JP5533765B2 - High-strength cold-rolled steel sheet with excellent local deformability and its manufacturing method - Google Patents

Description

  The present invention relates to a high-strength cold-rolled steel sheet excellent in local deformability such as bending, stretch flange, and burring, and is mainly used for automobile parts.

  In order to reduce carbon dioxide emissions from automobiles, the weight of automobile bodies is being reduced using high-strength steel sheets. In addition, in order to ensure the safety of passengers, high strength steel plates are often used in automobile bodies in addition to mild steel plates. Furthermore, in order to further reduce the weight of automobile bodies in the future, it is necessary to increase the use strength level of high-strength steel sheets more than before. For example, to use high-strength steel sheets for undercarriage parts, local parts for burring processing The deformability must be improved.

  However, generally, if the strength of the steel plate is increased, the formability is lowered, and as in Non-Patent Document 1, the uniform elongation that is important for drawing and stretch forming is reduced. On the other hand, as disclosed in Non-Patent Document 2, a method is disclosed in which uniform elongation is ensured even with the same strength by compounding the metal structure of a steel plate.

  On the other hand, a metal structure control method for a steel sheet that improves local ductility, represented by bending, hole expanding, and burring, is also disclosed. Non-Patent Document 3 discloses that if the hardness difference is reduced, it is effective for bendability and hole expansion.

  This is to improve the hole expansion property by making a single structure by controlling the structure, but in order to make a single structure, heat treatment from an austenite single phase as in Non-Patent Document 4 is a manufacturing method. Basic. Further, Non-Patent Document 4 also discloses a technique for obtaining an appropriate fraction of ferrite and bainite by controlling the metal structure by cooling control after hot rolling for compatibility with ductility, and controlling the precipitate and the transformation structure. There is disclosure.

  On the other hand, strength and ductility are controlled by controlling the finishing temperature of hot rolling, the rolling reduction and temperature range of finishing rolling, promoting the recrystallization of austenite, suppressing the development of rolling texture, and randomizing the crystal orientation. Patent Document 1 discloses a technique for improving hole expansibility.

JP 2009-263718 A

Kishida, Nippon Steel Technical Report (1999) No.371, p.13 O. Matsumura et al, Trans. ISIJ (1987) vol. 27, p. 570 Kato et al., Steel Research (1984) vol.312, p.41 K. Sugimoto et al, ISIJ International (2000) Vol.40, p.920

As described above, the factors that degrade the local deformability are various “non-uniformities” such as inter-structure hardness differences, non-metallic inclusions, and developed rolling texture. Among them, the one having the greatest influence is the hardness difference between the structures shown in Non-Patent Document 3 above, and the other developed dominant texture is the developed rolling texture shown in Patent Document 1. .
Since these elements are entangled in a complex manner and the local deformability of the steel sheet is determined, in order to maximize the increase in local deformability due to texture control, the structure control is performed together and the hardness difference between structures It is necessary to eliminate the resulting non-uniformity as much as possible.

  Therefore, in the present invention, it is possible to improve the local ductility of the high-strength steel sheet and improve the anisotropy in the steel sheet together with the texture control by making the bainite area ratio 95% or more. A high-strength cold-rolled steel sheet having excellent local deformability and a method for producing the same are provided.

  According to the conventional knowledge, as described above, improvement of hole expansibility and bendability has been achieved by inclusion control, refinement of precipitates, homogenization / single phase structure of the structure, and reduction of hardness difference between structures. It was. However, with this alone, there is a concern about the influence on anisotropy in a high-strength steel sheet to which Nb, Ti, or the like is added. This causes problems such as sacrificing other formability factors and limiting the direction in which the blank before molding is taken, and the application is also limited.

Therefore, the present inventors have investigated and studied the effects in detail, focusing attention on the influence of the texture of the steel sheet newly in order to improve the hole expandability and the bending workability. As a result, it was clarified that by controlling the strength of each orientation of a specific crystal orientation group, the local deformability was dramatically improved without greatly reducing the elongation and strength. The point to be emphasized is that the improvement cost of the local deformability by the texture control is largely dependent on the steel structure, and after ensuring the strength of the steel by making the bainite area ratio 95% or more, It is also clarified that the improvement cost of local deformability is maximized.
In addition, it has been found that the size of each grain unit greatly affects the local ductility in a structure in which the strength of each orientation of a specific crystal orientation group is controlled.

In general, in a structure in which low-temperature generation phases (bainite, martensite, etc.) are mixed, the definition of crystal grains is very vague and difficult to quantify. On the other hand, the present inventors have found that the problem of quantification can be solved by using a grain unit measured as follows.
That is, the grain unit as used in the present invention is measured by measuring the orientation of a steel sheet by EBSP (Electron Back Scattering Pattern), for example, at a magnification of 1500 times and a measurement step of 0.5 μm or less. The position where the azimuth difference between adjacent measurement points exceeds 15 ° is defined as the grain boundary of the grain unit. Then, the equivalent circle diameter d of the measured grain unit is obtained, the individual volume is obtained by 4 / 3πd ³ , and the volume average diameter can be obtained by weighted average of the volume. Grains can be quantified and evaluated.

The present invention is configured based on the above-mentioned knowledge, and the main points thereof are as follows.
(1) In mass%,
C: 0.002% or more, 0.20% or less,
Si: 0.001% or more, 2.5% or less,
Mn: 0.01% or more and 4.0% or less,
P: 0.001% or more, 0.15% or less,
S: 0.0005% or more, 0.03% or less,
Al: 0.001% or more, 2.0% or less,
N: 0.0005% or more, 0.01% or less,
O: 0.0005% or more, 0.01% or less,
Si + Al: containing less than 1.0%, consisting of the balance iron and inevitable impurities,
The area ratio of bainite in the metal structure of the steel sheet is 95% or more, and {100} <011> to {223 of the plate surface at a thickness of at least 5/8 to 3/8 from the surface of the steel sheet in the texture of the steel sheet. } The average value of the X-ray random intensity ratio of the <110> orientation group is less than 4.0, the X-ray random intensity ratio of the crystal orientation of {332} <113> is 5.0 or less, and the metal of the steel plate A high-strength cold-rolled steel sheet having excellent local deformability, wherein the volume average size of grain units in the structure is 7 μm or less.

(2) Further, the ratio of the length dL in the rolling direction to the length dt in the plate thickness direction of the grain unit of bainite, and the ratio of grains having dL / dt of 3.0 or less is 50% or more. A high-strength cold-rolled steel sheet excellent in local deformability as described in (1) above.
(3) Furthermore, in mass%,
Ti: 0.001% or more, 0.20% or less,
Nb: 0.001% or more, 0.20% or less,
V: 0.001% or more, 1.0% or less,
W: A high-strength cold-rolled steel sheet having excellent local deformability as described in (1) or (2) above, containing one or more of 0.001% or more and 1.0% or less.
(4) Furthermore, in mass%,
B: 0.0001 or more and 0.0050% or less Mo: 0.001 or more and 1.0% or less,
Cr: 0.001 or more and 2.0% or less,
Cu: 0.001 or more, 2.0% or less,
Ni: 0.001 or more and 2.0% or less,
Co: 0.0001 or more and 1.0% or less,
Sn: 0.0001 or more and 0.2% or less,
Zr: 0.0001 or more and 0.2% or less,
As: A high-strength cold-rolled steel sheet excellent in local deformability according to any one of the above (1) to (3), which contains one or more of 0.0001 or more and 0.50% or less.
(5) Furthermore, in mass%,
Mg: 0.0001 or more and 0.010% or less,
REM: 0.0001 or more and 0.1% or less,
Ca: A high-strength cold-rolled steel sheet excellent in local deformability according to any one of (1) to (4), which contains one or more of 0.0001 or more and 0.010% or less.
(6) The surface of the high-strength cold-rolled steel sheet according to any one of (1) to (5) is provided with a hot-dip galvanized layer or an alloyed hot-dip galvanized layer, and has excellent local deformability. High strength cold rolled steel sheet.

(7) In producing the high-strength cold-rolled steel sheet according to any one of (1) to (5),
After melting into a predetermined steel plate component, it is cast into a steel ingot or slab and subjected to rough rolling at a temperature range of 1000 ° C. or higher and 1200 ° C. or lower and at least once at least 20%, and austenite grain size When the temperature determined by the steel plate component in (Equation 1) in finish rolling is T1 ° C., the total rolling reduction in the temperature range of T1 + 30 ° C. or higher and T1 + 200 ° C. or lower is 50% or higher. Thereafter, the total rolling reduction in the temperature range of T1 ° C. or more and less than T1 + 30 ° C. is set to 30% or less, and then the hot rolling is finished at the Ar3 transformation temperature or more, and T1 ° C. or more and T1 + 30 ° C. or less after the final pass of the reduction. After the retention time t (seconds) in the temperature range of (2) satisfies (Equation 2), cooling is performed, followed by pickling, and rolling at 30% to 70% in the cold. After annealing for 1 to 300 seconds in the temperature range of Ae3 to 950 ° C, the average cooling rate in the temperature range of Ae3 to 500 ° C is set to 10 ° C / s or more and 200 ° C / s or less, and further 350 ° C or more, A method for producing a high-strength cold-rolled steel sheet excellent in local deformability, characterized by holding for at least t2 seconds satisfying (Equation 4) at an overaging temperature of 500 ° C. or less. However, the maximum value of t2 is 400 seconds.
T1 (℃) = 850 + 10 × (C + N) × Mn + 350 × Nb + 250 × Ti + 40 × B + 10 × Cr + 100 × Mo + 100 × V
... (Formula 1)
t <t1 (Formula 2)
Here, t1 is expressed by (Expression 3).
t1 = 0.001 ((Tf−T1) × P1) ² −0.109 ((Tf−T1) × P1) +3.1 (Equation 3)
Here, Tf and P1 are a temperature and a reduction ratio at the time of final reduction in a temperature range of T1 + 30 ° C. or higher and T1 + 200 ° C. or lower, respectively.
log (t2) = 0.0002 (T2-425) ² + 1.18 (Formula 4)
Here, T2 is an overaging treatment temperature.

(8) In the manufacturing method according to the above (7), the rolling ratio of the final pass of rolling in a temperature range of T1 + 30 ° C. or higher and T1 + 200 ° C. or lower is 25% or more, and has high local deformability. A method for producing a high strength cold-rolled steel sheet.
(9) In the manufacturing method according to (7) or (8) above, after the hot rolling is finished at an Ar3 transformation temperature or higher, the cooling temperature change is 40 ° C or higher within t seconds represented by (Formula 2). A method for producing a high-strength cold-rolled steel sheet having excellent local deformability, characterized in that rapid cooling is started immediately after the cooling rate becomes 50 ° C./s or higher at 150 ° C. or lower.
(10) A hot-dip galvanized layer or an alloyed hot-dip galvanized layer is formed on the surface of the cold-rolled steel sheet produced by the production method according to any one of (7) to (9) above. A method for producing a high-strength cold-rolled steel sheet having excellent local deformability.

  Note that the grain unit determines the position where the orientation difference between adjacent measurement points exceeds a certain angle in the analysis of the orientation of the steel sheet by EBSP (Electron Back Scattering Pattern) as the grain boundary of the grain unit.

  According to the present invention, by controlling the texture and steel structure of a steel sheet, a high-strength cold-rolled steel sheet having excellent local deformability such as bending, stretch flange, and burring is obtained.

The relationship between the average value of the X-ray random intensity ratio of {100} <011> to {223} <110> orientation group and the thickness / minimum bending radius is shown. The relationship between the X-ray random intensity ratio of the {332} <113> orientation group and the thickness / minimum bending radius is shown. The relationship between the rolling frequency | count of 40% or more in rough rolling and the austenite grain size of rough rolling is shown. The relationship between the rolling reduction of T1 + 30 to T1 + 200 ° C. and the average value of the X-ray random intensity ratio of {100} <011> to {223} <110> orientation groups is shown. The relationship between the rolling reduction of T1 + 30-T1 + 200 degreeC and the X-ray random intensity ratio of the crystal orientation of {332} <113> is shown. The relationship between the strength of the invention steel and the comparative steel and the hole expandability is shown. The relationship between the strength and bendability of the steel of the present invention and the comparative steel is shown.

The contents of the present invention will be described in detail below.
First, the X-ray random intensity ratio of the plate surface at a thickness of 5/8 to 3/8 from the surface of the cold rolled steel plate will be described.
In the cold-rolled steel sheet of the present invention, the average value of the X-ray random intensity ratios of {100} <011> to {223} <110> orientation groups on the plate surface at a thickness of 5/8 to 3/8 from the surface of the steel plate. , And {332} <113> crystal orientation X-ray random intensity ratios are particularly important characteristic values.

  As shown in FIG. 1, {100} <011> to {100 when the X-ray diffraction of the plate surface at a thickness of 5/8 to 3/8 from the surface of the steel plate is performed and the intensity ratio of each orientation with respect to the random sample is obtained. The average value of the 223} <110> azimuth group is less than 4.0, and it is possible to satisfy the plate thickness / bending radius ≧ 1.5 required for processing of the skeleton part that is required most recently. In addition, when the steel structure satisfies the requirements of claim 1, that is, the bainite fraction of 95% or more, the thickness / bending radius ≧ 2.5 is satisfied. When hole expansibility and small critical bending characteristics are required, the average value of {100} <011> to {223} <110> orientation groups is preferably less than 3.0.

  If this value is 4.0 or more, the anisotropy of the mechanical properties of the steel sheet becomes extremely strong, which in turn improves the local deformability only in one direction, but the material in a direction different from that significantly deteriorates, and the thickness / bending radius. ≧ 1.5 cannot be satisfied. On the other hand, it is difficult to realize by the current general continuous hot rolling process, but when it is less than 0.5, there is a concern about deterioration of local deformability.

The main orientations included in this orientation group are {100} <011>, {116} <110>, {114} <110>, {113} <110>, {112} <110>, {335} <110> and {223} <110>.
The X-ray random intensity ratio in each direction is measured using an apparatus such as X-ray diffraction or EBSD (Electron Back Scattering Diffraction). {110} Using a three-dimensional texture calculated by the vector method based on a pole figure, or a plurality of pole figures (preferably three or more) among {110}, {100}, {211}, {310} pole figures What is necessary is just to obtain | require from the three-dimensional texture calculated by the series expansion method.

For example, the X-ray random intensity ratio of the respective crystal orientations in the latter method, in phi _{2 =} 45 ° cross-section of the three-dimensional texture (001) [1-10], (116) [1-10], ( 114) [1-10], (113) [1-10], (112) [1-10], (335) [1-10], (223) [1-10] strengths can be used as they are. .
The average value of {100} <011> to {223} <110> orientation group is an arithmetic average of each of the above-mentioned orientations. When the strengths of all the above directions cannot be obtained, {100} <011>, {116} <110>, {114} <110>, {112} <110>, {223} <110> Alternatively, an arithmetic average of each direction may be substituted. (1 with an upper bar representing “minus 1” is represented by “−1”.)

  Furthermore, for the same reason, the X-ray random intensity ratio of {332} <113> crystal orientation of the plate surface at 5/8 to 3/8 plate thickness from the surface of the steel plate is not less than 5.0 as shown in FIG. must not. Desirably, if it is 3.0 or less, the plate thickness / bending radius ≧ 1.5 required for the processing of the most recently required skeleton parts is satisfied. In addition, when the steel structure satisfies the requirement of claim 1, that is, the bainite fraction is 95% or more, the thickness / bending radius ≧ 2.5 is satisfied. On the other hand, if the X-ray random intensity ratio of the {332} <113> crystal orientation exceeds 5.0, the anisotropy of the mechanical properties of the steel sheet becomes extremely strong, which improves the local deformability only in a certain direction. However, the material in a direction different from that is significantly deteriorated, and the thickness / bending radius ≧ 1.5 cannot be surely satisfied. On the other hand, it is difficult to realize by the current general continuous hot rolling process, but when it is less than 0.5, there is a concern about deterioration of local deformability.

  The reason why the X-ray intensity of the crystal orientation described above is important for the shape freezing property during bending is not necessarily clear, but it is presumed to be related to the sliding behavior of the crystal during bending deformation. .

A sample subjected to X-ray diffraction is obtained by reducing the thickness of a steel sheet from a surface to a predetermined thickness by mechanical polishing or the like, and then removing distortion by chemical polishing or electrolytic polishing and at the same time having a thickness of 3/8 to 5/8. What is necessary is just to adjust and measure a sample according to the above-mentioned method so that a suitable surface may become a measurement surface in the range.
As a matter of course, the above-mentioned limitation of the X-ray intensity is satisfied not only in the vicinity of the plate thickness ½ but also as much as possible, so that the spread performance is further improved.
However, the measurement of 3/8 to 5/8 from the surface of the steel sheet can generally represent the material properties of the entire steel sheet, so this shall be specified. The crystal orientation represented by {hkl} <uvw> indicates that the normal direction of the plate surface is parallel to <hkl> and the rolling direction is parallel to <uvw>.

Then, the grain unit of a steel plate is described.
As a result of earnest examination of the texture control in the cold-rolled steel sheet, the inventors of the present invention have an extremely large influence on the local ductility under the condition that the texture is controlled as described above. It was found that a dramatic improvement in local ductility can be obtained.
The reason for this is not clear, but randomizing the texture of the hot-rolled steel strip and refinement of the grain units suppress local strain concentration that occurs on the micro order, and increase the homogenization of deformation to increase the micro locality. This is because the strain concentration can be suppressed and the strain can be uniformly dispersed even in the micro order.

At this time, as for the contribution of the grain unit, even if the number is small, the larger the grain unit, the greater the deterioration of local ductility. For this reason, local ductility and a strong correlation are obtained when the size of a grain unit is not a normal size average but a volume average defined by a weighted average of volumes. In order to obtain this effect, it is necessary that the volume average particle size is 7 μm or less. Further, in order to ensure the hole expandability at a high level, 5 μm or less is desirable.
In addition, it is as having mentioned above about the measuring method of a grain unit.

  As a result of further pursuing local ductility, the present inventors have also found that the local ductility is improved when the grain unit is excellent in equiaxedness after satisfying the above texture and grain unit. As an index representing this equiaxed property, in a grain represented by a grain unit, the ratio of the length dL in the cold rolling direction and the length dt in the sheet thickness direction of the grain, dL / dt is 3.0 or less. At least 50% or more of all bainite grains are required for the proportion of grains having excellent equiaxedness. If it is less than 50%, the local ductility deteriorates.

Subsequently, the limiting conditions of the components will be described. In addition,% of content is the mass%.
C has a lower limit of 0.02% in order to make 95% or more of the steel structure bainite. Further, since C is an element that increases the strength, it is preferably 0.025% or more for securing the strength. On the other hand, if the amount of C exceeds 0.20%, the weldability may be impaired, or the workability may be extremely deteriorated due to an increase in the hard structure, so the upper limit is made 0.20%. Further, if the C content exceeds 0.10%, the moldability deteriorates, so the C content is preferably 0.10% or less.

  Si is an effective element for increasing the mechanical strength of the steel sheet, but if it exceeds 2.5%, workability deteriorates or surface flaws occur, so this is the upper limit. Moreover, since chemical conversion processability will fall when there is much Si amount, it is preferable to set it as 1.20% or less. On the other hand, since it is difficult to make Si less than 0.001% in practical steel, this is the lower limit.

  Mn is also an effective element for increasing the mechanical strength of the steel sheet, but if it exceeds 4.0%, workability deteriorates, so this is the upper limit. On the other hand, since it is difficult to make Mn less than 0.01% in practical steel, this is the lower limit. In addition to Mn, when an element such as Ti that suppresses the occurrence of hot cracking due to S is not sufficiently added, it is desirable to add an amount of Mn that satisfies Mn / S ≧ 20 by mass%. Furthermore, Mn is an element that expands the austenite temperature to the low temperature side with an increase in the content thereof, improves the hardenability, and facilitates the formation of a continuous cooling transformation structure having excellent burring properties. Since this effect is hardly exhibited when the Mn content is less than 1%, it is desirable to add 1% or more.

  The upper limits of P and S are 0.15% or less for P and 0.03% or less for S, respectively. This is to prevent workability deterioration and cracking during hot rolling or cold rolling. The lower limit was set to 0.001% for P and 0.0005% for S as possible values for P and S by current general refining (including secondary refining).

  Al is added in an amount of 0.001% or more for deoxidation. When deoxidation is sufficiently necessary, addition of 0.01% or more is preferable. Al is also an element that significantly increases the γ → α transformation point. However, if the amount is too large, the weldability becomes poor, so the upper limit is made 2.0%. Preferably, it is 1.0% or less.

  N and O are impurities, and both are made 0.01% or less so as not to deteriorate the workability. The lower limit was set to 0.0005%, which is possible for both elements by current general refining (including secondary refining). However, 0.001% or more is preferable in order to suppress an extreme increase in steelmaking cost.

  If Si and Al are contained excessively, cementite precipitation during the overaging treatment is suppressed and the retained austenite fraction becomes too large, so the total amount of Si and Al added is less than 1%.

Furthermore, when obtaining strength by precipitation strengthening, fine carbonitrides are preferably generated. In order to obtain precipitation strengthening, addition of Ti, Nb, V, and W is effective, and one or more of these may be contained.
In order to obtain this effect by adding Ti, Nb, V, and W, the addition of 0.001% for Ti, 0.001% for Nb, 0.001% or more for V, and 0.001% or more for W should be added. is necessary. When precipitation strengthening is particularly necessary, it is desirable to add Ti 0.01% or more, Nb 0.005% or more, V 0.01% or more, and W 0.01% or more. However, even if it is excessively added, the increase in strength is saturated, and in addition, it is difficult to control the crystal orientation after cold rolling annealing by suppressing recrystallization after hot rolling. Hereinafter, it is necessary that Nb is 0.20% or less, V is 1.0% or less, and W is 1.0% or less.

  Addition of one or more of B, Mo, Cr, Cu, Ni, Co, Sn, Zr, As is effective when strength is ensured by increasing the hardenability of the structure and controlling the second phase It is. In order to obtain this effect, it is necessary to add 0.0001% or more of B, 0.001% or more of Mo, Cr, Cu, and Ni, and 0.0001% or more of Co, Sn, Zr, and As. However, excessive addition adversely degrades workability, so the upper limit of B is 0.0050%, the upper limit of Mo is 1.00%, the upper limit of Cr, Cu, Ni is 2.0%, and the upper limit of Co is 1.0%, the upper limit of Sn and Zr is 0.2%, and As is 0.50%.

  In order to improve local forming ability, Mg, REM, and Ca are important additive elements for detoxifying inclusions. Each lower limit for obtaining this effect was made 0.0001%. On the other hand, excessive addition leads to deterioration of cleanliness, so 0.010% for Mg, 0.1% for REM, and 0.010% for Ca were made the upper limit.

  Even if the surface treatment is applied to the high-strength cold-rolled steel sheet of the present invention, the effect of improving the local deformability is not lost. Electroplating, hot dipping, vapor deposition plating, organic film formation, film lamination, organic salt / inorganic salt treatment The effect of the present invention can be obtained by any of non-chromic treatment and the like.

Next, the metal structure of the cold rolled steel sheet of the present invention will be described.
The metal structure of the cold-rolled steel sheet of the present invention has an area ratio of bainite of 95% or more, preferably a bainite single phase. This is because it is possible to achieve both strength and hole expansibility by using bainite as the metal structure. Further, since this structure is generated by transformation at a relatively high temperature, it is not necessary to cool to a low temperature during production, and this structure is preferable from the viewpoint of material stability and productivity.

  As the balance, 5% or less of pro-eutectoid ferrite, pearlite, martensite, and retained austenite are allowed. Proeutectoid ferrite is not a problem as long as it is sufficiently precipitation strengthened, but depending on the component, it may become soft, and when the area ratio exceeds 5%, the hole expandability slightly decreases due to the hardness difference from bainite. . In addition, when the area ratio of pearlite exceeds 5%, strength and workability may be impaired. When the area ratio of martensite and retained austenite that becomes martensite by processing-induced transformation is 1% or more and more than 5%, the interface between bainite and a structure harder than bainite becomes the starting point of cracking. Hole expandability deteriorates.

  Therefore, if the area ratio of bainite is 95% or more, the area ratio of the remaining pro-eutectoid ferrite, pearlite, martensite, and residual γ is 5% or less, so that the balance between strength and hole expansibility becomes good. However, martensite needs to be less than 1% as described above.

Here, the bainite in the present invention is the Japan Iron and Steel Institute Basic Research Group, Bainite Research Section / Edition; Recent Research on the Bainite Structure and Transformation Behavior of Low Carbon Steels-Final Report of Bainite Research Group (1994) The continuous cooling transformation structure (Zw) in the intermediate stage between the microstructure containing polygonal ferrite and pearlite produced by the diffusion mechanism and the martensite produced by the non-diffusion shearing mechanism as described in A defined microstructure.
That is, the continuous cooling transformation structure (Zw) is mainly composed of Bainitic ferrite (α ° _B ), Granular bainitic ferrite (α _B ), Quasi, as described in the above-mentioned references 125 to 127 as an optical microscope observation structure. It is composed of -polygonal ferrite (α _q ), and is further defined as a microstructure containing a small amount of retained austenite (γ _r ) and Martensite-austenite (MA).

Note that α _q is not distinguished from PF because the internal structure does not appear by etching as in polygonal ferrite (PF), but the shape is ashular. Here, α _q is a grain whose ratio (lq / dq) satisfies lq / dq ≧ 3.5 when the perimeter length lq of the target crystal grain and its equivalent circle diameter is dq.
The continuous cooling transformation structure (Zw) in the present invention is defined as a microstructure containing one or more of α ° _B , α _B , α _q , γ _r and MA. Note that a small amount of γ _r and MA is 3% or less in total.

This continuously cooled transformation structure (Zw) may be difficult to distinguish by optical microscope observation during etching using a Nital reagent. In that case, it is determined by using the ^{EBSP-OIM TM.}
The EBSP-OIM (Electron Back Scatter Diffraction Pattern-Orientation Image Microscopy registered trademark) method is formed by irradiating an electron beam onto a highly inclined sample in a scanning electron microscope (SEM) and backscattering. It consists of a device and software that measures the crystal orientation of the irradiation point in a short time by photographing the Kikuchi pattern with a high-sensitivity camera and processing the computer image.
The EBSP method can quantitatively analyze the microstructure and crystal orientation of the bulk sample surface, and the analysis area can be analyzed up to a minimum resolution of 20 nm as long as it is within the region that can be observed with the SEM, depending on the resolution of the SEM. The analysis by the EBSP-OIM method is performed by mapping several tens of thousands of regions to be analyzed in a grid at equal intervals over several hours. For polycrystalline materials, the crystal orientation distribution and crystal grain size in the sample can be seen. In the present invention, an image that can be discriminated from an image mapped with the azimuth difference of each packet as 15 ° may be conveniently defined as a continuous cooling transformation structure (Zw).

Moreover, the structure fraction of pro-eutectoid ferrite was calculated | required with the KAM (Kernel Average Misorientation) method equipped with EBSP-OIM.
In the KAM method, between six adjacent pixels (first approximation) or further outside 12 (second approximation) and further outside 18 (third approximation) pixels of a regular hexagonal pixel in the measurement data. Is calculated for each pixel, and the difference is averaged and the value is the value of the center pixel.
By performing this calculation so as not to cross the grain boundary, a map expressing the orientation change in the grain can be created. That is, this map represents a strain distribution based on local orientation changes in the grains. In the present invention, the analysis condition is such that the azimuth difference between adjacent pixels in EBSP-OIM is calculated as a third approximation, and the azimuth difference is 5 ° or less. It is.

In the present invention, the pro-eutectoid ferrite is defined as the microstructure up to the surface area fraction of the pixel calculated as the above-mentioned third difference of the first approximation 1 ° or less.
This is because the polygonal pro-eutectoid ferrite transformed at high temperature is formed by diffusion transformation, so the dislocation density is small and the intra-granular distortion is small, so the intra-granular difference in crystal orientation is small. From the various investigation results, the polygonal ferrite volume fraction obtained by optical microscope observation and the area fraction of the area obtained by the third approximation of the first difference of 1 ° measured by the KAM method were obtained. .

Next, the manufacturing method of the cold rolled steel sheet of this invention is described.
In order to achieve excellent local deformability, it is important to form a texture with an X-ray random intensity ratio and to make steel sheets that satisfy the conditions of grain refinement, equiaxed graining, and homogenization. The details of the manufacturing conditions for simultaneously satisfying these requirements are described below.

The production method preceding hot rolling is not particularly limited. That is, various secondary smelting may be performed following the smelting by a blast furnace or an electric furnace, and then the casting may be performed by a method such as a thin slab casting in addition to a normal continuous casting and an ingot method. In the case of continuous casting, after cooling to low temperature once, it may be heated again and then hot rolled, or the cast slab may be continuously hot rolled. Scrap may be used as a raw material.
Moreover, in hot rolling, a sheet bar may be joined after rough rolling, and finish rolling may be performed continuously. At this time, the coarse bar may be wound once in a coil shape, stored in a cover having a heat retaining function as necessary, and rewound again to perform bonding.

The high-strength steel sheet excellent in local deformability of the present invention is obtained when the following requirements are satisfied.
First, the austenite grain size after rough rolling, that is, before finish rolling is important, and it is desirable that the austenite grain size before finish rolling is small, and if it is 200 μm or less, it greatly contributes to refinement of grain units and homogenization of the main phase. It has been found.

In order to obtain the austenite grain size before finish rolling of 200 μm or less, as shown in FIG. 3, rough rolling in a temperature range of 1000 ° C. or more and 1200 ° C. or less is performed at least once at a reduction ratio of 20% or more. It has also been found that a predetermined austenite particle size can be obtained. However, in order to further improve the homogeneity and the local ductility, it is necessary to perform rolling at least once at a rolling rate of at least 40% in a rough rolling rate in a temperature range of 1000 ° C. or higher and 1200 ° C. or lower.
Finer grains can be obtained as the rolling reduction ratio and the number of rolling reductions are increased, and in order to obtain this effect more efficiently, it is desirable to obtain an austenite grain size of 100 μm or less, and for this purpose, 40% It is desirable to perform the above rolling twice or more. However, the reduction exceeding 70% or the rough rolling exceeding 10 times may cause a decrease in temperature or excessive generation of scale.

In this way, reducing the austenite grain size before finish rolling improves the local deformability through the control of austenite recrystallization in the subsequent finish rolling, finer grain size of the final structure, and equiaxing. It is effective for.
This is presumed to be due to the function of the austenite grain boundary after rough rolling (that is, before finish rolling) as one of the recrystallization nuclei during finish rolling.

  In order to confirm the austenite grain size after rough rolling, it is desirable to cool the plate piece before finishing rolling as much as possible, and the plate piece is cooled at a cooling rate of 10 ° C./s or more. The structure of the cross section is etched to raise the austenite grain boundary and measured with an optical microscope. At this time, 20 fields of view or more are measured by image analysis or a point count method at a magnification of 50 times or more.

Further, the average value of the X-ray random intensity ratios of {100} <011> to {223} <110> orientation groups on the plate surface at a thickness of 5/8 to 3/8 from the surface of the steel plate, and {332} <113 > In order to make the X-ray random intensity ratio of the crystal orientation within the above range, the T1 temperature determined by the steel plate components in the finish rolling after the rough rolling T1 (° C.) = 850 + 10 × (C + N) × Mn + 350 × Nb + 250 × Ti + 40 × B + 10 × Cr + 100 × Mo
+ 100 × V (Formula 1)
Based on the above, it is necessary to make the sum of the rolling reductions in the temperature range of T1 + 30 ° C. or more and T1 + 200 ° C. or less 50% or more, and subsequently suppress the rolling reduction at T1 or more and less than T1 + 30 ° C. as much as possible. Thus, the local deformability of the final product can be ensured. FIGS. 4 to 5 show the relationship between the rolling reduction in each temperature region and the X-ray random intensity ratio in each direction.

  That is, as shown in FIGS. 4 and 5, large pressure in a temperature range of T1 + 30 ° C. or more and T1 + 200 ° C. or less, and subsequent light pressure at T1 or more and less than T1 + 30 ° C. are shown in Table 2 of Examples below. 3, the average value of the X-ray random intensity ratios of {100} <011> to {223} <110> orientation groups of the plate surface at a thickness of 5/8 to 3/8 from the surface of the steel plate, By controlling the X-ray random intensity ratio of the {332} <113> crystal orientation, the local deformability of the final product is dramatically improved.

This T1 temperature itself is obtained empirically. Based on the T1 temperature, the inventors have empirically found that recrystallization in the austenitic region of each steel is promoted. In order to obtain a better local deformability, it is important to accumulate strain due to large reduction, and the total reduction ratio is 50% or more. Furthermore, it is desirable to take a reduction of 70% or more. On the other hand, taking a reduction ratio of more than 90% adds to securing temperature and adding excessive rolling.
Furthermore, in order to increase the homogeneity of the hot-rolled sheet and increase the local ductility to the maximum, the rolling rate in the final pass needs to be 25% or more in rolling in the temperature range of T1 + 30 ° C. or higher and T1 + 200 ° C. or lower. There is. However, when higher workability is required, the final two passes must be 25% or more.

  Further, in order to promote uniform recrystallization by releasing the accumulated strain, after a large pressure at T1 + 30 ° C. or higher and T1 + 200 ° C. or lower, the processing amount in the temperature range of T1 ° C. or higher and lower than T1 + 30 ° C. is minimized. It is necessary to set the rolling reduction ratio at T1 ° C or higher and lower than T1 + 30 ° C to 30% or lower. From the plate shape, a rolling reduction rate of 10% or higher is desirable. Is preferably 0%. In addition, if the rolling reduction at T1 ° C. or more and less than T1 + 30 ° C. is large, recrystallized austenite grains expand, and if the retention time is short, recrystallization does not proceed sufficiently and local deformability deteriorates. End up. That is, in the manufacturing conditions of the present invention, the austenite is uniformly and finely recrystallized in finish rolling to control the texture of the product and improve the local deformability such as hole expandability and bendability. .

When rolling is performed at a temperature lower than the temperature range specified above or a large reduction ratio is taken, the austenite texture develops, and finally the steel sheet surface obtained [ Each crystal orientation of the predetermined X-ray intensity level described in (1) of the means] cannot be obtained.
On the other hand, when rolling is performed at a temperature higher than the above-mentioned temperature range or a small reduction ratio is taken, coarse grains and mixed grains are formed, and the area ratio of crystal grains exceeding 20 μm increases.

  Whether the rolling specified above is performed or not can be determined by the actual result or calculation from the rolling load, sheet thickness measurement, etc., and the temperature can be measured if there is an inter-stand thermometer, or the line It can be obtained by a calculation simulation considering processing heat generation from the speed and rolling reduction, or both.

The hot rolling performed as described above ends at a temperature of Ar ₃ or higher. When hot rolling is finished at Ar ₃ or less, it becomes two-phase rolling into austenite and ferrite, and accumulation in {100} <011> to {223} <110> orientation groups becomes strong, and as a result, local deformability is increased. Deteriorates significantly.

  Furthermore, in order to refine the grain unit and suppress the stretched grain, it is necessary to suppress the maximum heat generation amount during the reduction at T1 + 30 ° C. to T1 + 200 ° C., that is, the temperature increase (° C.) due to the reduction to 18 ° C. or less. It is desirable to use cooling between stands.

Cooling after reduction in the final rolling stand of rolling in a temperature range of T1 + 30 ° C. or more and less than T1 + 200 ° C. has a large effect on the austenite grain size, and this is the equiaxed grain fraction and coarse grain of the structure after cold rolling annealing. It has a strong influence on the fraction.
It is desirable that the time t (seconds) from the final pass under the reduction in the temperature range of T1 + 30 ° C. or higher and lower than T1 + 200 ° C. to 5 seconds or less. If it is longer than 5 seconds, the austenite grains become coarse and the strength and local ductility are lowered. Furthermore, the time until the start of cooling needs to satisfy (Equation 2) with respect to the final reduction temperature Tf and the rolling rate P1. When the time t becomes t1 or more, coarsening progresses and local ductility is remarkably reduced.
t <t1 (Formula 2)
Here, t1 is a numerical value that can be obtained by the following (formula 3).
t1 = 0.001 ((Tf−T1) × P1) ² −0.109 ((Tf−T1) × P1) +3.1 (Equation 3)

  In order to suppress grain growth after recrystallization as much as possible, the temperature change of rapid cooling immediately after starting within t seconds defined by (Equation 2) is set to 40 ° C. or more and 150 ° C. or less. Immediately after the cooling immediately after the quenching, no particular regulation is provided, and the effects of the present invention can be obtained even if a cooling pattern for controlling the structure for each purpose is taken.

  The hot-rolled original sheet manufactured as described above is cold-rolled at 30% or more and 70% or less. When the rolling reduction is 30% or less, it is difficult to cause recrystallization in the subsequent annealing step, the equiaxed grain fraction is lowered, and the grains after annealing are coarsened. When rolling exceeds 70%, the anisotropy becomes strong because the texture is developed during annealing. For this reason, it is 70% or less.

The cold-rolled steel sheet is then held in the temperature range of Ae 3 to 950 ° C. for 1 to 300 seconds in order to obtain austenite single phase steel or almost austenite single phase steel. If the temperature is lower or shorter than this, the fraction of the bainite structure is not 95% or more in the subsequent cooling step, and the increase in local ductility due to texture control is reduced. On the other hand, when the temperature exceeds 950 ° C. or the holding time exceeds 300 seconds, the crystal grains become coarse, and the area ratio of grains of 20 μm or less increases. Ae ₃ [° C.] is calculated by the following (formula 5) according to the content [mass%] of C, Mn, Si, Cu, Ni, Cr, and Mo. In addition, when not containing a selective element, it calculates as 0.
Ae ₃ = 911-239C-36Mn + 40Si-28Cu-20Ni-12Cr + 63Mo (Formula 5)

Thereafter, primary cooling is performed to a temperature of 500 ° C. or lower so that the average cooling rate in the temperature range between Ae3 and 500 ° C. is 10 ° C./s or higher and 200 ° C./s or lower.
When the primary cooling rate is less than 10 ° C./s, ferrite is excessively generated, and the fraction of the bainite structure cannot be increased to 95% or more. Therefore, the increase in local ductility due to texture control is reduced. On the other hand, even if the cooling rate exceeds 200 ° C./s, the controllability of the cooling end point temperature is remarkably deteriorated, so that it is 200 ° C./s or less. Preferably, in order to reliably suppress ferrite transformation and pearlite transformation, the average cooling rate at HF to 0.5HF + 250 ° C. does not exceed the average cooling rate at 0.5HF + 250 ° C. to 500 ° C.

In order to promote bainite transformation, over-aging heat treatment is performed in the temperature range of 350 ° C. to 500 ° C. following primary cooling. The time held in this temperature range is t2 seconds or longer that satisfies the following (Formula 4) according to the overaging treatment temperature T2. However, considering the applicable temperature range of (Equation 4), the maximum value of t2 is 400 seconds.
Log (t2) = 0.0002 (T2-425) ² +1.18 (Formula 4)
In the present invention, holding means not only isothermal holding but also retention in a temperature range of 500 to 350 ° C. That is, after cooling to 350 ° C., it may be heated to 500 ° C., or may be cooled to 500 ° C. and then cooled to 3500 ° C.

  The steel sheet according to the present invention can be applied to stretch forming and composite forming mainly composed of bending, such as bending, stretching, and drawing.

The technical contents of the present invention will be described with reference to examples of the present invention.
As examples, the results of studies using steels having the composition shown in Table 1 and satisfying the constituents of claims of the present invention from A to T and comparative steels from a to i will be described.
After casting, these steels are reheated as they are or once cooled to room temperature, heated to a temperature range of 1000 ° C. to 1300 ° C., and then hot-rolled under the conditions shown in Table 2 to obtain an Ar 3 transformation temperature. The hot rolling is finished as described above, and after cooling under the conditions shown in Table 2 to obtain a hot-rolled steel sheet having a thickness of 2 to 5 mm, pickling and cold-rolling, then cold-rolling to a thickness of 1.2 to 2.3 mm After rolling and annealing under the annealing conditions shown in Table 2, 0.5% skin pass rolling was performed for material evaluation.

Table 1 shows the chemical composition of each steel, and Table 2 shows the production conditions. Table 3 shows the structure and mechanical properties of each.
As an index of local deformability, the hole expansion rate and the critical bending radius by 60 ° V-bending were used. The bending test was C direction bending. The tensile test and the bending test were compliant with JIS Z 2241 and Z 2248 (V block 90 ° bending test), and the hole expansion test was compliant with the iron standard JFS T1001. The X-ray random intensity ratio was measured at 0.5 μm pitch in the 3/8 to 5 / region of the cross section parallel to the rolling direction using the above-mentioned EBSP. In Table 3, B means bainite, P means pearlite, F means proeutectoid ferrite, M means martensite, and rA means retained austenite.

  As shown in FIGS. 6 and 7, it is understood that only those satisfying the definition of the present invention can have both excellent hole expansibility and bendability.

Claims (10)

% By mass
C: 0.02% or more, 0.20% or less,
Si: 0.001% or more, 2.5% or less,
Mn: 0.01% or more and 4.0% or less,
P: 0.001% or more, 0.15% or less,
S: 0.0005% or more, 0.03% or less,
Al: 0.001% or more, 2.0% or less,
N: 0.0005% or more, 0.01% or less,
O: 0.0005% or more, 0.01% or less,
Si + Al: containing less than 1.0%, consisting of the balance iron and inevitable impurities,
The area ratio of bainite in the metal structure of the steel sheet is 95% or more,
In the texture of the steel sheet, the average value of the X-ray random intensity ratios of {100} <011> to {223} <110> orientation groups of the plate surface at a thickness of 5/8 to 3/8 at least from the surface of the steel plate is Less than 4.0 and the {332} <113> crystal orientation X-ray random intensity ratio is 5.0 or less,
Further, a high-strength cold-rolled steel sheet having excellent local deformability, wherein the volume average particle size in the metal structure of the steel sheet is 7 μm or less.   Further, the ratio of the length dL in the rolling direction to the length dt in the sheet thickness direction, the proportion of grains having dL / dt of 3.0 or less in the grain unit of bainite is 50% or more. Item 5. A high-strength cold-rolled steel sheet excellent in local deformability according to item 1. Furthermore, in mass%,
Ti: 0.001% or more, 0.20% or less,
Nb: 0.001% or more, 0.20% or less,
V: 0.001% or more, 1.0% or less,
W: A high-strength cold-rolled steel sheet excellent in local deformability according to claim 1 or 2, containing one or more of 0.001% or more and 1.0% or less. Furthermore, in mass%,
B: 0.0001 or more and 0.0050% or less Mo: 0.001 or more and 1.0% or less,
Cr: 0.001 or more and 2.0% or less,
Cu: 0.001 or more, 2.0% or less,
Ni: 0.001 or more and 2.0% or less,
Co: 0.0001 or more and 1.0% or less,
Sn: 0.0001 or more and 0.2% or less,
Zr: 0.0001 or more and 0.2% or less,
The high-strength cold-rolled steel sheet excellent in local deformability according to any one of claims 1 to 3, containing As: 0.0001 or more and 0.50% or less. Furthermore, in mass%,
Mg: 0.0001 or more and 0.010% or less,
REM: 0.0001 or more and 0.1% or less,
The high-strength cold-rolled steel sheet excellent in local deformability according to any one of claims 1 to 4, containing one or more of Ca: 0.0001 or more and 0.010% or less.   A high-strength cold steel having excellent local deformability, comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the high-strength cold-rolled steel sheet according to any one of claims 1 to 5. Rolled steel sheet. In producing the high-strength cold-rolled steel sheet according to any one of claims 1 to 5,
After melting into a predetermined steel plate component, it is cast into a steel ingot or slab and subjected to rough rolling at a temperature range of 1000 ° C. or higher and 1200 ° C. or lower and at least once at least 20%, and austenite grain size Is 200 μm or less,
Then, if the temperature determined by the steel plate component in (Formula 1) in finish rolling is T1 ° C., the total rolling reduction in the temperature range of T1 + 30 ° C. or more and T1 + 200 ° C. or less is 50% or more, and then T1 ° C. or more. The total rolling reduction in the temperature range of less than T1 + 30 ° C. is set to 30% or less, and then the hot rolling is finished at the Ar3 transformation temperature or higher, and in the temperature range of T1 ° C. or higher and T1 + 30 ° C. or lower after the final reduction. After the stop time t (seconds) satisfies (Equation 2), cooling is performed,
Subsequently, pickling, rolling 30% or more and 70% or less in the cold, and then annealing in the temperature range of Ae3 to 950 ° C for 1 to 300 seconds, then the temperature range of Ae3 to 500 ° C The average cooling rate is 10 ° C./s or more and 200 ° C./s or less, and is further maintained at t2 seconds or more and 400 seconds or less satisfying (Equation 4) at 350 ° C. or more and 500 ° C. or less overaging heat treatment temperature. A method for producing a high-strength cold-rolled steel sheet having excellent local deformability. However, the maximum value of t2 is 400 seconds.
T1 (℃) = 850 + 10 × (C + N) × Mn + 350 × Nb + 250 × Ti + 40 × B + 10 × Cr + 100 × Mo + 100 × V
... (Formula 1)
t <t1 (Formula 2)
Here, t1 is expressed by (Expression 3).
t1 = 0.001 ((Tf−T1) × P1) ² −0.109 ((Tf−T1) × P1) +3.1 (Equation 3)
Here, Tf and P1 are a temperature and a reduction ratio at the time of final reduction in a temperature range of T1 + 30 ° C. or higher and T1 + 200 ° C. or lower, respectively.
log (t2) = 0.0002 (T2-425) ² + 1.18 (Formula 4)
Here, T2 is the overaging temperature, and the maximum value of t2 is 400.   The high strength cold-rolled steel sheet having excellent local deformability, wherein the rolling rate in the final pass of rolling in the temperature range of T1 + 30 ° C or higher and T1 + 200 ° C or lower is 25% or more in the manufacturing method according to claim 7. Manufacturing method.   The manufacturing method according to claim 7 or 8, wherein after the hot rolling is finished at an Ar3 transformation temperature or higher, the cooling temperature change is 40 ° C or higher and 150 ° C or lower within t seconds shown by the (Formula 2). A method for producing a high-strength cold-rolled steel sheet having excellent local deformability, characterized in that rapid cooling is started immediately after the speed reaches 50 ° C./s or more.   A local deformability characterized by forming a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the cold-rolled steel plate produced by the production method according to any one of claims 7 to 9. An excellent method for producing high-strength cold-rolled steel sheets.

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