly or once cooled and then heated to 1050 ° C or higher and subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature is 820 ° C to a reduction rate of 85% or more. After performing hot rolling to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is completed at 620 to 720 ° C. After winding in a temperature range of 400 to 630 ° C, pickling, cold rolling with a rolling reduction of 40 to 70%, and passing through a continuous annealing line, after annealing at a maximum heating temperature of 760 to 870 ° C, 630 A method for producing a high-strength steel sheet with less property deterioration after cutting, characterized in that the temperature is cooled between 3 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more and is maintained in a temperature range of 450 ° C. to 300 ° C. for 30 seconds or more .   The cast slab according to claim 7 is directly or once cooled and then heated to 1050 ° C or higher and subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature is 820 ° C to a reduction rate of 85% or more. After performing hot rolling to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is completed at 620 to 720 ° C. After winding in a temperature range of 400 to 630 ° C., pickling, cold rolling with a rolling reduction of 40 to 70%, and annealing at a maximum heating temperature of 760 to 870 ° C. when passing through a continuous hot dip galvanizing line After cooling between 630 ° C. and 570 ° C. at an average cooling rate of 3 ° C./second or more (zinc plating bath temperature−40) ° C. to (zinc plating bath temperature + 50) ° C., before or after immersion in the zinc plating bath Either later or both, Method for producing a less high strength galvanized steel sheet characteristic degradation after cutting, characterized in that the holding in a temperature range of galvanizing bath temperature +50) ℃ ~300 ℃ 30 seconds or more.   The cast slab according to claim 7 is directly or once cooled and then heated to 1050 ° C or higher and subjected to hot rolling to a reduction rate of 70% or more, and then the finishing temperature is 820 ° C to a reduction rate of 85% or more. After performing hot rolling to 930 ° C, water cooling is started, water cooling is performed at an average cooling rate of 720 to 800 ° C at a cooling rate of 25 ° C / second or more, and water cooling is completed at 620 to 720 ° C. After winding in a temperature range of 400 to 630 ° C., pickling, cold rolling with a rolling reduction of 40 to 70%, and annealing at a maximum heating temperature of 760 to 870 ° C. when passing through a continuous hot dip galvanizing line After cooling from 630 ° C to 570 ° C at an average cooling rate of 3 ° C / second or more to (zinc plating bath temperature -40) ° C to (zinc plating bath temperature +50) ° C, a temperature of 460 to 540 ° C is necessary. Is alloyed and immersed in a galvanizing bath. Characteristic deterioration after cutting, characterized by holding for at least 30 seconds in a temperature range of (zinc plating bath temperature +50) ° C. to 300 ° C. either before, after immersion, after alloying treatment or at all A method for producing few high-strength galvannealed steel sheets.   9. The method for producing a high-strength steel sheet with less characteristic deterioration after cutting according to claim 8, wherein the high-strength steel sheet is produced by the method of claim 8 and then zinc-based electroplating is performed.

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