High-strength steel sheet and high-strength zinc-coated steel sheet which have excellent ductility and stretch-flangeability and manufacturing method thereof

TECHNICAL FIELD

The present invention relates to a high-strength steel sheet and a high-strength zinc-coated steel-sheet which have excellent ductility and stretch-flangeability and a manufacturing method thereof.

Priority is claimed on Japanese Patent Application Nos. 2010-208329 and 2010-208330, filed Sep. 16, 2010, the content of which is incorporated herein by reference.

BACKGROUND ART

In recent years, there has been an increasing demand for a high-strength steel sheet used in a vehicle or the like, and a high-strength cold-rolled steel sheet with a maximum tensile stress of 900 MPa or more is also being used.

Generally, as the strength of a steel sheet is enhanced, ductility and stretch-flangeability are lowered, and workability is degraded. However, a high-strength steel sheet with sufficient workability has been demanded in recent years.

As a conventional technique for enhancing ductility and stretch-flangeability of a high-strength steel sheet, a high-tensile galvanized steel sheet, which has a composition containing by mass percentage, C: 0.05 to 0.20%, Si: 0.3 to 1.8%, Mn: 1.0 to 3.0%, S: 0.005% or less, the remainder composed of Fe and inevitable impurities, has a composite structure including ferrite, tempered martensite, retained austenite, and low temperature transformation phase, and contains by volume percentage 30% or more of ferrite, 20% or more of tempered martensite, 2% or more of retained austenite, in which average crystal grain sizes of ferrite and tempered martensite are 10 μm or less, is an exemplary example (see Patent Document 1, for example).

In addition, as a conventional technique for enhancing workability of a high-strength steel sheet, a high-tensile cold-rolled steel sheet, in which amounts of C, Si, Mn, P, S, Al, and N are adjusted, which further contains 3% or more of ferrite and a total of 40% or more of bainite containing carbide and martensite containing carbide as metal structures of the steel sheet containing one or more of Ti, Nb, V, B, Cr, Mo, Cu, Ni, and Ca as necessary, in which the total amount of ferrite, bainite, and martensite is 60% or more, and which further has a structure in which the number of ferrite grains containing cementite, martensite, or retained austenite therein corresponds to 30% or more of the total number of ferrite grains and has tensile strength of 780 MPa or more, is an exemplary example (see Patent Document 2, for example).

Moreover, as a conventional technique for enhancing stretch-flangeability of a high-strength steel sheet, a steel sheet in which a difference in hardness between a hard part and a soft part of the steel sheet is reduced is an exemplary example. For example, Patent Document 3 discloses a technique in which the standard deviation of hardness in the steel sheet is reduced and uniform hardness is given to the entire steel sheet. Patent Document 4 discloses a technique in which hardness in the hard part is lowered by heat treatment and the difference in hardness from that in the soft part is reduced. Patent Document 5 discloses a technique in which the difference in hardness from the soft part is reduced by configuring the hard part of relatively soft bainite.

Furthermore, as a conventional technique for enhancing stretch-flangeability of a high-strength steel sheet, a steel sheet, which has a structure containing by an area ratio 40 to 70% of tempered martensite and a remainder composed of ferrite, in which a ratio between an upper limit value and a lower limit value of Mn concentration in a cross-section in a thickness direction of the steel sheet is reduced (see Patent Document 6, for example) may be exemplified.

CITATION LIST
Patent Documents

[Patent Document 1] Japanese Unexamined Patent Application, First Publication No. 2001-192768

[Patent Document 2] Japanese Unexamined Patent Application, First Publication No. 2004-68050

[Patent Document 3] Japanese Unexamined Patent Application, First Publication No. 2008-266778

[Patent Document 4] Japanese Unexamined Patent Application, First Publication No. 2007-302918

[Patent Document 5] Japanese Unexamined Patent Application, First Publication No. 2004-263270

[Patent Document 6] Japanese Unexamined Patent Application, First Publication No. 2010-65307

SUMMARY OF INVENTION
Technical Problem

According to the conventional techniques, however, workability of the high-strength steel sheet with a maximum tensile strength of 900 MPa or more is insufficient, and it has been desired to further enhance ductility and stretch-flangeability and to thereby further enhance workability.

The present invention is made in view of such circumstances, and an object thereof is to provide a high-strength steel sheet, which has excellent ductility and stretch-flangeability and has excellent workability while high strength is secured such that the maximum tensile strength becomes 900 MPa or more, and a manufacturing method thereof.

Solution to Problem

The present inventor conducted intensive study in order to solve the above problems. As a result, the present inventor found that it is possible to secure a maximum tensile strength as high as 900 MPa or more and significantly enhance ductility and stretch-flangeability (hole expanding property) by allowing the steel sheet to have a large hardness difference by increasing a micro Mn distribution inside the steel sheet and have a sufficiently small average crystal grain size by controlling dispertion in the hardness distribution.

[1] A high-strength steel sheet which has excellent ductility and stretch-flangeability, including by mass percentage: 0.05 to 0.4% of C; 0.1 to 2.5% of Si; 1.0 to 3.5% of Mn; 0.001 to 0.03% of P; 0.0001 to 0.01% of S; 0.001 to 2.5% of Al; 0.0001 to 0.01% of N; 0.0001 to 0.008% of O; and a remainder composed of iron and inevitable impurities, wherein a steel sheet structure contains by volume fraction 10 to 50% of a ferrite phase, 10 to 50% of a tempered martensite phase, and a remaining hard phase, wherein when a plurality of measurement regions with diameters of 1 μm or less are set in a range from ⅛ to ⅜ of thickness of the steel sheet, hardness measurement values in the plurality of measurement regions are arranged in an ascending order to obtain a hardness distribution, an integer N0.02, which is a number obtained by multiplying a total number of the hardness measurement values by 0.02 and, if present, by rounding up a decimal number, is obtained, a hardness of a measurement value which is an N0.02-th largest value from a smallest hardness measurement value is regarded as a 2% hardness, an integer N0.98 which is a number obtained by multiplying the total number of the hardness measurement values by 0.98 and, if present, by rounding down the decimal number is obtained, and a hardness of a measurement value which is an N0.98-th largest value from the smallest hardness measurement value is regarded as a 98% hardness, the 98% hardness is 1.5 or more times as high as the 2% hardness, wherein a kurtosis K* of the hardness distribution between the 2% hardness and the 98% hardness is equal to or more than −1.2 and equal to or less than −0.4, and wherein an average crystal grain size in the steel sheet structure is 10 μm or less.

[2] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to [1], wherein a difference between a maximum value and a minimum value of Mn concentration in a base iron in a thickness range from ⅛ to ⅜ of the steel sheet is equal to or more than 0.4% and equal to or less than 3.5% when converted into the mass percentage.

[3] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to [1] or [2], wherein when a section from the 2% hardness to the 98% hardness is equally divided into 10 parts, and 10 1/10-sections are set, a number of the hardness measurement values in each 1/10-section is 2 to 30% of a number of all measurement values.

[4] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [3], wherein the hard phase includes any one of or both a bainitic ferrite phase and a bainite phase of 10 to 45% by a volume fraction, and a fresh martensite phase of at 10% or less.

[5] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [4], wherein the steel sheet structure further includes 2 to 25% of a retained austenite phase.

[6] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [5], further including by mass percentage one or more of 0.005 to 0.09% of Ti; and 0.005 to 0.09% of Nb.

[7] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [6], further including by mass percentage one or more of: 0.0001 to 0.01% of B; 0.01 to 2.0% of Cr; 0.01 to 2.0% of Ni; 0.01 to 2.0% of Cu; and 0.01 to 0.8% of Mo.

[8] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [7], further including by mass percentage: 0.005 to 0.09% of V.

[9] The high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [1] to [8], further including one or more of Ca, Ce, Mg, and REM at 0.0001 to 0.5% by mass percentage in total.

[10] A high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability, wherein the high-strength zinc-coated steel sheet is produced by forming a zinc-coated layer on a surface of the high-strength steel sheet according to any one of [1] to [9].

[11] A manufacturing method of a high-strength steel sheet which has an excellent ductility and a stretch-flangeability, the method including: a hot rolling process in which a slab containing the chemical constituents according to any one of [1] or [6] to [9] is heated up to 1050° C. or higher directly or after cooling once, a hot rolling is performed thereon at a higher temperature of one of 800° C. and an Ar₃transformation point, and a winding is performed in a temperature range of 750° C. or lower such that an austenite phase in a structure of a rolled material after rolling occupies 50% by volume or more; a cooling process in which the steel sheet after the hot rolling is cooled from a winding temperature to (the winding temperature−100)° C. at a rate of 20° C./hour or lower while a following Equation (1) is satisfied; and a process in which continuous annealing is performed on the steel sheet after the cooling, wherein in the process in which continuous annealing is performed, the steel sheet is annealed at a maximum heating temperature of 750 to 1000° C., a first cooling in which the steel sheet is cooled from the maximum heating temperature to a ferrite transformation temperature range or lower and maintained in the ferrite transformation temperature range for 20 to 1000 seconds is subsequently performed, a second cooling in which the steel sheet is cooled at a cooling rate of 10° C./second or higher on average in a bainite transformation temperature range and cooling is stopped within a range from a martensite transformation start temperature−120° C. to the martensite transformation start temperature is subsequently performed, the steel sheet after the second cooling is maintained in a range from a second cooling stop temperature to the martensite transformation start temperature for 2 to 1000 seconds, the steel sheet is subsequently reheated up to a reheating stop temperature, which is equal to or more than a bainite transformation start temperature−100° C., at a rate of temperature increase of 10° C./second or higher on average in the bainite transformation temperature range, and a third cooling in which the steel sheet after the reheating is cooled from the reheating stop temperature to a temperature which is lower than the bainite transformation temperature range and maintained in the bainite transformation temperature range for 30 seconds or more is performed:

$\begin{matrix} [Equation 1] \\ {[\int_{T_{c - 100}}^{T_{c}} 9.47 \times 10^{5} \cdot \exp (- \frac{18480}{T + 273}) \cdot t (T) \cdot ⅆ T]}^{0.5} \geq 1.0 & (1) \end{matrix}$

[where t(T) in Equation (1) represents maintaining time (seconds) of the steel sheet at a temperature T° C. in the cooling process after the winding.]

[12] The manufacturing method of the high-strength steel sheet which has excellent ductility and stretch-flangeability according to [11], wherein the winding temperature after the hot rolling is equal to or more than a Bs point and equal to or less than 750° C.

[13] The manufacturing method of the high-strength steel sheet which has excellent ductility and stretch-flangeability according to [11] or [12], further including between the cooling process and the continuous annealing process: a cold rolling process in which the steel sheet is subjected to acid pickling and a cold rolling at rolling reduction from 35 to 80%.

[14] The manufacturing method of the high-strength steel sheet which has excellent ductility and stretch-flangeability according to any one of [11] to [13], wherein a sum of a time during which the steel sheet is maintained in the bainite transformation temperature range in the second cooling and a time during which the steel sheet is maintained in the bainite transformation temperature range in the reheating is 25 seconds or less.

[15] A manufacturing method of a high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability, wherein the steel sheet is dipped into a zinc plating bath in the reheating in manufacturing the high-strength steel sheet based on the manufacturing method according to any one of [11] to [14].

[16] A manufacturing method of a high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability, wherein the steel sheet is dipped into a zinc plating bath in the bainite transformation temperature range in the third cooling in manufacturing the high-strength steel sheet based on the manufacturing method according to any one of [11] to [14].

[17] A manufacturing method of a high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability, wherein a zinc electroplating is performed after manufacturing the high-strength steel sheet based on the manufacturing method according to any one of [11] to [14].

[18] A manufacturing method of a high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability, wherein a hot-dip zinc-plating is performed after manufacturing the high-strength steel sheet based on the manufacturing method according to any one of [11] to [14].

Advantageous Effects of Invention

The high-strength steel sheet of the present invention contains predetermined chemical constituents, when a plurality of measurement regions with diameters of 1 μm or less are set in a range from ⅛ to ⅜ of a thickness of the steel sheet, hardness measurement values in the plurality of measurement regions are arranged in ascending order to obtain a hardness distribution, an integer N0.02 which is a number obtained by multiplying a total number of the hardness measurement values by 0.02 and, if present, by rounding up a decimal number, is obtained, a hardness of a measurement value which is an N0.02-th largest value from the smallest hardness measurement value is regarded as a 2% hardness, an integer N0.98 which is a number obtained by multiplying the total number of the hardness measurement values by 0.98 and, if present, rounding down a decimal number, is obtained, and a hardness of a measurement value which is an N0.98-th largest value from the smallest hardness measurement value is regarded as a 98% hardness, the 98% hardness is 1.5 or more times as high as the 2% hardness, a kurtosis K* of the hardness distribution between the 2% hardness and the 98% hardness is equal to or less than −0.40, an average crystal grain size in the steel sheet structure is 10 μm or less, and therefore, the steel sheet which has excellent ductility and stretch-flangeability is obtained while tensile strength which is as high as 900 MPa or more is secured.

In addition, a micro Mn distribution inside the steel sheet increases by winding the steel sheet after the hot rolling around a coil at 750° C. and cooling the steel sheet from the winding temperature to (the winding temperature−100)° C. at a cooling rate of 20° C./hour or lower while the above Equation (1) is satisfied, in the process for producing a hot-rolled coil from the slab containing the predetermined chemical constituents in the manufacturing method of the high-strength steel sheet according to the present invention.

In addition, since the process in which continuous annealing is performed on the steel sheet with increased Mn distribution includes a heating process in which the steel sheet is annealed at a maximum heating temperature of 750 to 1000° C., a first cooling process in which the steel sheet is cooled from the maximum heating temperature to a ferrite transformation temperature range or lower and maintained in a ferrite transformation temperature range for 20 to 1000 seconds, a second cooling process in which the steel sheet after the first cooling process is cooled at a cooling rate of 10° C./second or higher on average in a bainite transformation temperature range and cooling is stopped within a range from a martensite transformation start temperature−120° C. to the martensite transformation start temperature, a maintaining process in which the steel sheet after the second cooling process is maintained in a range from a second cooling stop temperature to the Ms point or lower for 2 to 1000 seconds, a reheating process in which the steel sheet after the maintaining process is reheated up to a reheating stop temperature, which is equal to or more than a bainite transformation start temperature−80° C., at a rate of temperature increase of 10° C./second or higher on average in the bainite transformation temperature range, and a third cooling process in which the steel sheet after the reheating process is cooled from the reheating stop temperature to a temperature which is lower than the bainite transformation temperature range and maintained in the bainite transformation temperature range for 30 seconds or more, the steel sheet structure is controlled such that the hardness difference inside the steel sheet is large and the average crystal grain size is sufficiently small, and it is possible to obtain the high-strength cold-rolled steel sheet which has excellent ductility and stretch-flangeability (hole expanding property) and has excellent workability while securing a maximum tensile strength of 900 MPa or more.

Furthermore, it is possible to obtain the high-strength zinc-coated steel sheet which has excellent ductility and stretch-flangeability (hole expanding property) and has excellent workability while securing the maximum tensile strength as high as 900 MPa or more by adding the process for forming the zinc-pated layer.

BRIEF DESCRIPTION OF DRAWINGS

FIG. 1 is a graph showing a relationship between hardness classified into a plurality of levels and a number of measurement values in each level, which is obtained by converting each measurement value while a difference between a maximum hardness measurement value and a minimum hardness measurement value is regarded as 100%, in relation to an example of a high-strength steel sheet according to the present invention.

FIG. 2 is a diagram for comparing the hardness distribution in the high-strength steel sheet according to the present invention with a normal distribution.

FIG. 3 is a graph schematically showing a relationship between a transformation rate and elapsed time of transformation treatment when the difference between a maximum value and a minimum value of Mn concentration in base iron is relatively large.

FIG. 4 is a graph schematically showing a relationship between a transformation rate and elapsed time of transformation treatment when a difference between a maximum value and a minimum value of Mn concentration in base iron is relatively small.

FIG. 5 is a graph illustrating temperature history of a cold-rolled steel sheet when the sheet is made to pass through a continuous annealing line, which shows a relationship between the temperature of the cold-rolled steel sheet and time.

DESCRIPTION OF EMBODIMENTS

The high-strength steel sheet according to the present invention is a steel sheet, which includes predetermined chemical components, in which an average crystal grain size in the structure thereof is 10 μm or less, 98% hardness is 1.5 or more times as high as 2% hardness in a hardness distribution when a plurality of measurement regions with diameters of 1 μm or less is set in a thickness range from ⅛ to ⅜ thereof, and measurement values of hardness in the plurality of measurement regions are aligned in an order from a smallest measurement value, and kurtosis K* of the hardness distribution between the 2% hardness region and the 98% hardness region is −0.40 or less. An example of hardness distribution in the high-strength steel sheet according to the present invention is shown in FIG. 1.

(Definition of Hardness)

Hereinafter, definition of hardness will be described, and 2% hardness and 98% hardness will be described first. Measurement values of hardness are obtained in the plurality of measurement regions set in a thickness range from ⅛ to ⅜ of the steel sheet, and an integer N0.02, which is a number obtained by multiplying the total number of the measurement values of hardness by 0.02 and, if present, by rounding up a decimal number, is obtained. In addition, when a number obtained by multiplying the total number of the measurement values of hardness by 0.98 includes a decimal number, an integer N0.98 is obtained by rounding down the decimal number. Then, hardness of an N0.02-th largest measurement value from the minimum hardness measurement value in the plurality of measurement regions is regarded as the 2% hardness. In addition, a hardness of an N0.98-th largest measurement value from the minimum hardness measurement value in the plurality of measurement regions is regarded as the 98% hardness. In the high-strength steel sheet of the present invention, the 98% hardness is preferably 1.5 or more times as high as the 2% hardness, and the kurtosis K* of the hardness distribution between the 2% hardness and the 98% hardness is preferably −0.40 or less.

Each diameter of the measurement regions is limited to 1 μm or less in setting the plurality of measurement regions in order to exactly evaluate dispertion in hardness resulting from a steel sheet structure including a ferrite phase, a bainite phase, a martensite phase, and the like. Since the average crystal grain size in the steel sheet structure is 10 μm or less in the high-strength steel sheet of the present invention, it is necessary to obtain hardness measurement values in narrower measurement regions than the average crystal grain size in order to exactly evaluate the dispertion in hardness resulting from the steel sheet structure, and specifically, it is necessary to set regions with diameters of 1 μm or less as the measurement regions. When the hardness is measured using an ordinary Vickers tester, an indentation size is too large to exactly evaluate the dispertion in hardness resulting from the structure.

Accordingly, the “hardness measurement value” in the present invention represents a value evaluated based on the following method. That is, a measurement value obtained by measuring hardness under an indentation load of 1 g using a dynamic micro-hardness tester provided with a Berkovich type three-sided pyramid indenter based on an indentation depth measurement method is used for the high-strength steel sheet of the present invention. The hardness measurement position is set to a range from ⅛ to ⅜ around ¼ of a sheet thickness in the sheet thickness cross-section which is parallel to a rolling direction of the steel sheet. In addition, the total number of the hardness measurement values ranges from 100 to 10000, and is preferably equal to or more than 1000. The thus measured indentation size has a diameter of 1 μm or less on the assumption that the indentation shape is a circular shape. When the indentation shape is rectangular shape or a triangular shape other than the circular shape, the dimension of the indentation shape in the longitudinal direction may be 1 μm or less.

In addition, the “average crystal grain size” in the present invention represents the size measured by the following method. That is, a grain size measured based on an EBSD (Electron BackScattering Diffraction) method is preferably used for the high-strength steel sheet of the present invention. A grain size observation surface ranges from ⅛ to ⅜ around ¼ of the sheet thickness in the sheet thickness cross-section which is parallel to the rolling direction of the steel sheet. In addition, it is preferable to calculate the average crystal grain size by applying a intercept method to a grain boundary map for the observation surface obtained by regarding a boundary, at which a crystal orientation difference between adjacent measurement points in a bcc crystal orientation becomes 15° or more, as a grain boundary.

In order to obtain a steel sheet which has excellent ductility, it is important to utilize a structure such as ferrite, which has excellent ductility, as the steel sheet structure. However, the structure which has excellent ductility is soft. Accordingly, it is necessary to employ a steel sheet structure containing a soft structure and a hard structure such as martensite in order to obtain a steel sheet with high ductility while having sufficient strength.

In the steel sheet with the steel sheet structure including both the soft structure and the hard structure, strain caused by deformation is more easily accumulated in the soft part and is not easily distributed to the hard part when a hardness difference between the soft part and the hard part is larger, and therefore ductility is enhanced.

Since the 98% hardness is 1.5 or more times as high as the 2% hardness in the high-strength steel sheet of the present invention, the hardness difference between the soft part and the hard part is sufficiently large, and therefore, it is possible to obtain sufficiently high ductility. In order to obtain further higher ductility, the 98% hardness is preferably 3.0 or more times as high as the 2% hardness, more preferably more than 3.0 times, further more preferably 3.1 or more times, further more preferably 4.0 or more times, and still further more preferably 4.2 or more times. When the measurement value of the 98% hardness is less than 1.5 times of the measurement value of the 2% hardness, the hardness difference between the soft part and the hard part is not sufficiently large, and therefore, ductility is insufficient. Meanwhile, the measurement value of the 98% hardness is 4.2 or more times of the measurement value of the 2% hardness, the hardness difference between the soft part and the hard part is sufficiently large, and both ductility and a hole expanding property are further enhanced, which is preferable.

As described above, the hardness difference between the soft part and the hard part is preferably larger from the standpoint of ductility. However, if regions with the large hardness difference are in contact with each other, a strain gap caused by deformation of the steel sheet occurs at the border part, and a micro-crack is easily generated. Since the micro-crack may become a start point of cracking, stretch-flangeability is degraded. In order to suppress the degradation of stretch-flangeability resulted from the large hardness difference between the soft part and the hard part, it is effective to reduce number of borders at which the regions with the large hardness difference are in contact with each other and shorten the length of each border at which the regions with the large hardness difference are in contact with each other.

Since the average crystal grain size of the high-strength steel sheet of the present invention, which is measured by the EBSD method, is 10 μm or less, the border, at which the regions with the large hardness differences are in contact with each other, in the steel sheet is shortened, degradation of stretch-flangeability resulting from the large hardness difference between the soft part and the hard part is suppressed, and excellent stretch-flangeability can be obtained. In order to obtain further excellent stretch-flangeability, the average crystal grain size is preferably 8 μm or less, and more preferably 5 μm. If the average crystal grain size exceeds 10 μm, the effect of shortening the border, at which the regions with the large hardness difference are in contact with each other, in the steel sheet is not sufficient, and it is not possible to sufficiently suppress the degradation of stretch-flangeability.

In addition, in order to reduce the number of the borders at which the regions with the large hardness difference are in contact with each other, the steel sheet structure having a variety of narrow distribution of hardness, in which dispertion of the hardness distribution in the steel sheet is small, may be employed.

According to the high-strength steel sheet of the present invention, the dispertion in the hardness distribution in the steel sheet is reduced by setting the kurtosis K* of the hardness distribution to be −0.40 or less, it is possible to reduce the borders at which the regions with the large hardness difference are in contact with each other and thereby to obtain excellent stretch-flangeability. In order to obtain further excellent stretch-flangeability, the kurtosis K* is preferably −0.50 or less, and more preferably −0.55 or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the kurtosis K*, it is difficult to set K* to be less than −1.20, and therefore, this value is regarded as the lower limit.

In addition, the kurtosis K* is a value which can be obtained by the following Equation (2) based on the hardness distribution and is a numerical value obtained as a result of evaluation of the hardness distribution by comparing the hardness distribution with the normal distribution. A case in which the kurtosis is a negative value denotes that a hardness distribution curve is relatively flat, and a large absolute value denotes that the hardness distribution deviates further from the normal distribution.

$\begin{matrix} [Equation 2] \\ K^{*} = {\frac{(N_{0.98} - N_{0.02} + 1) (N_{0.98} - N_{0.02} + 2)}{(N_{0.98} - N_{0.02}) (N_{0.98} - N_{0.02} - 1) (N_{0.98} - N_{0.02} - 2)}} \cdot \sum_{i = N_{0.02}}^{N_{0.98}} {(\frac{H_{i} H^{*}}{s^{*}})}^{4} - \frac{3 {(N_{0.98} - N_{0.02})}^{2}}{(N_{0.98} - N_{0.02} - 1) (N_{0.98} - N_{0.02} - 2)} & (2) \end{matrix}$

Hi: hardness of an i-th largest measurement point from a measurement value of the minimum hardness

H*: average hardness from the N0.02-th largest measurement point from the minimum hardness to the N0.98-th largest measurement point

s*: standard deviation from the N0.02-th largest measurement point from the minimum hardness to the N0.98-th largest measurement point

In addition, when the kurtosis K* exceeds −0.40, the steel sheet structure is not a structure which has a sufficient variety of sufficiently narrow distribution of hardness, dispertion in the hardness distribution in the steel sheet becomes larger, the number of the borders at which the regions with the large hardness difference are in contact with each other increases, and it is not possible to sufficiently suppress degradation of stretch-flangeability.

Next, detailed description will be given of the dispertion in the hardness distribution in the steel sheet with reference to FIG. 1. FIG. 1 is a graph showing a relationship between hardness classified into a plurality of levels and a number of measurement values in each level, which is obtained by converting each measurement value while a difference between a maximum hardness measurement value and a minimum hardness measurement value of the hardness is regarded as 100%, in relation to an example of a high-strength steel sheet according to the present invention. In the graph shown in FIG. 1, the horizontal axis represents hardness, and the vertical axis represents a number of measurement values in each level. In addition, a solid line of the graph shown in FIG. 1 is obtained by connecting the point representing the numbers of the measurement values in each level.

In the high-strength steel sheet of the present invention, it is preferable that all numbers of the measurement values in divided ranges D, which are obtained by equally dividing a range from the 2% hardness to the 98% hardness into 10 parts, in the graph shown in FIG. 1 be within a range from 2% to 30% of the number of all measurement values.

In such a high-strength steel sheet, the line joining up the numbers of the measurement values in the levels becomes a smooth curve with no steep peaks and valleys in the graph shown in FIG. 1, and the dispertion in the hardness distribution in the steel sheet is significantly reduced. Accordingly, such a high-strength steel sheet has less borders at which the regions with large hardness difference are in contact with each other, and excellent stretch-flangeability can be obtained.

In addition, if any of the numbers of the measurement values in a divided range D, which has been equally divided into 10 parts, is outside the range from 2% to 30% of the number of total measurement values in the graph shown in FIG. 1, the line joining up the numbers of the measurement values in the levels may easily include a steep peak or a valley, and an effect that stretch-flangeability is enhanced due to low dispertion in the hardness distribution in the steel sheet is reduced.

Specifically, for example, when only a number of the measurement values in a divided range D near the center exceeds 30% of the number of all measurement values among the equally divided 10 regions D, the line joining up the numbers of the measurement numbers in the levels has a peak in the divided range D near the center.

In addition, if only a number of the measurement values in the divided range D near the center are less than 2% of the number of all measurement values, the line joining up the numbers of the measurement values in the levels has a valley in the divided range D near the center, and many structures have large hardness differences, in which the hardness in different divided ranges D arranged on both sides of the valley is included.

In the high-strength steel sheet of the present invention, all numbers of the measurement values in the divided ranges D are preferably 25% or less of the number of all measurement values, and more preferably 20% or less, in order to further enhance stretch-flangeability. In order to still further enhance stretch-flangeability, all numbers of the measurement values in the divided ranges D are preferably 4% or more of the number of all measurement values, and more preferably 5% or more.

The hardness distribution in the high-strength steel sheet of the present invention will be compared with a general normal distribution and described in detail. The kurtosis K* of the normal distribution is generally considered to be 0. On the other hand, the kurtosis of the hardness distribution in the steel sheet according to the present invention is −0.4 or less, and therefore, it is obvious that the distribution is different from the normal distribution. The hardness distribution in the steel sheet according to the present invention is flatter and has a wider bottom as compared with the normal distribution as shown in FIG. 2. Since the high-strength steel sheet of the present invention has such a hardness distribution, and the ratio of the 98% hardness to the 2% hardness, which correspond to both sides of the bottom of the distribution, is 1.5 or more times which is extremely large, the hardness difference between the soft part and the hard part in the steel sheet structure is sufficiently large, and high ductility can be obtained. That is, the present inventor found that the hole expanding property is further enhanced when the ratio between the 98% harness and the 2% hardness is larger in the hardness distribution in which the kurtosis is −0.4 or less unlike the conventional hardness distribution. On the other hand, the hole expanding property is considered to be further enhanced as the hardness ratio in the structure is smaller, according to the conventional technique. The conventional technique was based on the assumption of the hardness distribution which is close to the normal distribution, which is basically different from the technique proposed in the present invention.

(Mn Distribution)

In the high-strength steel sheet of the present invention, it is preferable that a difference between a maximum value and a minimum value of Mn concentration in the base iron at a thickness from ⅛ to ⅜ of the steel sheet be equal to or more than 0.40% and equal to or less than 3.50% when converted into a mass percentage in order to obtain the aforementioned hardness distribution.

The difference between the maximum value and the minimum value of the Mn concentration in the base iron at the thickness from ⅛ to ⅜ of the steel sheet is defined as 0.40% or more when converted into a mass percentage because phase transformation proceeds more slowly during continuous annealing after cold rolling as the difference between the maximum value and the minimum value of the Mn concentration is larger and it is possible to reliably generate each transformation product at a desired volume fraction and to thereby obtain the high-strength steel sheet with the aforementioned hardness distribution. More specifically, it is possible to generate a transformation product with relatively high hardness such as martensite in place of a transformation product with relatively low hardness such as ferrite in a balanced manner, and therefore, a sharp peak is not present in the hardness distribution in the high-strength steel sheet, that is, the kurtosis decrease, and a flat hardness distribution curve as shown in FIG. 1 can be obtained. In addition, the width of the hardness distribution is widened by generating various transformation products in a balanced manner, and it is thus possible to set the 98% hardness to be 1.5 or more times as high as the 2% hardness, preferably 3.0 or more times, more preferably more than 3.0 times, further more preferably 3.1 or more times, still further preferably 4.0 or more times, and still further preferably 4.2 or more times.

For example, transformation of a ferrite phase will be described as an example. In a heat treatment process for causing transformation of the ferrite phase, the phase transformation from austenite to ferrite starts relatively early in a region where the Mn concentration is low. On the other hand, the phase transformation from austenite to ferrite starts relatively slowly in the region where the Mn concentration is high as compared with the region where the Mn concentration is low. Therefore, the phase transformation from the austenite to ferrite proceeds more slowly in the steel sheet as the Mn concentration in the steel sheet is more non-uniform and the concentration difference is larger. In other words, a transformation rate, during a period when the volume percentage of the ferrite phase reaches, for example, 50% from 0%, becomes lower.

The above phenomenon similarly occurs in the tempered martensite phase and the remaining hard phase as well as the ferrite phase.

FIG. 3 schematically shows a relationship between a transformation rate and elapsed time of transformation treatment. In the case of the phase transformation from austenite to ferrite, for example, the transformation rate represents a volume percentage of ferrite in the steel sheet structure, and the elapsed time of the transformation treatment represents elapsed time of heat treatment for causing ferrite transformation. In the example of the present invention shown in FIG. 3, the difference between the maximum value and the minimum value of the Mn concentration is relatively large, and a gradient of the curve showing the transformation rate in the entire steel sheet is small (the transformation rate is low). On the other hand, in the comparative example shown in FIG. 4, the difference between the maximum value and the minimum value of the Mn concentration is relatively small, and the gradient of the curve showing the transformation rate in the entire steel sheet is large (the transformation rate is high). For this reason, although the transformation treatment may be terminated during a period from x₁to x₂in order to control the transformation rate (volume percentage) in a range from y₁to y₂(%) in the example shown in FIG. 3, it is necessary to terminate the transformation treatment during a period from x₃to x₄and it is difficult to control treatment time in the example shown in FIG. 4.

When the difference in the Mn concentration is less than 0.40%, it is not possible to sufficiently suppress the transformation rate and achieve a sufficient effect, and therefore, this is set as the lower limit. The difference in the Mn concentration is preferably 0.60% or more, and more preferably 0.80% or more. Although the phase transformation can be more easily controlled as the difference in the Mn concentration is larger, it is necessary to excessively increase the amount of Mn added to the steel sheet in order that the difference in the Mn concentration exceeds 3.50%, and it is preferable that the difference in the Mn concentration be 3.50% or less since there is a concern of cracking of a cast slab and degradation of a welding property. In view of the welding property, the difference in the Mn concentration is more preferably 3.40% or less, and more preferably 3.30% or less.

A method of determining a difference between the maximum value and the minimum value of Mn at the thickness from ⅛ to ⅜ is as follows. First, a sample is obtained while a sheet thickness cross-section which is parallel to the rolling direction of the steel sheet is regarded as an observation surface. Then, EPMA analysis is performed in a thickness range from ⅛ to ⅜ around a thickness of ¼ to measure an Mn amount. The measurement is performed while a probe diameter is set to 0.2 to 1.0 vim and measurement time per one point is set to 10 ms or longer, and the Mn amounts are measured at 1000 or more points based on line analysis or surface analysis.

In the measurement results, points at which the Mn concentration exceeds three times the added Mn concentration are considered to be points at which inclusions such as manganese sulfide are observed. In addition, points at which the Mn concentration is less than ⅓ times the added Mn concentration are considered to be points at which inclusions such as aluminum oxide are observed. Since such Mn concentrations hardly affect the phase transformation behavior in the base iron, the maximum value and the minimum value of the Mn concentration are respectively obtained after the measurement results of the inclusions are excluded from the measurement results. Then, the difference between the thus obtained maximum value and minimum value of the Mn concentration is calculated.

The method of measuring the Mn amount is not limited to the above method. For example, an EMA method or direct observation using a three-dimensional atom probe (3D-AP) may be performed to measure the Mn concentration.

(Steel Sheet Structure)

In addition, the steel sheet structure of the high-strength steel sheet of the present invention includes 10 to 50% of a ferrite phase and 10 to 50% of a tempered martensite phase and a remaining hard phase by volume fractions. In addition, the remaining hard phase includes 10 to 60% of one of or both a bainitic ferrite phase and a bainite phase and 10% or less of a fresh martensite phase by volume fractions. Furthermore, the steel sheet structure may contain 2 to 25% of a retained austenite phase. When the high-strength steel sheet of the present invention has such a steel sheet structure, the hardness difference inside the steel sheet becomes much larger, the average crystal grain size becomes sufficiently small, and therefore, the high-strength steel sheet has further higher strength and excellent ductility and strength-flangeability (hole expanding property).

“Ferrite”

Ferrite is a structure which is effective in enhancing ductility and is preferably contained in the steel sheet structure at 10 to 50% by a volume fraction. The volume fraction of ferrite contained in the steel sheet structure is preferably 15% or more, and more preferably 20% or more in view of ductility. In addition, the volume fraction of ferrite contained in the steel sheet structure is preferably 45% or less, and more preferably 40% or less in order to sufficiently enhance the tensile strength of the steel sheet. When the volume fraction of ferrite is less than 10%, there is a concern that sufficient ductility may not be achieved. On the other hand, ferrite has a soft structure, and therefore, yield stress is lower in some cases when the volume fraction exceeds 50%.

“Bainitic Ferrite and Bainite”

Bainitic ferrite and bainite are structures with a hardness between the hardness of soft ferrite and the hardness of hard tempered martensite and fresh martensite. The high-strength steel sheet of the present invention may contain any one of bainitic ferrite and bainite or may contain both. In order to flatten the hardness distribution inside the steel sheet, a total amount of bainitic ferrite and bainite contained in the steel sheet structure is preferably 10 to 45% by volume fraction. The sum of volume fractions of bainitic ferrite and bainite contained in the steel sheet structure is preferably 15% or more, and more preferably 20% or more in view of stretch-flangeability. In addition, the sum of the volume fractions of bainitic ferrite and bainite is preferably 40% or less, or more preferably 35% or less in order to obtain a satisfactory balance between ductility and yield stress.

When the sum of the volume fractions of bainitic ferrite and bainite is less than 10%, bias occurs in the hardness distribution, and there is a concern that stretch-flangeability may be degraded. On the other hand, when the sum of the volume fractions of bainitic ferrite and bainite exceeds 45%, it becomes difficult to generate appropriate amounts of ferrite and tempered martensite, and the balance between ductility and yield stress is degraded, which is not preferable.

“Tempered Martensite”

Tempered martensite is a structure which greatly enhances the tensile strength and is preferably contained in the steel sheet structure at 10 to 50% by a volume fraction. When the volume fraction of tempered martensite contained in the steel sheet structure is less than 10%, there is a concern that sufficient tensile strength may not be obtained. On the other hand, when the volume fraction of the tempered martensite contained in the steel sheet structure exceeds 50%, it becomes difficult to secure ferrite and retained austenite necessary for enhancing ductility. In order to sufficiently enhance the ductility of the high-strength steel sheet, the volume fraction of tempered martensite is preferably 45% or less, and more preferably 40% or less. In addition, in order to secure tensile strength, the volume fraction of tempered martensite is preferably 15% or more, and more preferably 20% or more.

“Retained Austenite”

Retained austenite is a structure which is effective in enhancing ductility and is preferably contained in the steel sheet structure at 2 to 25% by a volume fraction. When the volume fraction of retained austenite contained in the steel sheet structure is 2% or more, more sufficient ductility can be obtained. In addition, when the volume fraction of retained austenite is 25% or less, the welding property is enhanced without a need for adding a large amount of austenite stabilizer such as C or Mn. In addition, although it is preferable that retained austenite be contained in the steel sheet structure of the high-strength steel sheet according to the present invention since retained austenite is effective in enhancing ductility, retained austenite may not be contained when sufficient ductility can be obtained.

“Fresh Martensite”

Since fresh martensite functions as a start point of fracture and degrades stretch-flangeability while fresh martensite greatly enhances tensile strength, fresh martensite is preferably contained in the steel sheet structure at 10% or less by a volume fraction. In order to enhance stretch-flangeability, the volume fraction of fresh martensite is preferably 5% or less, and more preferably 2% or less.

“Others”

The steel sheet structure of the high-strength steel sheet according to the present invention may contain structures such as pearlite and coarse cementite other than the above structures. However, when large amounts of pearlite and coarse cementite are contained in the steel sheet structure of the high-strength steel sheet, ductility is degraded. For this reason, the volume fraction of pearlite and coarse cementite contained in the steel sheet structure is preferably 10% or less in total, and more preferably 5% or less.

The volume fraction of each structure contained in the steel sheet structure of the high-strength steel sheet according to the present invention can be measured based on the following method, for example.

In relation to the volume fraction of retained austenite, X-ray analysis is performed while a surface at a thickness of ¼, which is parallel to the sheet surface of the steel sheet, is regarded as an observation surface, an area fraction is calculated, and the result thereof can be regarded as the volume fraction.

In relation to the volume fractions of ferrite, bainitic ferrite, bainite, tempered martensite, and fresh martensite, a sample is obtained while a sheet thickness cross-section which is parallel to the rolling direction of the steel sheet is regarded as an observation surface, the observation surface is ground, subjected to nital etching, and observed with a Field Emission Scanning Electron Microscope (FE-SEM) in a thickness range from ⅛ to ⅜ around ¼ of the sheet thickness to measure area fractions, and the results thereof can be regarded as the volume fractions.

In addition, an area of the observation surface observed with the FE-SEM can be a 30 μm sided square, for example, and each structure in the observation surface can be distinguished from each other as follows.

Ferrite is a lump of crystal grains and is a region inside which iron carbide with a long diameter of 100 nm or more is not present. In addition, the volume fraction of ferrite is a sum of the volume fraction of ferrite remaining at the highest heating temperature and the volume fraction of ferrite which is newly produced in a ferrite transformation temperature range. However, it is difficult to directly measure the volume fraction of ferrite during the production. For this reason, a small piece of the cold-rolled steel sheet before passing though the continuous annealing line is cut, the small piece is annealed based on the same temperature history as that when the small piece is made to pass through the continuous annealing line, dispertion in the volume of ferrite in the small piece is measured, and a numerical value calculated with the use of the result is regarded as the volume fraction, in the present invention.

In addition, bainitic ferrite is a group of lath-shaped crystal grains, and iron carbide with a long diameter of 20 nm or more is not contained inside the lath.

In addition, bainite is a group of lath-shaped crystal grains, and a plurality of compounds of iron carbide with a long diameter of 20 nm or more is contained inside the lath, and carbide belongs to a single variant, namely an iron carbide group extending in a same direction. Here, the iron carbide group extending in the same direction denotes that the differences in the extending direction of the iron carbide group are within 5°.

In addition, tempered martensite is a group of lath-shaped crystal grains, a plurality of compounds of iron carbide with a long diameter of 20 nm or more is contained inside the lath, and carbide belongs to a plurality of variants, namely a plurality of iron carbide groups extending in different directions.

Moreover, bainite and tempered martensite can be easily distinguished from each other by observing iron carbide inside the lath-shaped crystal grain using the FE-SEM and examining the extending directions thereof.

In addition, fresh martensite and retained austenite are not sufficiently eroded by the nital etching. Therefore, fresh martensite and retained austenite are apparently distinguished from the aforementioned structures (ferrite, bainitic ferrite, bainite, tempered martensite) in the observation with the FE-SEM.

Accordingly, the volume fraction of fresh martensite is obtained as a difference between an area fraction of a region observed with the FE-SEM, which has not yet been eroded, and an area fraction of retained austenite measured with X rays.

(Concerning Definition of Chemical Compositions)

Next, description will be given of chemical constituents (compositions) of the high-strength steel sheet of the present invention. In addition, [%] in the following description represents [mass %].

“C: 0.050 to 0.400%”

C is contained in order to enhance the strength of the high-strength steel sheet. However, if the C content exceeds 0.400%, a sufficient welding property is not obtained. In view of the welding property, the C content is preferably 0.350% or less, and more preferably 0.300% or less. On the other hand, if the C content is less than 0.050%, the strength is lowered, and it is not possible to secure the maximum tensile strength of 900 MPa or more. In order to enhance the strength, the C content is preferably 0.060% or more, and more preferably 0.080% or more.

“Si: 0.10 to 2.50%”

Si is added in order to suppress temper softening of martensite and enhance the strength of the steel sheet. However, if the Si content exceeds 2.50%, embrittlement of the steel sheet is caused, and ductility is degraded. In view of ductility, the Si content is preferably 2.20% or less, and more preferably 2.00% or less. On the other hand, if the Si content is less than 0.10%, hardness of tempered martensite is lowered to a large degree, and it is not possible to secure a maximum tensile strength of 900 MPa or more. In order to enhance the strength, the lower limit value of Si is preferably 0.30% or more, and more preferably 0.50% or more.

“Mn: 1.00 to 3.50%”

Since Mn is an element which enhances the strength of the steel sheet, and it is possible to control the hardness distribution in the steel sheet by controlling the Mn distribution in the steel sheet, Mn is added to the steel sheet of the present invention. However, if the Mn content exceeds 3.50%, a coarse Mn concentrated part is generated at the center in the sheet thickness of the steel sheet, embrittlement easily occurs, and problems such as cracking of a cast slab easily occur. In addition, if the Mn content exceeds 3.50%, the welding property is also degraded. For this reason, it is necessary that the Mu content be 3.50% or less. In view of the welding property, the Mn content is preferably 3.20% or less, and more preferably 3.00% or less. On the other hand, if the Mn content is less than 1.00%, a large amount of soft structures are formed during cooling after annealing, which makes it difficult to secure the maximum tensile strength of 900 MPa or more, and therefore, it is necessary that the Mn content be 1.00% or more. In order to enhance the strength, the Mn content is preferably 1.30% or more, and more preferably 1.50% or more.

“P: 0.001 to 0.030%”

P tends to be segregated at the center in the sheet thickness of the steel sheet and brings about embrittlement of a welded part. If the P content exceeds 0.300%, significant embrittlement of the welded part occurs, and therefore the P content is limited to 0.030% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the P content, 0.001% is set as the lower limit value since manufacturing costs greatly increase when the P content is less than 0.001%.

“S: 0.0001 to 0.0100%”

S adversely affects the welding property and manufacturability during casting and hot rolling. For this reason, the upper limit of S content is set to 0.0100% or less. In addition, since S is bonded to Mn to form coarse MnS and lowers the stretch-flangeability, S is preferably contained at 0.0050% or less, and more preferably contained at 0.0025% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of S content, 0.0001% is set as the lower limit value since manufacturing costs greatly increase when the S content is less than 0.0001%.

“Al: 0.001% to 2.500%”

Al is an element which suppresses production of iron carbide and enhances the strength. However, if an Al content exceeds 2.50%, a ferrite fraction in the steel sheet excessively increases, and the strength is rather lowered, therefore the upper limit of the Al content is set to 2.500%. The Al content is preferably 2.000% or less, and more preferably 1.600% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the Al content, 0.001% is set as the lower limit since an effect as a deoxidizing agent can be obtained when the Al content is 0.001% or more. In order to obtain sufficient effect as the deoxidizing agent, the Al content is preferably 0.005% or more, and more preferably 0.010% or more.

“N: 0.0001 to 0.0100%”

Since N forms coarse nitride and degrades the stretch-flangeability, it is necessary to suppress the added amount thereof. If the N content exceeds 0.0100%, this tendency is more evident, and therefore, the range of the N content is set to 0.0100% or less. In addition, since N causes a blow hole during welding in many cases, it is preferable that the amount of N is as small as possible. Although the effects of the present invention can be achieved without particularly determining the lower limit of the N content, 0.0001% is set as the lower limit value since manufacturing costs greatly increase when the N content is less than 0.0001%.

“O: 0.0001 to 0.0080%”

Since O forms oxide and degrades the stretch-flangeability, it is necessary to suppress the added amount thereof. If the O content exceeds 0.0080%, the degradation of the stretch-flangeability is more evident, and therefore, the upper limit of the O content is set to 0.0080% or less. The O content is preferably 0.0070% or less, and more preferably 0.0060% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the O content, 0.0001% is set as the lower limit value since manufacturing costs greatly increase when the O content is less than 0.0001%.

The high-strength steel sheet of the present invention may further contain the following elements as necessary.

“Ti: 0.005 to 0.090%”

Ti is an element which contributes to enhancement of the strength of the steel sheet by precipitation strengthening, fine grain strengthening by suppressing growth of the ferrite crystal grains, and dislocation strengthening by suppressing recrystallization. However, if a Ti content exceeds 0.090%, the number of precipitate of carbonitride increases, formability is degraded, and therefore, the Ti content is preferably 0.090% or less. In view of the formability, the Ti content is preferably 0.080% or less, and more preferably 0.70% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the Ti content, the Ti content is preferably 0.005% or more in order to sufficiently obtain the effect of Ti enhancing the strength. In order to further enhance the strength of the steel sheet, the Ti content is preferably 0.010% or more, and more preferably 0.015% or more.

“Nb: 0.005 to 0.090%”

Nb is an element which contributes to enhancement of the strength of the steel sheet by precipitation strengthening, fine grain strengthening by suppressing growth of ferrite crystal grains, and dislocation strengthening by suppressing recrystallization. However, if the Nb content exceeds 0.090%, the number of precipitate of carbonitride increases, formability is degraded, and therefore, the Nb content is preferably 0.090% or less. In view of formability, the Nb content is preferably 0.070% or less, and more preferably 0.050% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the Nb content, the Nb content is preferably 0.005% or more in order to sufficiently obtain the effect of Nb enhancing the strength. In order to further enhance the strength of the steel sheet, the Nb content is preferably 0.010% or more, and more preferably 0.015% or more.

“V: 0.005 to 0.090%”

V is an element which contributes to enhancement of the strength of the steel sheet by precipitation strengthening, fine grain strengthening by suppressing growth of ferrite crystal grains, and dislocation strengthening by suppressing recrystallization. However, if the V content exceeds 0.090%, the number of precipitate of carbonitride increases, formability is degraded, and therefore, the Nb content is preferably 0.090% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the V content, the V content is preferably 0.005% or more in order to sufficiently obtain the effect of V enhancing the strength.

“B: 0.0001 to 0.0100%”

Since B delays phase transformation from austenite in a cooling process after hot rolling, it is possible to effectively cause distribution of Mn to proceed by adding B. If the B content exceeds 0.0100%, workability at a high temperature deteriorates, productivity is lowered, and therefore, the B content is preferably 0.0100% or less. In view of the productivity, the B content is preferably 0.0050% or less, and more preferably 0.0030% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the B content, the B content is preferably 0.0001% or more in order to sufficiently obtain the effect of B delaying the phase transformation. In order to delay the phase transformation, the B content is preferably 0.0003% or more, and more preferably 0.0005% or more.

“Mo: 0.01 to 0.80%”

Since Mo delays phase transformation from austenite in a cooling process after hot rolling, it is possible to effectively cause distribution of Mn to proceed by adding Mo. If the Mo content exceeds 0.80%, workability at a high temperature deteriorates, productivity is lowered, and therefore, the Mo content is preferably 0.80% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the Mo content, the Mo content is preferably 0.01% or more in order to sufficiently obtain the effect of Mo delaying the phase transformation.

“Cr: 0.01 to 2.00%” “Ni: 0.01 to 2.00%” “Cu: 0.01 to 2.00%”

Cr, Ni, and Cu are elements which enhance contribution to the strength, and one kind or two or more kinds therefrom can be added instead of a part of C and/or Si. If the content of each element exceeds 2.00%, the acid pickling property, the welding property, the workability at a high temperature, and the like are degraded, and therefore, the content of Cr, Ni, and Cu is preferably 2.00% or less, respectively. Although the effects of the present invention can be achieved without particularly determining the lower limit of the content of Cr, Ni, and Cu, the content of Cr, Ni, and Cu is preferably 0.10% or more, respectively, in order to sufficiently obtain the effect of enhancing the strength of the steel sheet.

“Total Content of one kind or two or more kinds from Ca, Ce, Mg, and REM from 0.0001 to 0.5000%”

Ca, Ce, Mg, and REM are elements which are effective in enhancing formability, and it is possible to add one kind or two or more kinds therefrom. However, if the total amount of one or more of Ca, Ce, Mg, and REM exceeds 0.5000%, there is a concern that ductility may deteriorate, on the contrary, and therefore, the total content of the elements is preferably 0.5000% or less. Although the effects of the present invention can be achieved without particularly determining the lower limit of the content of one or more of Ca, Ce, Mg, and REM, the total content of the elements is preferably 0.0001% or more in order to sufficiently obtain the effect of enhancing formability of the steel sheet. In view of the formability, the total content of one or more of Ca, Ce, Mg, and REM is preferably 0.0005% or more, and more preferably 0.0010% or more. In addition, REM is an abbreviation for Rare Earth Metals and represents an element belonging to lanthanoid series. In the present invention, REM and Ce are added in the form of misch metal in many cases, and there is a case in which elements in the lanthanoid series are contained in combination in addition to La and Ce. Even if such elements in the lanthanoid series other than La and Ce are included as inevitable impurities, the effects of the present invention can be achieved. In addition, the effects of the present invention can be achieved even if metal La and Ce are added.

In addition, the high-strength steel sheet of the present invention may be configured as a high-strength zinc-coated steel sheet by forming a zinc-plated layer or an alloyed zinc-plated layer on the surface thereof. By forming the zinc-plated layer on the surface of the high-strength steel sheet, the high-strength steel sheet obtains excellent corrosion resistance. The high-strength steel sheet has excellent corrosion resistance, and excellent adhesion of a coating can be obtained, since the alloyed zinc-plated layer is formed on the surface thereof.

(Manufacturing Method of High-Strength Steel Sheet)

Next, description will be given of a manufacturing method of the high-strength steel sheet of the present invention.

Firstly, in order to manufacture the high-strength steel sheet of the present invention, slab containing the aforementioned chemical constituents (compositions) is firstly casted.

As the slab subjected to hot rolling, continuous cast slab or slab manufactured by a thin slab caster can be used. The manufacturing method of the high-strength steel sheet of the present invention can be adapted to a process such as continuous casting-direct rolling (CC-DR) in which hot rolling is performed immediately after the casting.

In the hot rolling process, it is necessary that a slab heating temperature be 1050° C. or higher. If the slab heating temperature is excessively low, a finish rolling temperature is below an Ar₃transformation temperature, two phase region rolling of ferrite and austenite is performed, a hot-rolled sheet structure becomes a duplex grain structure in which non-uniform grains are mixed, the non-uniform structure remains even after cold rolling and annealing processes, and therefore, ductility and bendability are degraded. In addition, since lowering of the finish rolling temperature causes excessive increase in rolling load, and there is a concern that it may become difficult to perform rolling or a shape of the steel sheet after the rolling may be defective, it is necessary that the slab heating temperature be 1050° C. or higher. Although the effects of the present invention can be achieved without particularly determining the upper limit of the slab heating temperature, it is preferable that the upper limit of the slab heating temperature be 1350° C. or lower since setting of an excessively high heating temperature is not economically preferable.

In addition, the Ar₃temperature is calculated based on the following equation.

Ar₃=901−325×C+33×Si−92×(Mn+Ni/2+Cr/2+Cu/2+Mo/2)+52×Al

In the above equation, C, Si, Mn, Ni, Cr, Cu, Mo, and Al represent content [mass %] of the elements.

In relation to the finish rolling temperature of the hot rolling, a higher temperature among 800° C. and the Ar₃point is set as a lower limit thereof, and 1000° C. is set as an upper limit thereof. If the finish rolling temperature is lower than 800° C., the rolling load during the finish rolling increases, and there is a concern that it may become difficult to perform the hot rolling or the shape of the hot-rolled steel sheet obtained after the hot rolling may be defective. In addition, if the finish rolling temperature is lower than the Ar₃point, the hot rolling becomes two phase region rolling of ferrite and austenite, and the structure of the hot-rolled steel sheet becomes a structure in which non-uniform grains are mixed.

On the other hand, although the effects of the present invention can be achieved without particularly determining the upper limit of the finish rolling temperature, it is necessary to set the slab heating temperature to an excessively high temperature when the finish rolling temperature is set to an excessively high temperature in order to secure the finish rolling temperature. For this reason, it is preferable that the upper limit temperature of the finish rolling temperature be 1000° C. or lower.

A winding process after the hot rolling and a cooling process before and after the winding process are significantly important to distribute Mn. The above Mn distribution in the steel sheet can be obtained by causing the micro structure during slow cooling after the winding to be a two phase structure of ferrite and austenite and performing processing thereon at a high temperature for long time to cause Mn to be diffused from ferrite to austenite.

In order to control the distribution of the Mn concentration in the base iron at the thickness from ⅛ to ⅜ of the steel sheet, it is necessary that the volume fraction of austenite is 50% or more at the thickness from ⅛ to ⅜ when the steel sheet is wound up. If the volume fraction of austenite at the thickness from ⅛ to ⅜ is less than 50%, austenite disappears immediately after the winding due to progression of the phase transformation, and therefore, the Mn distribution does not sufficiently proceed, and the above Mn concentration distribution in the steel sheet cannot be obtained. In order that the Mn distribution effectively proceeds, the volume fraction of austenite is preferably 70% or more, and more preferably 80% or more. On the other hand, if the volume fraction of austenite is 100%, the phase transformation proceeds after the winding, ferrite is produced, the Mn distribution is started, and therefore the upper limit is not particularly provided for the volume fraction of austenite.

In order to enhance the austenite fraction when the steel sheet is wound up, it is necessary that the cooling rate during a period from completion of the hot rolling to the winding be 10° C./second or higher on average. If the cooling rate is lower than 10° C./second, ferrite transformation proceeds during the cooling, and there is a possibility that the volume fraction of austenite during the winding may become less than 50%. In order to enhance the volume fraction of austenite, the cooling rate is preferably 13° C./second or higher, and more preferably 15° C./second or higher. Although the effects of the present invention can be achieved without particularly determining the upper limit of the cooling rate, it is preferable that the cooling rate be 200° C./second or lower since a special facility is required to obtain a cooling rate of higher than 200° C./second and manufacturing costs significantly increase.

Since a thickness of oxide formed on the surface of the steel sheet excessively increases and the acid pickling property is degraded if the steel sheet is wound up at a temperature which exceeds 800° C., the winding temperature is set to 750° C. or lower. In order to enhance the acid pickling property, the winding temperature is preferably 720° C. or lower, and more preferably 700° C. or lower. On the other hand, if the winding temperature is lower than Bs point, the strength of the hot-rolled steel sheet is excessively enhanced, it becomes difficult to perform cold rolling, and therefore, the winding temperature is set to the Bs point or higher. In addition, the winding temperature is preferably 500° C. or higher, more preferably 550° C. or higher, and further more preferably 600° C. or higher in order to enhance the austenite fraction after the winding.

Moreover, since it is difficult to directly measure the volume fraction of austenite during the production, a small piece is cut from the slab before the hot rolling, the small piece is rolled or compressed at the same temperature and rolling reduction as those in the final pass of the hot rolling and cooled with water immediately after cooling at the same cooling rate as that during a period from the hot rolling and the winding, phase fractions of the small piece are measured, and a sum of the volume fractions of as-quenched martensite, tempered martensite, and retained austenite is regarded as a volume fraction of austenite during the winding, in determining the volume fraction of austenite during the winding according to the present invention.

The cooling process of the steel sheet after the winding is important to control the Mn distribution. The Mn distribution according to the present invention can be obtained by cooling the steel sheet from the winding temperature to (winding temperature−100)° at a rate of 20° C./hour or lower while the austenite fraction is set to 50% or more during the winding and the following equation (3) is satisfied. Equation (3) is an index representing the degree of progression of the Mn distribution between ferrite and austenite and represents that the Mn distribution further proceeds as the value of the left side becomes greater. In order to further cause the Mn distribution to proceed, the value of the left side is preferably 2.5 or more, and more preferably 4.0 or more. Although the effects of the present invention can be achieved without particularly determining the upper limit of the value of the left side, it is preferable that the upper limit is 50.0 or less since it is necessary to retain heat for long time to keep the value over 50.0 and the manufacturing costs significantly increase.

$\begin{matrix} [Equation 3] \\ {[\int_{T_{c - 100}}^{T_{c}} 9.47 \times 10^{5} \cdot \exp (- \frac{18480}{T + 273}) \cdot t (T) \cdot ⅆ T]}^{0.5} \geq 1.0 & (3) \end{matrix}$

Tc: winding temperature (° C.)

T: steel sheet temperature (° C.)

t(T): maintaining time at temperature T (second)

In order to cause the Mn distribution to proceed between ferrite and austenite, it is necessary to maintain a state where both the two phases coexist. If the cooling rate from the winding temperature to (winding temperature−100)° C. exceeds 20° C./hour, the phase transformation excessively proceeds, austenite in the steel sheet may disappear, and therefore, the cooling rate from the winding temperature to (winding temperature−100)° C. is set to 20° C./hour or lower. In order to cause the Mn distribution to proceed, the cooling rate from the winding temperature to (winding temperature−100)° C. is preferably 17° C./hour or lower, and more preferably 15° C./hour or lower. Although the effects of the present invention can be achieved without particularly determining the lower limit of the cooling rate, it is preferable that the lower limit be 1° C./hour or higher since it is necessary to perform heat retaining for a long period of time in order to keep the cooling rate at lower than 1° C./hour and the manufacturing costs significantly increase.

In addition, the steel sheet may be reheated after the winding within a range of satisfying Equation (3) and the cooling rate.

Acid pickling is performed on the thus manufactured hot-rolled steel sheet. Acid pickling is important to enhance a phosphatability of the cold-rolled high-strength steel sheet as a final product and a hot dipping zinc-plating property of the cold-rolled steel sheet for a galvanized steel sheet or a galvannealed a steel sheet since oxide on the surface of the steel sheet can be removed by pickling. In addition, the acid pickling may be performed once or a plurality of times.

Next, the hot-rolled steel sheet after the acid pickling is subjected to cold rolling at rolling reduction from 35 to 80% and is made to pass through a continuous annealing line or a continuous galvanizing line. By setting the rolling reduction to 35% or higher, it is possible to maintain the flattened shape and enhance the ductility of the final product.

In order to enhance the stretch-flangeability, it is preferable that regions where the Mn concentration is high and regions where the Mn concentration is low have a narrow distribution in distributing Mn in the subsequent process. In order to do so, it is effective to increase the rolling reduction during the cold rolling, recrystallize ferrite during temperature increase, and make grain diameters be fine. In such a viewpoint, the rolling reduction is preferably 40% or higher, and more preferably 45% or higher.

On the other hand, in the case of cold rolling at the rolling reduction of 80% or lower, the cold rolling load is not excessively large, and it is not difficult to perform the cold rolling. For this reason, the upper limit of the rolling reduction is set to 80% or lower. In view of the cold rolling load, the rolling reduction is preferably 75% or lower.

In addition, the effects of the present invention can be achieved without particularly determining the number of rolling passes and rolling reduction of each pass. In addition, the cold rolling may be omitted.

Next, the obtained cold-rolled steel sheet is caused to pass through the continuous annealing line to manufacture the high-strength cold-rolled steel sheet. In relation to a process in which the cold-rolled steel sheet is caused to pass through the continuous annealing line, a detailed description will be given of a temperature history of the steel sheet when the steel sheet is caused to pass through the continuous annealing line, with reference to FIG. 5.

FIG. 5 is a graph illustrating the temperature history of the cold-rolled steel sheet when the cold-rolled steel sheet is caused to pass through the continuous annealing line, which is a graph showing the relationship between the temperature of the cold-rolled steel sheet and time. In FIG. 5, a range from (the Ae3 point −50° C.) to the Bs point is shown as a “ferrite transformation temperature region”, a range from the Bs point to the Ms point is shown as the “bainite transformation temperature range”, and a range from the Ms point to a room temperature is shown as the “martensite transformation temperature range”.

In addition, the Bs point is calculated based on the following equation:

Bs point [° C.]=820−290 C/(1−VF)−37Si−90Mn−65Cr−50Ni+70Al

In the above equation, VF represents the volume fraction of ferrite, and C, Mn, Cr, Ni, Al, and Si represent added amounts [mass %] of the elements.

In addition, the Ms point is calculated based on the following equation:

Ms point [° C.]=541−474 C/(1−VF)−15Si−35Mn−17Cr−17Ni+19Al

In the above equation, VF represents a volume fraction of ferrite, C, Si, Mn, Cr, Ni, and Al represent added amounts [mass %] of the elements. In addition, since it is difficult to directly measure the volume fraction of ferrite during the production, a small piece of the cold-rolled steel sheet before the cold-rolling sheet is made to pass through the continuous annealing line is cut and annealed based on the same temperature history as that when the small piece is caused to pass through the continuous annealing line, dispertion in the volume of ferrite in the small piece is measured, and a numerical value calculated using the result of the measurement is regarded as the volume fraction VF of ferrite, in determining the Ms point in the present invention.

As shown in FIG. 5, a heating process for annealing the cold-rolled steel sheet at a maximum heating temperature (T₁) ranging from 750° C. to 1000° C. is firstly performed in causing the cold-rolled steel sheet to pass through the continuous annealing line. If the maximum heating temperature T₁in the heating process is lower than 750° C., the amount of austenite is insufficient, and it is not possible to secure a sufficient amount of hard structures in the phase transformation during the subsequent cooling. From this viewpoint, the maximum heating temperature T₁is preferably 770° C. or higher. On the other hand, if the maximum heating temperature T₁exceeds 1000° C., the grain diameter of austenite becomes coarse, the transformation hardly proceeds during the cooling, and it becomes difficult to sufficiently obtain a soft ferrite structure, in particular. From this viewpoint, the maximum heating temperature T₁is preferably 900° C. or lower.

Next, a first cooling process for cooling the cold-rolled steel sheet from the maximum heating temperature T₁to the ferrite transformation temperature range or lower is performed as shown in FIG. 5. In the first cooling process, the cold-rolled steel sheet is maintained in the ferrite transformation temperature range for 20 seconds to 1000 seconds. In order to sufficiently produce a soft ferrite structure, it is necessary that the cold-rolled steel sheet be maintained for 20 seconds or longer in the ferrite transformation temperature range in the first cooling process, and the cold-rolled steel sheet is preferably maintained for 30 seconds or longer, and more preferably maintained for 50 seconds or longer. On the other hand, if the time during which the cold-rolled steel sheet is maintained in the ferrite transformation temperature range exceeds 1000 seconds, the ferrite transformation excessively proceeds, an amount of untransformed austenite decreases, and it is not possible to sufficiently obtain a hard structure.

In addition, a second cooling process in which the cold-rolled steel sheet after being maintained in the ferrite transformation temperature range for 20 seconds to 1000 seconds to cause ferrite transformation in the first cooling process is cooled at a second cooling rate and the cooling is stopped within a range from the Ms point −120° C. to the Ms point (the martensite transformation start temperature) is performed as shown in FIG. 5. By performing the second cooling process, it is possible to cause the martensite transformation of the untransformed austenite to proceed.

If the second cooling stop temperature T₂at which the second cooling process is stopped exceeds the Ms point, martensite is not produced. On the other hand, if the second cooling stop temperature T₂is lower than the Ms point−120° C., most parts of the untransformed austenite become martensite, and it is not possible to obtain a sufficient amount of bainite in the subsequent processes. In order to cause a sufficient amount of untransformed austenite to remain, the second cooling process stop temperature T₂is preferably the Ms point−80° C. or higher, and more preferably the Ms point−60° C. or higher.

In addition, it is preferable to prevent the bainite transformation from excessively proceeding in the bainite transformation temperature range, which is a temperature range between the ferrite transformation temperature range and the martensite transformation temperature range, in cooling the steel sheet from the ferrite transformation temperature range to the martensite transformation temperature range at the second cooling rate in the second cooling process. For this reason, it is necessary to set the second cooling rate in the bainite transformation temperature range to 10° C./second or higher on average, and the second cooling rate is preferably 20° C./second or higher, and more preferably 50° C./second or higher.

After performing the second cooling process which stops the cooling in a range from the Ms point−120° C. to the Ms point, as shown in FIG. 5, a maintaining process in which the steel sheet is maintained within a range from the second cooling stop temperature to the Ms point for 2 seconds to 1000 seconds in order to cause the martensite transformation to further proceed is performed. In the maintaining process, it is necessary to maintain the steel sheet for 2 seconds or longer in order to cause the martensite transformation to sufficiently proceed. If the time during which the steel sheet is maintained exceeds 1000 seconds in the maintaining process, hard lower bainite is produced, an amount of untransformed austenite is reduced, and bainite with a hardness which is close to that of ferrite cannot be obtained.

Moreover, after maintaining the steel sheet in within the range from the second cooling stop temperature to the Ms point and causing the martensite transformation to proceed as shown in FIG. 5, a reheating process for reheating the steel sheet is performed in order to produce bainite with a hardness between the hardness of ferrite and the hardness of martensite. A temperature T₃(reheating stop temperature) at which the reheating is stopped in the reheating process is set to the Bs point (Bainite transformation start temperature (the upper limit of the bainite transformation temperature range))−100° C. or higher in order to reduce the dispertion in the hardness distribution in the steel sheet.

In order to further reduce the dispertion in the hardness distribution in the steel sheet, it is preferable to produce soft bainite with a small hardness different from that of ferrite. In order to produce soft bainite, the bainite transformation is preferably caused to proceed at a temperature which is as high as possible. Accordingly, the reheating stop temperature T₃is preferably the Bs point−60° C. or higher, and is more preferably the Bs point or higher as shown in FIG. 5.

In the reheating process, it is necessary that the rate of temperature increase in the bainite transformation temperature range be 10° C./second or higher on average, and the rate of temperature increase is preferably 20° C./second or higher, and more preferably 40° C./second or higher. Since the bainite transformation excessively proceeds in a state of the low temperature range if the rate of temperature increase in the bainite transformation temperature range is low in the reheating process, hard bainite with a large hardness difference from that of ferrite is easily produced, and soft bainite with a small hardness difference from that of ferrite, which can reduce the dispertion in the hardness distribution in the steel sheet, is not easily produced. Accordingly, it is preferable that the rate of temperature increase in the bainite transformation temperature range be high in the reheating process.

According to this embodiment, a sum (total maintaining time) of the time during which the steel sheet is maintained in the bainite transformation temperature range in the second cooling process and the time during which the steel sheet is maintained in the bainite transformation range in the reheating process is preferably 25 seconds or shorter, and more preferably 20 seconds or shorter, in order to suppress the excessive progression of the bainite transformation in the second cooling process and the reheating process.

In addition, a third cooling process for cooling the steel sheet from the reheating stop temperature T3 to a temperature which is lower than the bainite transformation temperature range is performed after the reheating process as shown in FIG. 5. In the third cooling process, the steel sheet is maintained in the bainite transformation temperature range for 30 seconds or longer in order to cause the bainite transformation to proceed. In order to obtain a sufficient amount of bainite, the steel sheet is preferably maintained in the bainite transformation temperature range for 60 seconds or longer in the third process, and more preferably maintained for 120 seconds or longer. Although the upper limit of the time during which the steel sheet is maintained in the bainite transformation temperature range in the third cooling process is not particularly provided, the upper limit is preferably 2000 seconds or shorter, and more preferably 1000 seconds or shorter. If the time during which the steel sheet is maintained in the bainite transformation temperature range is 2000 seconds or shorter, it is possible to cool the steel sheet to the room temperature before completion of the bainite transformation of untransformed austenite and to thereby further enhance the yield stress and the ductility of the high-strength cold-rolled steel sheet by changing the untransformed austenite into martensite or retained austenite.

Moreover, a fourth cooling process for cooling the steel sheet from the temperature which is lower than the bainite transformation temperature range to room temperature is performed after the third cooling process as shown in FIG. 5. Although the cooling rate in the fourth cooling process is not particularly defined, it is preferable that the average cooling rate be 1° C./second or higher in order to change untransformed austenite into martensite or retained austenite.

As a result of the above processes, it is possible to obtain a high-strength cold-rolled steel sheet with high ductility and high stretch-flangeability.

Furthermore, a high-strength zinc-coated steel sheet may also be obtained in the present invention by performing zinc electroplating on the high-strength cold-rolled steel sheet obtained by causing the steel sheet to pass through the continuous annealing line based on the aforementioned method.

In addition, the high-strength zinc-coated steel sheet may also be manufactured in the present invention by the following method using the cold-rolled steel sheet obtained based on the above method.

That is, the high-strength zinc-coated steel sheet can be manufacturing in the same manner as the aforementioned case in which the cold-rolled steel sheet is caused to pass through the continuous annealing line except that the cold-rolled steel sheet is dipped into a zinc plating bath in the reheating process.

In so doing, it is possible to obtain the high-strength zinc-coated steel sheet with high ductility and high stretch-flangeability, the surface of which includes a zinc-plated layer formed thereon.

Furthermore, when the cold-rolled steel sheet is dipped into the zinc plating bath in the reheating process, the plated layer on the surface may be alloyed by setting the reheating stop temperature T₃during the reheating process to 460° C. to 600° C. and performing alloying processing in which the cold-rolled steel sheet after being dipped into the zinc plating bath is maintained at the reheating stop temperature T₃for two or more seconds.

By performing such alloying processing, Zn—Fe alloy obtained by alloying the zinc plating layer is formed on the surface, and the high-strength zinc-coated steel sheet with the alloyed zinc plated layer provided on the surface thereof can be obtained.

In addition, the manufacturing method of the high-strength zinc-coated steel sheet is not limited to the above example, and the high-strength zinc-coated steel sheet may be manufactured by performing the same processing as that in the aforementioned case in which the cold-rolled steel sheet is caused to pass through the continuous annealing line other than that the steel sheet is dipped into the zinc plating bath in the bainite transformation temperature range in the third cooling process, for example.

In so doing, the high-strength zinc-coated steel sheet with high ductility and high stretch-flangeability, the surface of which includes the zinc-plated layer formed thereon, can be obtained.

When the steel sheet is dipped into the zinc plating bath in the bainite transformation temperature range in the third cooling process, the plated layer on the surface may be alloyed by performing alloying processing in which the cold-rolled steel sheet after being dipped into the zinc plating bath is reheated again up to 460° C. to 600° C. and maintained for 2 seconds or longer.

Even when such alloying processing is performed, Zn—Fe alloy which is obtained by alloying the zinc plated layer is formed on the surface, and the high-strength zinc-coated steel sheet which includes the alloyed zinc plated layer on the surface thereof can be obtained.

In addition, rolling for shape correction may be performed on the cold-rolled steel sheet after the annealing in this embodiment. However, since work-hardening of the soft ferrite part occurs and the ductility is significantly degraded if the rolling reduction after the annealing exceeds 10%, the rolling reduction is preferably less than 10%.

In addition, the present invention is not limited to the above examples.

For example, plating of one or a plurality of Ni, Cu, Co, and Fe may be performed on the steel sheet before the annealing in order to enhance plating adhesion in the manufacturing method of the high-strength zinc-coated steel sheet according to the present invention.

EXAMPLES

Slab containing chemical constituents A to AQ shown in Tables 1, 2, 19, and 20 was cast, hot rolling was performed thereon under conditions (hot rolling slab heating temperature, finish rolling temperature) shown in Tables 3, 4, 21, 22, and 29, and winding was performed under conditions (cooling rate after rolling, winding temperature, cooling rate after winding) shown in Tables 3, 4, 21, 22, and 29. Then, after acid pickling, cold rolling was performed at “rolling reduction” shown in Tables 3, 21, and 22 to obtain the cold-rolled steel sheets with thicknesses in Experiment Examples a to bd and Experiment Examples ca to ds shown in Tables 3, 21, and 22. In addition, acid picking was performed after the winding, and cold rolling was not performed thereon to obtain the hot-rolled steel sheet with thicknesses in Experiment Examples dt to dz shown in Table 29.

Thereafter, the cold-rolled steel sheet in Experiment Examples a to bd and Experiment Examples ca to ds and the hot-rolled steel sheet in Experiment Examples dt to dz were caused to pass through the continuous annealing line to manufacture the steel sheets in Experiment Examples 1 to 134.

In causing the steel sheets to pass through the continuous annealing line, the high-strength cold-rolled steel sheets in Experiment Examples 1 to 134 were obtained based on the following method under conditions shown in Tables 5 to 12, 23 to 25, 30, and 31 (a maximum heating temperature in a heating process, maintaining time in a ferrite transformation temperature range in a first cooling process, a cooling rate in bainite transformation temperature range in a second cooling process, a cooling stop temperature in the second cooling process, maintaining time in a maintaining process, a rate of temperature increase in the bainite transformation temperature range and the reheating stop temperature in a reheating process, maintaining time in the bainite transformation temperature range in a third cooling process, the cooling rate in a fourth cooling process, a sum of a time during which the steel sheet is maintained in the bainite transformation temperature range in the second cooling process and a time during which the steel sheet is maintained in the bainite transformation range in the reheating process (total maintaining time)).

That is, the heating process for annealing the cold-rolled steel sheet in Experiment Examples a to bd and Experiment Examples ca to ds and the hot-rolled steel sheet in Experiment Examples dt to dz, the first cooling process for cooling the cold-rolled steel sheet from the maximum heating temperature to the ferrite transformation temperature range or lower, the second cooling process for cooling the cold-rolled steel sheet after the first cooling process, the maintaining process for maintaining the cold-rolled steel sheet after the second cooling process, the reheating process for reheating the cold-rolled steel sheet after the maintaining process up to the reheating stop temperature, the third cooling process for cooling the cold-rolled steel sheet after the reheating process from the reheating stop temperature to the temperature which is lower than the bainite transformation temperature range, in which the cold-rolled steel sheet is maintained in the bainite transformation temperature range for 30 seconds or longer, and the fourth cooling process for cooling the steel sheet from the temperature which is lower than the bainite transformation temperature range to the room temperature are performed.

As a result of the above processes, the high-strength cold-rolled steel sheets and the high-strength hot-rolled steel sheets in Experiment Examples 1 to 134 were obtained.

Thereafter, a part of Experiment Examples in which the steel sheets were caused to pass through the continuous annealing line, namely the cold-rolled steel sheets in Experiment Examples 60 to 63 were subjected to the zinc electroplating based on the following method to manufacture the zinc-electroplated steel sheet (EG) in Experiment Examples 60 to 63.

First, alkaline degreasing, rinsing with water, acid pickling, and rinsing with water were performed on the steel sheet, which had passed through the continuous annealing line, as pre-processing for plating. Thereafter, electrolytic treatment was performed on the steel sheet after the pre-processing using a liquid circulation type electroplating device with a plating bath containing zinc sulfate, sodium sulfate, and sulfuric acid at a current density of 100 A/dm²up to a predetermined plating thickness, and Zn plating was performed.

In relation to the cold-rolled steel sheets in Experiment Examples 64 to 68, the cold-rolled steel sheets were dipped into the zinc plating bath in the reheating process when the cold-rolled steel sheet was caused to pass through the continuous annealing line and the high-strength zinc-coated steel sheets were obtained.

In addition, in relation to the cold-rolled steel sheets in Experiment Examples 69 to 73, the cold-rolled steel sheets after being dipped into the zinc plating bath in the reheating process were subjected to the alloying processing, in which the cold-rolled steel sheets were maintained at the “reheating stop temperature T₃” shown in Table 11 for the “maintaining time” shown in Table 12 to alloy the plated layer on the surface thereof, and the high-strength zinc-coated steel sheets with alloyed zinc-plated layers were obtained.

In relation to the cold-rolled steel sheet in Experiment Examples 74 to 77, the cold-rolled steel sheets were dipped into the zinc plating bath in the third cooling process when the cold-rolled steel sheets were caused to pass through the continuous annealing line, and the high-strength zinc-coated steel sheets were obtained.

In relation to the cold-rolled steel sheets in Experiment Examples 78 to 82, the cold-rolled steel sheets after being dipped into the zinc plating bath in the third cooling process were subjected to the alloying process in which the cold-rolled steel sheets were reheated again up to the “alloying temperature Tg” shown in Table 12 and maintained for the “maintaining time” shown in Table 12 to alloy the plated layers on the surfaces thereof, and the high-strength zinc-coated steel sheets with alloyed zinc-plated layers were obtained.

In relation to the hot-rolled steel sheet in Experiment Example 130, the high-strength zinc-coated steel sheet with the alloyed zinc-plated layer was obtained by dipping the steel sheet which was made to pass through the continuous annealing line into the zinc plating bath, then performing thereon alloying processing in which the steel sheet was reheated again up to the “alloying temperature Tg” shown in Table 31 and maintained for the “maintaining time” shown in Table 31, and thereby alloyed the plated layer on the surface thereof.

In relation to the hot-rolled steel sheet in Experiment Example 132, the high-strength zinc-coated steel sheet with the alloyed zinc-plated layer was obtained by dipping the hot-rolled steel sheet into the zinc plating bath when the hot-rolled steel sheet was caused to pass through the continuous annealing line, performing thereon alloying processing in which the hot-rolled steel sheet was reheated again up to the “alloying temperature Tg” shown in Table 31 and maintained for the “maintaining time” shown in Table 31, and thereby alloying the plated layer on the surface thereof.

In relation to the hot-rolled steel sheet in Example 134, the steel sheet which was caused pass through the continuous annealing line was dipped into the zinc plating bath, and the high-strength zinc-coated steel sheet was obtained.

In relation to the thus obtained high-strength steel sheets in Experiment Examples 1 to 134, micro structures were observed, and volume fractions of ferrite (F), bainitic ferrite (BF), bainite (B), tempered martensite (TM), fresh martensite (M), and retained austenite (retained γ) were obtained based on the following method. In addition, “B+BF” in the tables represents a total volume fraction of ferrite and bainitic ferrite.

In relation to the volume fraction of retained austenite, an observation surface at a thickness of ¼, which was parallel to the plate surface of the steel sheet, was regarded as an observation surface, X-ray analysis was performed thereon, and an area fraction was calculated and regarded as the volume fraction thereof.

In relation to the volume fractions of ferrite, bainitic ferrite, bainite, tempered martensite, and fresh martensite, a sheet thickness cross-section which was parallel to the rolling direction of the steel sheet was regarded as an observation surface, a sample was collected therefrom, grinding and nital etching were performed on the observation surface, a region surrounded by sides of 30 μm was set at a thickness range from ⅛ to ⅜ around ¼ of the sheet thickness, the region was observed with FE-SEM, and area fractions were measured and regarded as the volume fractions thereof.

The results are shown in Tables 13, 14, 17, 26, and 32.

In relation to the high-strength steel sheets in Experiment Example 1 to 134, sheet thickness cross-section which were parallel to the rolling direction of the steel sheets were finished as mirror surfaces, and EPMA analysis was performed in a range from ⅛ to ⅜ around ¼ of the sheet thicknesses to measure the Mn amounts. The measurement was performed while the probe diameter was set to 0.5 μm and a measurement time for one point was set to 20 ms, and the Mn amounts were measured for 40000 points in the surface analysis. The results are shown in Tables 15, 16, 18, 27, 28, and 33. After removing inclusion measurement results from the measurement results, maximum values and minimum values of the Mn concentration were respectively obtained, and differences between the obtained maximum values and the minimum values of the Mn concentration were calculated. The results will be shown in Tables 15, 16, 18, 27, 28, and 33.

In relation to each of the high-strength steel sheets in Experiment Examples 1 to 134, “a ratio (H98/H2) of a measurement value of the 2% hardness (H2) with respect to a measurement value of the 98% hardness (H98), which was obtained by converting the measurement values while a difference between a maximum measurement value and a minimum measurement value of hardness was regarded as 100%, a kurtosis (K*) between the measurement value of the 2% hardness and the measurement value of the 98% hardness, an average crystal grain size, and whether or not the number of all measurement values in each divided range, which were obtained by equally dividing a range from the 2% hardness to the 98% hardness into 10 parts, were in a range from 2% to 30% of the number of all measurement values in a graph representing a relationship between the hardness classified into a plurality of levels and a number of measurement values in each level when each measurement value was converted while a difference between a maximum value and a minimum value of the hardness measurement values was regarded as 100%” were exemplified. The results are shown in Tables 15, 16, 18, 27, 28, and 33.

In addition, the hardness was measured using a dynamic micro-hardness tester provided with a Berkovich type three-sided pyramid indenter under an indentation load of 1 g based on an indentation depth measurement method. The hardness measurement position was set to a range from ⅛ to ⅜ around ¼ of the sheet thickness in the sheet thickness cross-section which was parallel to the rolling direction of the steel sheet. In addition, the number of measurement values (point number of indentations) was in the range from 100 to 10000 and preferably 1000 or more.

In addition, the average crystal grain size was measured using an EBSD (Electron BackScattering Diffreaction) method. A crystal grain size observation surface was set a range from ⅛ to ⅜ around ¼ of the sheet thickness in the sheet thickness cross-section which was parallel to the rolling direction of the steel sheet. Then, a border, at which a crystal orientation difference between measurement points which were adjacent in the bcc crystal orientation on the observation surface was 15° or more, on the observation surface was regarded as a crystal grain boundary, and crystal grain size was measured. Then, the average crystal grain size was calculated by applying a intercept method to the result (map) of the obtained crystal grain boundary. The results are shown in Tables 13, 14, 17, 26, and 32.

Moreover, tensile test pieces based on JIS Z 2201 were collected from the high-strength steel sheets in Experiment Examples 1 to 134, tensile tests were performed thereon based on JIS Z 2241, and maximum tensile strength (TS) and ductility (EL) were measured. The results are shown in Tables 15, 16, 18, 27, 28, and 33.

As shown in Tables 15, 16, 18, 27, 28, and 33, it was confirmed that the measurement value of the 98% hardness was 1.5 or more times as high as the measurement value of the 2% hardness, that the kurtosis (K*) between the measurement value of the 2% hardness and the measurement value of the 98% hardness was −0.40 or less, that the average crystal grain size was 10 μm or less, and that the steel sheet had excellent maximum tensile strength (TS), ductility (EL), and stretch-flangeability (λ), in Examples of the present invention.

On the other hand, in Experiment Examples 9, 14, 17, 25, 30, 36, 39, 56 to 59, 85, 86, 89, 90, 93, 94, 101, 102, 117, 120, and 123 as Comparative Examples of the present invention, there was no steel sheet in which all the maximum tensile strength (TS), the ductility (EL), and the stretch-flangeability (λ) were sufficient as shown below. Particularly, in Experiment Example 102, the total of the volume fractions of bainite and bainitic ferrite was 50% or more, the K* value was −0.4 or more, that is, the hardness distribution was close to the normal distribution, and therefore, the ductility was low even at a hardness ratio of 4.2.

In Experiment Example 9, the maintaining time in the bainite transformation temperature range was short in the third cooling process in the continuous annealing line, and the bainite transformation did not sufficiently proceed. For this reason, the ratios of bainite and bainitic ferrite were low in Experiment Example 9, the kurtosis (K*) exceeded −0.40, the hardness distribution was not flat and had a “valley”, and therefore, the stretch-flangeability λ deteriorated.

In Experiment Example 14, the rolling reduction in the cold rolling process was below the lower limit, and the degree of flatness of the steel sheet deteriorated. In addition, since the rolling reduction was low, recrystallization did not proceed in the continuous annealing line, the average crystal grain size became coarse, and therefore, the stretch-flangeability λ was lowered.

In Experiment Example 17, the maintaining time in the ferrite transformation temperature range was short in the first cooling process, and the ferrite transformation did not sufficiently proceed. For this reason, a fraction of soft ferrite was low, H98/H2 was below the lower limit, the hardness difference between the hard part and the soft part was small, and the ductility EL deteriorated, in Experiment Example 17.

In Experiment Example 25, since the maintaining time in the ferrite transformation temperature range was long, the ferrite transformation excessively proceeded. In Experiment Example 25, the cooling termination temperature exceeded the Ms point in the second cooling process, and tempered martensite was not sufficiently obtained. For this reason, the stretch-flangeability λ was lowered in Experiment Example 25.

In Experiment Example 30, the cooling termination temperature was below the lower limit in the second cooling process, and it was not possible to cause the bainite transformation to proceed in the third cooling process. For this reason, the ratios of bainite and bainitic ferrite were low, the hardness distribution has a “valley”, and therefore, the stretch-flangeability λ deteriorated in Experiment Example 30.

In Experiment Example 36, the maximum heating temperature exceeded the upper limit, and the cooling termination temperature in the second cooling process was below the lower limit. For this reason, a fraction of tempered martensite increased, the soft structures such as ferrite were not present, and therefore, H98/H2 was below the lower limit, the hardness difference between the hard part and the soft part was small, and the ductility EL deteriorated, in Experiment Example 36.

Experiment Example 39 was an example in which the average cooling rate in the bainite transformation temperature range was low in the second cooling process and the bainite transformation excessively proceeded in the process. In Experiment Example 39, tempered martensite was not present, and therefore, the tensile strength TS was insufficient.

The chemical constituents of the steel sheets in Experiment Examples 56 to 59 were not within the range of definition.

More specifically, the C content in the steel W in Experiment Example 56 was below the lower limit defined in this invention. For this reason, the ratio of soft structure was high, and the tensile strength TS was insufficient, in Experiment Example 56.

In Experiment Example 57, the C content in the steel X exceeded the upper limit. For this reason, the rate of the soft structure was low, and the ductility EL was insufficient, in Experiment Example 57.

In Experiment Example 58, the Si content in the steel Y was below the lower limit. For this reason, the strength of tempered martensite was low, and the tensile strength TS was insufficient in Experiment Example 58.

In Experiment Example 59, the Mn content in the steel Z was below the lower limit. For this reason, a tempering property was significantly lowered, it was not possible to obtain tempered martensite and martensite which had soft structures, and therefore, the tensile strength TS was insufficient, in Experiment Example 59.

In Experiment Examples 85 and 102, the cooling rate from the completion of the hot rolling to the winding was below the lower limit. For this reason, the phase transformation excessively proceeded before the winding, most parts of austenite in the steel sheet disappeared, the Mn distribution did not proceed, and a predetermined micro structure was not obtained in the continuous annealing line, in Experiment Examples 85 and 102. For this reason, the kurtosis K* exceeds the upper limit, and the stretch-flangeability λ was insufficient.

In Experiment Example 86, the maintaining time in the maintaining process in the martensite transformation temperature range in the continuous annealing line was below the lower limit. For this reason, the ratio of tempered martensite was low, the kurtosis (K*) exceeded −0.40, the hardness distribution was not flat and had a “valley”, and therefore, the stretch-flangeability λ was lowered, in Experiment Example 86.

In Experiment Example 89, the winding temperature was below the lower limit. For this reason, the Mn distribution did not proceed, and the predetermined micro structure was not obtained in the continuous annealing line in Experiment Example 89. For this reason, the kurtosis K* exceeded the upper limit, and the stretch-flangeability λ was insufficient.

In Experiment Example 90, the reheating stop temperature in the reheating process in the continuous annealing line was below the lower limit. For this reason, the hardness of produced bainite and bainitic ferrite excessively increased, the hardness difference between the hardness of ferrite and the hardness of bainite and bainitic ferrite increased, the kurtosis (K*) exceeded −0.40, the hardness distribution had a “valley”, and therefore, the stretch-flangeability λ was lowered.

In Experiment Example 93, the cooling rate after the winding exceeded the upper limit. For this reason, the Mn distribution did not proceed, and the predetermined micro structure was not obtained in the continuous annealing line, in Experiment Example 93. Therefore, the kurtosis K* exceeded the upper limit, and the stretch-flangeability λ was insufficient.

In Experiment Example 94, the average rate of temperature increase in the bainite transformation temperature range in the reheating process in the continuous annealing line exceeded the upper limit. For this reason, the hardness of produced bainite and bainitic ferrite excessively increased, the hardness difference between the hardness of ferrite and the hardness of bainite and bainitic ferrite increased, the kurtosis (K*) exceeded −0.40, the hardness distribution had a “valley”, and the therefore, the stretch-flangeability λ was lowered.

In Experiment Example 101, the maintaining time in the maintaining process in the martensite transformation temperature range in the continuous annealing line exceeded the upper limit. For this reason, hard lower bainite was produced, relatively soft bainite and/or bainitic ferrite was not obtained, the kurtosis (K*) exceeded −0.40, the hardness distribution had a “valley”, and therefore, the stretch-flangeability λ was lowered.

In Experiment Example 117, the maximum heating temperature in the continuous annealing line exceeded the upper limit. For this reason, soft ferrite was not obtained, H98/H2 was below the lower limit, the hardness difference between the hard part and the soft part was small, and the ductility EL deteriorated, in Experiment Example 117.

In Example 120, the maximum heating temperature in the continuous annealing line was below the lower limit. For this reason, less hard structure was obtained, and the strength TS deteriorated, in Experiment Example 120.

In Experiment Example 123, the cooling stop temperature in the second cooling process in the continuous annealing line exceeded the upper limit. For this reason, tempered martensite was not obtained, the kurtosis (K*) exceeded −0.40, the hardness distribution had a “valley”, and therefore, the stretch-flangeability λ was lowered, in Experiment Example 123.

INDUSTRIAL APPLICABILITY

Since the high-strength steel sheet of the present invention contains predetermined chemical constituents, the 98% hardness is 1.5 or more times as high as the 2% hardness, the kurtosis K* of the hardness distribution between the 2% hardness and the 98% hardness is −0.40 or less, the average crystal grain size in the steel sheet structure is 10 μm or less, and therefore, the steel sheet has excellent ductility and stretch-flangeability while tensile strength which is as high as 900 MPa or more is secured. Accordingly, the present invention can make very significant contributions to the industry since the strength of the steel sheet can be secured without degrading workability.

TABLE 1

Experiment
C
Si
Mn
P
S
Al
N
O

Example
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %

A
0.185
1.32
2.41
0.006
0.0016
0.043
0.0039
0.0008
Example

B
0.094
1.79
2.65
0.012
0.0009
0.017
0.0020
0.0011
Example

C
0.128
1.02
2.87
0.022
0.0007
0.127
0.0028
0.0014
Example

D
0.234
0.85
2.15
0.005
0.0004
0.233
0.0016
0.0011
Example

E
0.167
1.38
2.16
0.013
0.0021
0.026
0.0030
0.0009
Example

F
0.219
1.47
1.82
0.007
0.0020
0.061
0.0025
0.0020
Example

G
0.242
0.50
2.37
0.007
0.0043
1.175
0.0040
0.0022
Example

H
0.124
1.65
2.14
0.005
0.0043
0.032
0.0050
0.0010
Example

I
0.104
2.28
1.95
0.018
0.0046
0.030
0.0023
0.0018
Example

J
0.076
1.82
2.48
0.018
0.0013
0.064
0.0056
0.0009
Example

K
0.197
0.78
2.82
0.005
0.0021
1.310
0.0054
0.0008
Example

L
0.159
1.09
3.01
0.005
0.0040
0.029
0.0028
0.0016
Example

M
0.088
2.06
2.50
0.020
0.0032
0.015
0.0034
0.0017
Example

N
0.080
1.52
2.01
0.022
0.0023
0.046
0.0032
0.0018
Example

O
0.172
1.33
2.67
0.014
0.0032
0.086
0.0039
0.0043
Example

P
0.223
0.38
3.02
0.009
0.0037
2.304
0.0015
0.0012
Example

Q
0.137
2.08
2.12
0.013
0.0045
0.075
0.0020
0.0015
Example

R
0.143
1.13
1.59
0.004
0.0041
0.020
0.0060
0.0021
Example

S
0.173
0.85
2.37
0.010
0.0004
1.526
0.0048
0.0023
Example

T
0.167
1.95
1.79
0.009
0.0032
0.091
0.0016
0.0016
Example

U
0.211
0.41
2.56
0.012
0.0043
0.683
0.0034
0.0023
Example

V
0.226
1.26
1.68
0.003
0.0029
0.746
0.0014
0.0010
Example

W

0.025

1.99
2.19
0.014
0.0039
0.046
0.0058
0.0021
Comparative Example

X

0.519

1.22
1.84
0.018
0.0047
0.036
0.0033
0.0010
Comparative Example

Y
0.175

0.03

2.14
0.019
0.0036
0.050
0.0034
0.0008
Comparative Example

Z
0.205
0.93

0.57

0.009
0.0037
0.099
0.0020
0.0015
Comparative Example

TABLE 2

Experi-

ment
Ti
Nb
B
Cr
Ni
Cu
Mo
V
Ca
Ce
Mg
REM

Example
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %

A

Example

B

Example

C

0.0016

Example

D

0.0013

Example

E
0.017

Example

F
0.065

0.0014
0.0007

Example

G

0.046

Example

H
0.030

0.0016

0.0014

Example

I

0.0034

Example

J
0.021
0.019

Example

K

0.31

Example

L

0.25

Example

M

0.42

Example

N

0.29

Example

O

0.071

Example

P
0.053

0.0011

0.18

0.0032

Example

Q

0.42

0.22

0.0012
Example

R

1.29
0.10

0.0013

Example

S
0.028

0.0008
0.10
0.27
0.14
0.07

0.0007
0.0009

Example

T

0.027

0.78

0.086

0.0018

Example

U
0.017
0.050

0.60
0.10

0.0028

0.0015
Example

V

0.0029
1.11
0.50

0.039
0.0018
0.0018

Example

W

Comparative Example

X

Comparative Example

Y

Comparative Example

Z

Comparative Example

TABLE 3

Slab
Ar₃
Finish
Cooling

Heating
Transformation
Rolling
Rate After
Winding

Experiment
Chemical
Temperature
Point
Temperature
Rolling
Temperature

Example
Constituent
° C.
° C.
° C.
° C./second
° C.

a
A
1230
665
909
48
630

b
A
1265
665
937
114
576

c
A
1210
665
916
32
674

d
B
1245
687
909
48
526

e
B
1245
687
861
71
601

f
B
1255
687
851
19
606

g
C
1215
636
953
26
614

h
C
1240
636
902
77
617

i
D
1175
667
890
26
573

j
D
1165
667
890
61
528

k
E
1190
695
908
69
608

l
E
1205
695
918
29
654

m
E
1165
695
940
25
653

n
F
1225
714
865
36
561

o
F
1225
714
899
79
542

p
G
1210
682
929
67
555

q
G
1260
682
862
49
537

r
H
1165
720
897
14
581

s
H
1195
720
945
34
528

t
H
1170
720
903
38
663

u
I
1210
765
881
55
533

v
I
1175
765
924
26
613

w
I
1200
765
931
12
559

x
J
1260
712
901
72
627

y
J
1270
712
950
60
573

z
K
1210
657
916
64
547

Cooling
Volume

Cold-rolled

Left Side of
Rate After
Fraction of

Rolling
Sheet

Experiment
Equation
Winding
Austenite
Bs
Reduction
Thickness

Example
(1)
° C./hour
volume %
° C.
%
mm

a
11.2
14
82
492
50
1.6
Example

b
3
13
100
504
50
1.6
Example

c
29.2
15
90
498
68
0.8
Example

d
1.1
8
72
479
40
1.2
Example

e
6.1
12
83
484
60
1.2
Example

f
5.9
14
77
481
60
1.2
Example

g
5.7
18
88
491
60
1.2
Example

h
12.8
9
95
494
60
1.2
Example

i
2.7
13
58
494
50
1.6
Example

j
1.2
9
72
517
50
1.6
Example

k
11.4
8
79
515
60
1.6
Example

l
16
18
72
509
68
0.8
Example

m
24.4
11
78
514
5
2.3
Comparative

Example

n
2.2
12
79
526
50
2
Example

o
1.1
12
78
525
50
2
Example

p
1.5
14
93
595
50
2
Example

q
1.1
11
74
576
50
2
Example

r
2.7
15
78
522
50
2
Example

s
1.1
7
93
530
50
2
Example

t
18.6
19
100
533
72
0.8
Example

u
1.2
10
90
529
38
1.6
Example

v
8.1
13
86
527
38
1.6
Example

w
1.9
13
97
531
38
1.6
Example

x
9.5
15
100
512
38
1.6
Example

y
1.8
18
86
508
38
1.6
Example

z
1.5
12
83
540
50
1.6
Example

TABLE 4

Slab
Ar₃
Finish
Cooling

Heating
Transformation
Rolling
Rate After
Winding

Experiment
Chemical
Temperature
Point
Temperature
Rolling
Temperature

Example
Constituent
° C.
° C.
° C.
° C./second
° C.

aa
K
1165
657
916
59
574

ab
L
1235
598
923
20
521

ac
L
1170
598
908
79
616

ad
M
1245
692
893
71
576

ae
M
1215
692
900
35
611

af
N
1180
729
918
88
629

ag
N
1210
729
830
26
608

ah
N
1155
729
873
38
508

ai
O
1205
648
919
106
538

aj
O
1250
648
949
26
575

ak
O
1255
648
937
49
650

al
P
1165
675
941
58
617

am
P
1165
675
903
34
566

an
Q
1230
705
872
30
571

ao
Q
1210
705
958
68
615

ap
R
1200
683
872
72
607

aq
R
1150
683
899
25
580

ar
S
1265
707
884
25
532

as
S
1210
707
944
63
624

at
S
1205
707
933
96
573

au
T
1265
715
886
37
611

av
T
1160
715
960
68
589

aw
U
1185
614
920
20
620

ax
U
1215
614
909
43
640

ay
V
1190
679
871
54
580

az
V
1205
679
911
43
609

ba

W

1155
759
862
56
651

bb

X

1210
605
939
56
659

bc

Y

1225
651
938
58
655

bd

Z

1180
818
917
50
643

Cooling
Volume

Cold-rolled

Left Side of
Rate After
Fraction of

Rolling
Sheet

Experiment
Equation
Winding
Austenite
Bs
Reduction
Thickness

Example
(1)
° C./hour
volume %
° C.
%
mm

aa
2.2
15
89
545
50
1.6
Example

ab
1.1
6
78
439
50
1.2
Example

ac
9.2
13
100
452
50
1.2
Example

ad
2.7
14
91
492
60
0.8
Example

ae
7.4
15
67
482
60
0.8
Example

af
10.1
16
100
563
50
1.2
Example

ag
7.6
12
73
554
50
1.2
Example

ah
1.2
4
89
560
36
1.2
Example

ai
1.4
9
100
487
60
0.8
Example

aj
2.5
15
80
474
60
0.8
Example

ak
15.7
18
98
486
72
0.8
Example

al
9.7
13
94
618
68
0.8
Example

am
2.8
11
74
599
68
0.8
Example

an
2.7
14
80
481
50
1.6
Example

ao
5.8
20
84
483
50
1.6
Example

ap
8.7
11
84
523
50
1.6
Example

aq
3.4
14
87
524
50
1.6
Example

ar
2.1
5
62
581
50
1.6
Example

as
11
13
86
604
50
1.6
Example

at
2.8
13
89
606
38
1.6
Example

au
17.1
6
87
487
50
1.6
Example

av
4.6
12
79
481
50
1.6
Example

aw
7.8
17
74
540
40
1.6
Example

ax
20.6
11
88
553
40
1.6
Example

ay
3.7
12
78
493
60
1.2
Example

az
6.4
14
76
491
60
1.2
Example

ba
24.7
11
0
—
72
1.4
Comparative

Example

bb
31.6
11
87
439
50
1.4
Comparative

Example

bc
27.5
10
72
559
50
1.6
Comparative

Example

bd
17.7
12
23
483
50
1.6
Comparative

Example

TABLE 5

Second Cooling Process

First Cooling Process
Average

Maximum
Maintaining
Cooling Rate
Cooling
Cooling

Cold-

Heating
Time in Ferrite
in Bainite
Termination
Termination

rolled

Temperature
Transformation
Transformation
Temperature
Temperature -

Experiment
Steel
Chemical
Type of
(T1)
Temperature Range
Temperature Range
(T2)
Ms

Example
Sheet
Constituent
Steel
° C.
second
° C./second
° C.
° C.

1
a
A
CR
822
47
57
257
−52
Example

2
b
A
CR
835
82
64
181
−93
Example

3
c
A
CR
839
39
85
268
−48
Example

4
d
B
CR
845
84
68
236
−99
Example

5
e
B
CR
837
126
60
308
−40
Example

6
f
B
CR
848
79
62
291
−58
Example

7
g
C
CR
831
149
74
270
−64
Example

8
h
C
CR
843
164
74
259
−66
Example

9
h
C
CR
838
150
88
305
−23
Comparative

Example

10
i
D
CR
827
66
83
275
−54
Example

11
j
D
CR
840
78
78
271
−49
Example

12
k
E
CR
803
71
61
219
−94
Example

13
l
E
CR
808
75
79
304
−8
Example

14

m

E
CR
802
70
60
255
−51
Comparative

Example

15
n
F
CR
817
42
59
211
−83
Example

16
o
F
CR
833
49
62
228
−85
Example

17
o
F
CR
880
6
60
272
−81
Comparative

Example

18
p
G
CR
787
85
67
261
−78
Example

19
q
G
CR
865
24
78
282
−60
Example

20
r
H
CR
845
90
67
284
−62
Example

21
s
H
CR
837
77
67
302
−36
Example

22
t
H
CR
872
35
56
309
−62
Example

23
u
I
CR
921
53
68
271
−78
Example

24
v
I
CR
936
42
69
281
−88
Example

25
w
I
CR
888

1730
85

303

50

Comparative

Example

26
x
J
CR
879
67
75
338
−36
Example

27
y
J
CR
852
74
77
304
−69
Example

28
z
K
CR
860
284
62
261
−38
Example

29
aa
K
CR
962
457
85
278
−52
Example

30
aa
K
CR
906
171
88
142

−148
Comparative

Example

TABLE 6

First Cooling

Process
Second Cooling Process

Maintaining
Average

Maximum
Time in Ferrite
Cooling Rate
Cooling
Cooling

Heating
Transformation
in Bainite
Termination
Termination

Temperature
Temperature
Transformation
Temperature
Temperature -

Experiment
Cold-rolled
Chemical
Type of
(T1)
Range
Temperature Range
(T2)
Ms

Example
Steel Sheet
Constituent
Steel
° C.
second
° C./second
° C.
° C.

31
ab
L
CR
809
96
88
274
−47
Example

32
ac
L
CR
814
153
67
247
−67
Example

33
ad
M
CR
846
75
79
274
−70
Example

34
ae
M
CR
843
81
71
292
−58
Example

35
af
N
CR
862
62
56
332
−49
Example

36
ag
N
CR

1035
42
86

272

−139
Comparative

Example

37
ah
N
CR
891
70
71
303
−92
Example

38
ai
O
CR
830
74
70
234
−64
Example

39
aj
O
CR
840
70
1
253
−54
Comparative

Example

40
ak
O
CR
835
70
74
266
−43
Example

41
al
P
CR
905
249
64
207
−65
Example

42
am
P
CR
909
248
53
218
−77
Example

43
an
Q
CR
838
55
74
326
−15
Example

44
ao
Q
CR
837
47
54
225
−107
Example

45
ap
R
CR
820
69
88
302
−61
Example

46
aq
R
CR
856
44
77
221
−105
Example

47
ar
S
CR
888
65
53
304
−47
Example

48
as
S
CR
902
35
57
330
−35
Example

49
at
S
CR
879
55
85
249
−71
Example

50
au
T
CR
852
47
54
250
−58
Example

51
av
T
CR
844
59
71
246
−80
Example

52
aw
U
CR
812
114
57
246
−80
Example

53
ax
U
CR
837
202
55
260
−77
Example

54
ay
V
CR
873
178
61
240
−43
Example

55
az
V
CR
858
155
78
238
−66
Example

56

ba

W

CR
842
46
56
334
−32
Comparative

Example

57

bb

X

CR
830
65
58
168
−40
Comparative

Example

58

bc

Y

CR
825
81
87
258
−80
Comparative

Example

59

bd

Z

CR
870
54
85
222
−19
Comparative

Example

TABLE 7

Reheating Process

Average Rate of

Maintaining Time
Temperature

Maintaining Time
Increase in
Reheating
Reheating
Total Maintaining

in Martensite
Bainite
Stop
Stop
Time in Bainite

Transformation
Transformation
Temperature
Temperature -
Transformation

Experiment
Temperature Range
Temperature Range
(T3)
Bs
Temperature Range

Example
Second
° C./second
° C.
° C.
Second

1
8
18
489
10
12
Example

2
9
20
427
−30
11
Example

3
12
12
471
−12
15
Example

4
9
25
443
−20
6
Example

5
10
24
420
−51
5
Example

6
12
15
470
−2
10
Example

7
7
22
485
9
8
Example

8
7
24
427
−43
6
Example

9
6
20
409
−63
6
Comparative

Example

10
12
20
483
−50
10
Example

11
8
22
484
−44
10
Example

12
5
14
455
−40
13
Example

13
15
15
447
−48
11
Example

14
7
27
438
−53
8
Comparative

Example

15
5
22
475
−32
12
Example

16
6
26
467
−52
9
Example

17
9
25
507
−36
9
Comparative

Example

18
8
26
577
−11
13
Example

19
4
15
538
−53
16
Example

20
9
26
495
−15
8
Example

21
6
11
446
−59
12
Example

22
12
17
464
−61
8
Example

23
7
15
505
−2
13
Example

24
11
22
522
3
9
Example

25

0

17
447
−1
13
Comparative

Example

26
8
18
487
−14
8
Example

27
6
11
455
−45
9
Example

28
11
27
485
−31
10
Example

29
11
15
494
−42
13
Example

30
15
25
485
−26
10
Comparative

Example

TABLE 8

Maintaining
Reheating Process

Time
Average Rate of

Maintaining
Temperature

Time in
Increase in

Total Maintaining

Martensite
Bainite
Reheating
Reheating
Time in Bainite

Transformation
Transformation
Stop
Stop
Transformation

Temperature
Temperature
Temperature
Temperature -
Temperature

Experiment
Range
Range
(T3)
Bs
Range

Example
Second
° C./second
° C.
° C.
Second

31
3
28
467
26
6
Example

32
8
16
380
−56
6
Example

33
6
25
492
20
7
Example

34
11
21
483
7
8
Example

35
5
18
539
−6
12
Example

36
14
23
577
14
8
Comparative

Example

37
6
25
564
10
9
Example

38
10
25
428
−29
7
Example

39
9
23
467
5

161
Comparative

Example

40
12
15
450
−13
11
Example

41
10
16
546
−19
22
Example

42
6
14
518
−61
21
Example

43
13
14
437
−39
9
Example

44
8
12
479
8
14
Example

45
4
17
529
9
11
Example

46
11
20
453
−45
9
Example

47
5
25
581
−10
14
Example

48
7
22
593
−6
14
Example

49
7
11
530
−41
22
Example

50
9
26
401
−62
6
Example

51
5
16
431
−43
9
Example

52
10
23
515
−26
12
Example

53
9
27
509
−40
10
Example

54
6
18
437
−38
12
Example

55
7
15
468
−20
13
Example

56
7
23
513
3
9
Comparative

Example

57
5
19
460
2
17
Comparative

Example

58
9
27
512
−39
9
Comparative

Example

59
10
18
584
7
23
Comparative

Example

TABLE 9

Third Cooling Process

Maintaining Time in

Bainite Trans-
Martensite Trans-

Bainite Transformation
Fourth Cooling Process
formation Start
formation Start

Experiment
Temperature Range
Average Cooling Rate
Temperature (Bs)
Temperature (Ms)

Example
Second
° C./second
° C.
° C.

1
407
7
479
309
Example

2
179
7
457
274
Example

3
212
13
483
317
Example

4
304
5
463
335
Example

5
271
13
471
348
Example

6
409
9
472
349
Example

7
407
4
476
334
Example

8
339
5
470
324
Example

9
9
10
472
328
Comparative

Example

10
347
7
533
329
Example

11
331
8
528
320
Example

12
264
9
495
312
Example

13
370
4
495
312
Example

14
186
13
491
305
Comparative

Example

15
159
13
507
294
Example

16
329
11
519
313
Example

17
350
9
543
353
Comparative

Example

18
149
7
588
339
Example

19
285
7
591
342
Example

20
305
8
510
346
Example

21
209
13
505
338
Example

22
149
4
525
371
Example

23
374
10
507
349
Example

24
237
9
519
368
Example

25
295
12
448
253
Comparative

Example

26
244
13
501
374
Example

27
276
11
500
373
Example

28
248
5
516
299
Example

29
384
4
536
330
Example

30
139
11
511
290
Comparative

Example

TABLE 10

Third Cooling Process

Maintaining Time in

Bainite Trans-
Martensite Trans-

Bainite Transformation
Fourth Cooling Process
formation Start
formation Start

Experiment
Temperature Range
Average Cooling Rate
Temperature (Bs)
Temperature (Ms)

Example
Second
° C./second
° C.
° C.

31
201
8
441
321
Example

32
430
7
436
313
Example

33
194
10
472
344
Example

34
194
6
476
351
Example

35
408
9
545
382
Example

36
338
8
563
411
Comparative

Example

37
349
12
554
396
Example

38
171
10
457
299
Example

39
283
11
462
307
Comparative

Example

40
202
7
463
309
Example

41
324
6
565
272
Example

42
348
7
579
295
Example

43
310
6
476
341
Example

44
195
12
471
332
Example

45
172
13
520
363
Example

46
405
4
498
326
Example

47
273
10
591
351
Example

48
418
10
599
365
Example

49
164
4
571
320
Example

50
149
5
463
308
Example

51
174
8
474
326
Example

52
288
13
541
326
Example

53
327
11
549
338
Example

54
374
8
475
283
Example

55
218
5
488
304
Example

56
332
4
510
366
Comparative

Example

57
416
13
458
208
Comparative

Example

58
229
4
551
338
Comparative

Example

59
412
6
577
241
Comparative

Example

TABLE 11

First Cooling
Second Cooling Process

Process
Average

Maintaining
Cooling Rate in

Maximum
Time in Ferrite
Bainite
Cooling
Cooling

Heating
Transformation
Transformation
Termination
Termination

Temperature
Temperature
Temperature
Temperature
Temperature -

Experiment
Cold-rolled
Chemical
Type of
(T1)
Range
Range
(T2)
Ms

Example
Steel Sheet
Constituent
Steel
° C.
Second
° C./second
° C.
° C.

60
g
C
EG
831
49
74
270
−64

61
z
K
EG
860
84
62
261
−38

62
ab
L
EG
809
46
88
274
−47

63
ay
V
EG
873
78
61
240
−43

64
a
A
GI
835
56
51
291
−49

65
d
B
GI
840
82
72
301
−71

66
i
D
GI
822
50
57
266
−30

67
ag
N
GI
864
59
54
312
−93

68
al
P
GI
912
47
51
284
−55

69
b
A
GA
842
61
23
284
−50

70
e
B
GA
832
71
19
322
−44

71
n
F
GA
825
49
22
249
−84

72
w
I
GA
888
54
27
328
−49

73
x
J
GA
868
53
17
332
−46

74
c
A
GI
829
48
55
273
−71

75
r
H
GI
852
80
64
304
−65

76
p
G
GI
802
76
79
281
−51

77
u
I
GI
915
56
49
297
−74

78
h
C
GA
837
43
12
278
−81

79
k
E
GA
812
56
25
287
−57

80
s
H
GA
842
51
19
312
−56

81
ad
M
GA
836
52
17
278
−98

82
aj
O
GA
847
66
17
263
−70

Maintaining
Reheating Process

Process
Average Rate of

Total

Maintaining
Temperature

Maintaining

Time in
Increase in

Time in

Martensite
Bainite
Reheating
Reheating
Bainite

Transformation
Transformation
Stop
Stop
Transformation

Temperature
Temperature
Temperature
Temperature -
Temperature

Experiment
Range
Range
(T3)
Bs
Range

Example
Second
° C./second
° C.
° C.
Second

60
7
22
485
9
8
Example

61
11
27
485
−31
10
Example

62
3
28
467
26
6
Example

63
6
18
437
−38
12
Example

64
10
11
486
−12
16
Example

65
7
19
471
−15
13
Example

66
10
14
497
−16
18
Example

67
9
13
527
−32
12
Example

68
8
22
548
−58
15
Example

69
4
14
524
30
18
Example

70
3
12
492
10
16
Example

71
4
20
501
−30
17
Example

72
5
10
507
−17
18
Example

73
5
19
531
28
14
Example

74
10
25
467
−33
8
Example

75
11
29
483
−41
6
Example

76
9
28
542
−42
11
Example

77
9
18
521
0
11
Example

78
4
22
483
−8
17
Example

79
4
19
490
−25
14
Example

80
3
16
494
−29
16
Example

81
6
24
507
16
12
Example

82
5
20
501
24
16
Example

TABLE 12

Third

Cooling

Process
Fourth

Maintaining
Cooling

Martensite

Time in Bainite
Process

Transformation

Alloying Conditions

Transformation
Average
Bainite
Start

Alloying

Temperature
Cooling
Transformation
Temperature

Temperature
Maintaining

Experiment
Range
Rate
Start Rate (Bs)
(Ms)
Plating Bath
(Tg)
Time

Example
Second
° C./second
° C.
° C.
Position
° C.
Second

60
407
4
476
334
After
—
—
Example

Annealing

61
248
5
516
299
After
—
—
Example

Annealing

62
201
8
441
321
After
—
—
Example

Annealing

63
374
8
475
283
After
—
—
Example

Annealing

64
157
9
498
340
Reheating
—
—
Example

Process

65
136
4
486
372
Reheating
—
—
Example

Process

66
179
10
513
296
Reheating
—
—
Example

Process

67
103
8
559
405
Reheating
—
—
Example

Process

68
147
7
606
339
Reheating
—
—
Example

Process

69
59
7
494
334
Reheating
—
10
Example

Process

70
50
6
482
366
Reheating
—
10
Example

Process

71
67
6
531
333
Reheating
—
10
Example

Process

72
240
6
524
377
Reheating
—
10
Example

Process

73
267
6
503
378
Reheating
—
10
Example

Process

74
300
11
500
344
Third Cooling
—
—
Example

Process

75
278
4
524
369
Third Cooling
—
—
Example

Process

76
85
6
584
332
Third Cooling
—
—
Example

Process

77
62
5
521
371
Third Cooling
—
—
Example

Process

78
137
4
491
359
Third Cooling
504
7
Example

Process

79
51
4
515
344
Third Cooling
544
7
Example

Process

80
37
4
523
368
Third Cooling
508
7
Example

Process

81
86
4
491
376
Third Cooling
535
7
Example

Process

82
81
4
477
333
Third Cooling
532
7
Example

Process

TABLE 13

Micro Structure Observation Results

Average

Cold-

Volume Fraction
Crystal

Experiment
Rolled
Chemical
Type of
F
B
BF
B + BF
TM
M
Retained γ
Others
Grain

Example
Steel Sheet
Constituent
Steel
%
%
%
%
%
%
%
%
μm

1
a
A
CR
33
18
12
30
27
0
10
0
4.5
Example

2
b
A
CR
45
19
2
21
32
2
0
0
5.1
Example

3
c
A
CR
27
21
15
36
22
3
11
1
2.9
Example

4
d
B
CR
47
3
12
15
33
0
5
0
9.0
Example

5
e
B
CR
41
9
29
38
15
0
5
1
7.7
Example

6
f
B
CR
39
19
10
29
22
4
6
0
7.2
Example

7
g
C
CR
36
23
9
32
25
1
6
0
6.5
Example

8
h
C
CR
43
32
0
32
22
3
0
0
8.4
Example

9
h
C
CR
41
5
2
7
19
30
2
1
4.7
Comparative

Example

10
i
D
CR
14
16
26
42
27
3
14
0
3.8
Example

11
j
D
CR
20
24
19
43
23
0
14
0
3.3
Example

12
k
E
CR
40
0
12
12
35
I
10
2
3.3
Example

13
l
E
CR
41
8
31
39
13
0
7
0
2.6
Example

14
m
E
CR
43
20
11
31
19
2
5
0

21.7
Comparative

Example

15
n
F
CR
35
22
8
30
31
0
4
0
1.9
Example

16
o
F
CR
28
0
18
18
41
2
10
1
2.2
Example

17
o
F
CR
3
18
26
44
44
3
4
2
2.5
Comparative

Example

18
p
G
CR
14
31
5
36
45
1
3
1
1.2
Example

19
q
G
CR
16
27
16
43
31
1
8
1
8.0
Example

20
r
H
CR
40
4
19
23
25
0
11
1
5.6
Example

21
s
H
CR
42
10
24
34
14
3
7
0
4.7
Example

22
t
H
CR
16
1
33
34
41
0
9
0
2.0
Example

23
u
I
CR
46
0
24
24
24
0
6
0
8.1
Example

24
v
I
CR
30
3
18
21
40
0
7
2
8.7
Example

25
w
I
CR
75
1
5
6

0

18
1
0
6.9
Comparative

Example

26
x
J
CR
32
5
37
42
15
2
9
0
5.5
Example

27
y
J
CR
35
10
15
25
31
2
5
2
6.2
Example

28
z
K
CR
40
24
17
41
15
0
4
0
5.6
Example

29
aa
K
CR
23
22
16
38
26
3
9
1
3.1
Example

30
aa
K
CR
44
0
6
6
42
4
4
0
2.9
Comparative

Example

TABLE 14

Micro Structure Observation Results

Cold-

Average

Rolled

Volume Fraction
Crystal

Experiment
Steel
Chemical
Type of
F
B
BF
B + BF
TM
M
Retained γ
Others
Grain

Example
Sheet
Constituent
Steel
%
%
%
%
%
%
%
%
μm

31
ab
L
CR
21
21
23
44
24
2
8
1
3.9
Example

32
ac
L
CR
27
31
4
35
32
0
6
0
4.5
Example

33
ad
M
CR
47
0
17
17
23
5
7
1
6.1
Example

34
ae
M
CR
43
5
25
30
19
0
8
0
4.9
Example

35
af
N
CR
43
20
13
33
17
0
7
0
4.4
Example

36
ag
N
CR
0
0
8
8
84
3
5
0
1.3
Comparative

Example

37
ah
N
CR
29
5
16
21
42
1
6
1
9.2
Example

38
ai
O
CR
36
2
19
21
28
0
15
0
5.1
Example

39
aj
O
CR
35
14
37
51
0
1
13
0
5.8
Comparative

Example

40
ak
O
CR
32
14
25
39
17
4
8
0
2.8
Example

41
al
P
CR
45
3
21
24
23
3
5
0
4.7
Example

42
am
P
CR
41
4
15
19
31
1
7
1
5.0
Example

43
an
Q
CR
28
10
31
41
22
0
9
0
4.7
Example

44
ao
Q
CR
34
0
18
18
41
0
7
0
6.1
Example

45
ap
R
CR
19
20
17
37
32
2
10
0
5.5
Example

46
aq
R
CR
45
15
4
19
35
1
0
0
6.0
Example

47
ar
S
CR
30
22
18
40
22
0
7
1
3.8
Example

48
as
S
CR
21
5
15
20
19
38
2
0
1.1
Example

49
at
S
CR
43
13
13
26
24
2
5
0
5.7
Example

50
au
T
CR
38
7
22
29
22
0
11
0
3.9
Example

51
av
T
CR
29
26
0
26
36
5
4
0
3.5
Example

52
aw
U
CR
25
12
10
22
38
3
10
2
7.0
Example

53
ax
U
CR
17
18
8
26
42
1
14
0
6.6
Example

54
ay
V
CR
35
6
23
29
17
2
17
0
4.7
Example

55
az
V
CR
26
14
18
32
28
1
13
0
6.3
Example

56
ba

W

CR
83
4
8
12
0
0
0
5
8.9
Comparative

Example

57
bb

X

CR
2

45

20

65

23
0
4
6
0.8
Comparative

Example

58
bc

Y

CR
35
28
0
28
35
2
0
0
8.4
Comparative

Example

59
bd

Z

CR

65

27
5
32
0
2
1
0
7.6
Comparative

Example

TABLE 15

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment
H2
H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

1
125
482
3.86
−0.61
17
7
3.12
2.09
1.03
1131
22
49
Example

2
119
513
4.31
−0.99
19
7
2.75
1.98
0.77
1116
24
66
Example

3
131
493
3.77
−0.49
22
3
3.12
1.99
1.13
1171
21
46
Example

4
120
427
3.56
−0.84
17
7
3.01
2.50
0.51
943
24
78
Example

5
124
408
3.30
−0.88
24
5
3.18
2.01
1.17
973
21
70
Example

6
117
394
3.37
−0.48
22
6
3.23
2.25
0.98
925
24
53
Example

7
113
377
3.35
−0.56
19
6
3.52
2.59
0.93
957
23
62
Example

8
121
409
3.37
−0.63
22
5
3.78
2.33
1.45
1022
22
68
Example

9
119
421
3.54

−0.30

19

0

3.67
2.39
1.28
1032
22

19

Comparative

Example

10
102
404
3.96
−0.43
18
4
2.45
1.96
0.49
1035
25
55
Example

11
112
411
3.67
−0.52
19
5
2.40
1.83
0.57
1010
22
67
Example

12
138
431
3.12
−0.45
22
4
2.77
1.75
1.02
1023
21
50
Example

13
128
429
3.36
−0.98
19
6
2.99
1.81
1.18
1012
21
88
Example

14
120
398
3.32
−1.03
23
3
2.83
1.56
1.27
963
23

22

Comparative

Example

15
157
456
2.90
−0.46
16
6
2.05
1.57
0.48
1303
15
42
Example

16
168
433
2.57
−0.61
21
4
2.16
1.63
0.53
1145
16
54
Example

17
295
408

1.38

−0.43
19
4
2.07
1.65
0.42
1250
9
44
Comparative

Example

18
131
351
2.68
−0.51
20
5
2.67
2.05
0.62
1140
16
59
Example

19
117
409
3.50
−0.78
23
4
2.67
2.13
0.54
1236
20
60
Example

20
148
405
2.74
−1.07
18
5
2.55
1.93
0.62
927
21
89
Example

21
150
429
2.86
−0.84
26
3
2.38
1.86
0.52
1047
19
65
Example

22
154
399
2.59
−0.45
20
4
2.99
1.80
1.19
1237
15
45
Example

23
142
458
3.23
−0.69
21
4
2.25
1.60
0.65
1052
19
73
Example

24
137
376
2.74
−0.58
19
7
2.31
1.60
0.71
1063
19
59
Example

25
134
523
3.91
0.11
37
0
2.22
1.67
0.55
920
25

10

Comparative

Example

26
135
435
3.22
−0.68
23
6
3.04
1.92
1.12
1029
20
74
Example

27
146
439
3.01
−0.76
18
5
2.74
2.15
0.59
1098
19
62
Example

28
101
427
4.22
−0.85
18
7
3.10
2.47
0.63
1194
22
68
Example

29
111
391
3.52
−0.73
22
4
3.22
2.52
0.70
1178
19
59
Example

30
119
417
3.50
−0.22
19
1
3.30
2.57
0.73
1222
19
8
Comparative

Example

TABLE 16

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment
H2
H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

31
115
402
3.50
−0.84
24
3
3.44
2.75
0.69
1068
22
58
Example

32
112
377
3.38
−0.66
17
7
3.74
2.37
1.37
1061
20
62
Example

33
140
434
3.11
−0.97
19
7
2.85
2.06
0.79
948
23
84
Example

34
148
403
2.72
−0.51
21
5
2.96
2.15
0.81
922
22
66
Example

35
134
409
3.06
−0.60
19
4
2.47
1.63
0.84
914
23
64
Example

36
241
330

1.37

0.07
18
4
2.34
1.73
0.61
970
6
58
Comparative

Example

37
116
398
3.42
−0.49
23
4
2.33
1.84
0.49
996
23
60
Example

38
145
434
2.99
−1.01
21
3
3.06
2.37
0.69
990
22
70
Example

39
148
341
2.30
−0.46
24
5
3.11
2.46
0.65
865
21
59
Comparative

Example

40
165
389
2.35
−0.84
18
6
3.76
2.14
1.62
1114
16
61
Example

41
143
453
3.16
−0.74
25
3
3.67
2.45
1.22
1038
21
71
Example

42
140
388
2.78
−1.08
26
5
3.52
2.64
0.88
923
22
80
Example

43
128
378
2.97
−0.93
19
6
2.45
1.91
0.54
945
23
77
Example

44
121
387
3.21
−0.80
23
4
2.68
1.80
0.88
1000
21
76
Example

45
132
333
2.53
−0.71
22
4
1.93
1.16
0.77
1025
20
74
Example

46
121
371
3.08
−0.78
23
3
1.89
1.38
0.51
1014
19
53
Example

47
142
347
2.44
−0.64
18
7
2.66
2.13
0.53
953
19
60
Example

48
159
541
3.40
−0.53
34
3
3.02
2.06
0.96
1359
15
34
Example

49
143
421
2.94
−0.44
20
4
2.79
2.01
0.78
1021
21
56
Example

50
169
437
2.58
−0.63
16
7
2.20
1.50
0.70
1047
20
61
Example

51
158
445
2.81
−0.67
19
6
2.22
1.53
0.69
1338
14
48
Example

52
141
372
2.64
−1.07
21
4
3.07
1.94
1.13
993
19
70
Example

53
137
405
2.97
−0.62
17
7
3.52
1.96
1.56
1347
17
49
Example

54
152
410
2.70
−1.12
20
5
1.92
1.45
0.47
1147
19
69
Example

55
141
403
2.86
−0.63
20
3
1.98
1.34
0.64
990
21
58
Example

56
116
142

1.22

0.24
25
5
2.30
2.06

0.24

414
35
80
Comparative

Example

57
339
454

1.34

−0.30

22
0
2.47
1.38
1.09
1409
7

26

Comparative

Example

58
86
245
2.85
−0.59
19
6
2.72
1.82
0.90
795
22
55
Comparative

Example

59
143
203

1.42

−0.35

32
3
0.66
0.48

0.18

723
24
41
Comparative

Example

TABLE 17

Micro Structure Observation Results

Average

Volume Fraction
Crystal

Experiment
Cold-rolled
Chemical
Type of
F
B
BF
B + BF
TM
M
Retained γ
Others
Grain

Example
Steel Sheet
Constituent
Steel
%
%
%
%
%
%
%
%
μm

60
g
C
EG
39
21
14
35
19
0
7
0
6.0
Example

61
z
K
EG
35
20
22
42
18
0
4
1
6.0
Example

62
ab
L
EG
23
22
20
42
23
0
12
0
4.5
Example

63
ay
V
EG
33
7
22
29
20
2
15
1
4.8
Example

64
a
A
GI
38
22
10
32
20
1
8
1
4.5
Example

65
d
B
GI
43
8
11
19
30
0
7
1
8.4
Example

66
i
D
GI
20
10
30
40
30
0
10
0
4.5
Example

67
ag
N
GI
43
20
13
33
17
0
6
1
6.1
Example

68
al
P
GI
38
10
19
29
26
1
6
0
5.7
Example

69
b
A
GA
45
10
16
26
27
0
2
0
6.2
Example

70
e
B
GA
47
15
20
35
13
0
5
0
5.9
Example

71
n
F
GA
38
11
15
26
28
2
6
0
3.9
Example

72
w
I
GA
40
8
20
28
26
0
6
0
7.1
Example

73
x
J
GA
29
15
28
43
21
0
7
0
5.0
Example

74
c
A
GI
32
19
6
25
29
0
13
1
4.3
Example

75
r
H
GI
37
0
19
19
33
1
10
0
7.2
Example

76
p
G
GI
19
18
19
37
35
0
9
0
8.6
Example

77
u
I
GI
45
0
28
28
22
0
5
0
7.4
Example

78
h
C
GA
39
22
12
34
24
3
0
0
9.0
Example

79
k
E
GA
38
2
21
23
28
0
8
3
5.1
Example

80
s
H
GA
38
13
20
33
19
1
9
0
6.1
Example

81
ad
M
GA
41
2
11
13
37
1
8
0
6.7
Example

82
aj
O
GA
33
18
15
33
19
0
15
0
6.3
Example

TABLE 18

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment

H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
H2 Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

60
113
403
3.57
−0.63
17
5
3.35
2.42
0.93
940
25
77
Example

61
111
486
4.37
−0.63
18
6
3.05
2.54
0.51
1184
19
63
Example

62
95
458
4.82
−0.79
22
3
3.26
2.74
0.52
1070
22
60
Example

63
131
450
3.44
−0.58
18
6
2.02
1.44
0.58
1139
19
48
Example

64
132
467
3.54
−0.71
19
4
2.95
1.75
1.20
1101
21
51
Example

65
106
477
4.50
−0.71
18
5
2.97
2.53
0.44
923
28
76
Example

66
126
393
3.12
−0.82
17
6
2.37
1.91
0.46
1005
21
78
Example

67
115
467
4.06
−0.44
18
3
2.40
1.76
0.64
960
22
55
Example

68
135
448
3.32
−0.60
19
4
3.97
2.55
1.42
1027
19
74
Example

69
109
497
4.56
−0.68
21
3
2.88
1.87
1.01
1113
24
66
Example

70
141
466
3.31
−0.91
19
7
3.38
2.33
1.05
961
21
72
Example

71
142
448
3.15
−0.47
18
4
2.12
1.64
0.48
1261
16
36
Example

72
143
606
4.23
−0.72
20
3
2.30
1.77
0.53
937
23
85
Example

73
120
496
4.14
−0.98
18
6
3.18
2.19
0.99
1024
24
74
Example

74
131
487
3.71
−0.97
17
5
3.59
1.96
1.63
1208
20
60
Example

75
147
479
3.26
−0.45
20
3
2.50
1.90
0.60
909
23
60
Example

76
122
458
3.75
−1.03
19
5
2.68
2.24
0.44
1237
18
69
Example

77
129
506
3.92
−0.93
16
7
2.13
1.76
0.37
1042
20
84
Example

78
121
442
3.65
−0.65
19
3
4.05
2.23
1.82
1039
20
62
Example

79
118
487
4.13
−0.68
18
6
2.69
1.62
1.07
1003
23
81
Example

80
138
499
3.61
−0.74
21
3
2.39
1.92
0.47
1048
20
63
Example

81
143
515
3.60
−0.80
17
5
3.11
2.25
0.86
941
23
70
Example

82
129
462
3.58
−0.71
20
6
3.17
2.35
0.82
929
22
81
Example

TABLE 19

Experiment
C
Si
Mn
P
S
Al
N
O

Example
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %

AA
0.112
0.78
1.99
0.028
0.0022
0.054
0.0022
0.0020
Example

AB
0.193
1.26
2.52
0.015
0.0036
0.012
0.0025
0.0037
Example

AC
0.087
1.06
2.60
0.003
0.0033
0.050
0.0041
0.0014
Example

AD
0.144
1.75
1.93
0.018
0.0038
0.015
0.0054
0.0023
Example

AE
0.205
0.99
2.28
0.014
0.0021
0.114
0.0044
0.0018
Example

AF
0.235
0.75
1.75
0.014
0.0005
0.023
0.0007
0.0031
Example

AG
0.310
0.57
1.94
0.006
0.0035
0.341
0.0055
0.0021
Example

AH
0.187
1.39
2.34
0.023
0.0015
0.050
0.0045
0.0016
Example

AI
0.159
1.73
1.97
0.014
0.0006
0.031
0.0055
0.0025
Example

AJ
0.098
1.92
2.78
0.009
0.0039
0.056
0.0030
0.0023
Example

AK
0.237
1.34
1.46
0.015
0.0015
0.045
0.0050
0.0015
Example

AL
0.172
0.36
2.38
0.009
0.0010
1.054
0.0016
0.0019
Example

AM
0.130
0.84
2.20
0.010
0.0013
0.012
0.0053
0.0023
Example

AN
0.275
1.60
1.96
0.013
0.0032
0.025
0.0010
0.0019
Example

AO
0.193
1.17
1.84
0.021
0.0090
0.021
0.0019
0.0019
Example

AP
0.257
0.73
1.31
0.011
0.0049
0.050
0.0053
0.0022
Example

AQ
0.205
0.17
2.58
0.004
0.0002
1.719
0.0044
0.0023
Example

TABLE 20

Experiment
Ti
Nb
B
Cr
Ni
Cu
Mo
V
Ca
Ce
Mg
REM

Example
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %
mass %

AA

Example

AB

Example

AC
0.031

Example

AD
0.053

Example

AE

0.028

Example

AF

Example

AG

0.14

Example

AH

0.0041

Example

AI

0.0022

Example

AJ

0.32

Example

AK

0.93

Example

AL

0.23

Example

AM

Example

AN

Example

AO

Example

AP
0.009

1.23
0.12

Example

AQ

0.0027

Example

TABLE 21

Slab
Ar₃
Finish
Cooling

Heating
Transformation
Rolling
Rate After
Winding

Experiment
Chemical
Temperature
Point
Temperature
Rolling
Temperature

Example
Constituent
° C.
° C.
° C.
° C./second
° C.

ca
AA
1245
707
941
23
627

cb
AA
1250
707
931
33
684

cc
AA
1205
707
892
6
654

cd
AA
1210
707
901
36
607

ce
AB
1225
648
882
26
617

cf
AB
1185
648
940
37
636

cg
AB
1230
648
894
36
466

ch
AB
1185
648
896
27
628

ci
AC
1180
669
927
35
684

cj
AC
1250
669
943
29
645

ck
AC
1240
669
883
36
615

cl
AC
1205
669
876
31
641

cm
AD
1205
734
914
37
620

cn
AD
1195
734
903
48
718

co
AE
1235
657
892
27
673

cp
AE
1235
657
971
39
644

cq
AF
1250
688
917
31
614

cr
AF
1215
688
900
35
620

cs
AF
1185
688
925
32
644

ct
AF
1205
688
920
3
637

cu
AG
1235
634
890
37
653

cv
AG
1215
634
926
49
614

cw
AH
1250
671
920
28
660

cx
AH
1250
671
937
29
638

cy
AI
1225
725
919
48
674

cz
AI
1235
725
898
51
640

Cooling
Volume

Cold-rolled

Rate After
Fraction of

Rolling
Sheet

Experiment
Left Side of
Winding
Austenite
Bs
Reduction
Thickness

Example
Equation (1)
° C./hour
volume %
° C.
%
mm

ca
11.7
13
71
570
50
1.4
Example

cb
50.3
12
81
576
50
1.4
Example

cc
19.3
14

23

475
50
1.4
Comparative

Example

cd
7.1
13
80
575
50
1.4
Example

ce
10.3
11
84
481
60
1.2
Example

cf
15.2
13
85
482
60
1.2
Example

cg
0.1
10
70
467
60
1.2
Comparative

Example

ch
14.0
11
86
482
60
1.2
Example

ci
41.1
14
92
523
50
1.2
Example

cj
18.6
13
84
520
50
1.2
Example

ck
4.1

28

78
518
50
1.2
Comparative

Example

cl
17.0
13
79
518
50
1.2
Example

cm
14.2
9
81
531
60
1.0
Example

cn
74.0
15
82
532
60
1.0
Example

co
39.4
11
92
522
50
2.0
Example

cp
25.2
10
100
527
50
2.0
Example

cq
9.7
12
76
547
60
1.2
Example

cr
10.7
13
90
561
60
1.2
Example

cs
26.9
8
80
551
60
1.2
Example

ct
14.8
14

11

17
60
1.2
Comparative

Example

cu
29.0
10
84
541
50
1.6
Example

cv
14.1
8
87
545
50
1.6
Example

cw
32.0
10
100
507
45
1.1
Example

cx
25.2
8
92
502
45
1.1
Example

cy
42.6
11
82
525
50
1.6
Example

cz
13.8
15
81
524
50
1.6
Example

TABLE 22

Slab
Ar₃
Finish
Cooling

Heating
Transformation
Rolling
Rate After
Winding

Experiment
Chemical
Temperature
Point
Temperature
Rolling
Temperature

Example
Constituent
° C.
° C.
° C.
° C./second
° C.

da
AJ
1190
662
898
24
642

db
AJ
1200
662
966
26
653

dc
AK
1240
691
949
46
618

dd
AK
1245
691
910
53
605

de
AL
1225
627
890
51
667

df
AL
1215
627
922
45
620

dg
AM
1205
684
897
40
679

dh
AM
1230
684
943
40
703

di
AM
1245
684
919
42
677

dj
AN
1245
684
885
29
670

dk
AN
1200
684
914
35
615

dl
AN
1240
684
924
33
672

dm
AO
1215
708
886
25
664

dn
AO
1250
708
928
32
734

do
AO
1230
708
935
42
685

dp
AP
1220
659
892
32
630

dq
AP
1245
659
902
36
648

dr
AQ
1240
599
911
25
635

ds
AQ
1235
599
927
36
604

Cold-

Volume

rolled

Left Side of
Cooling Rate
Fraction of

Rolling
Sheet

Experiment
Equation
After Winding
Austenite
Bs
Reduction
Thickness

Example
(1)
° C./hour
volume %
° C.
%
mm

da
17.1
12
100
453
38
1.6
Example

db
21.3
14
89
450
50
1.2
Example

dc
16.0
8
94
523
50
1.2
Example

dd
5.9
15
86
516
38
1.6
Example

de
44.1
9
85
608
50
1.2
Example

df
8.6
15
79
603
50
1.2
Example

dg
38.7
14
91
550
43
1.2
Example

dh
78.6
11
80
545
43
1.2
Example

di
46.3
10
88
549
43
1.2
Example

dj
29.3
14
80
486
50
1.2
Example

dk
12.8
9
83
490
50
1.2
Example

dl
47.0
10
87
494
50
1.2
Example

dm
29.4
13
83
545
43
1.2
Example

dn
81.8
19
100
557
43
1.2
Example

do
58.4
10
92
552
43
1.2
Example

dp
16.5
10
83
527
50
1.6
Example

dq
26.9
9
90
534
50
1.6
Example

dr
17.3
11
75
623
50
1.6
Example

ds
6.1
14
67
613
50
1.6
Example

TABLE 23

First Cooling

Process
Second Cooling Process

Maintaining
Average Cooling

Maximum
Time in Ferrite
Rate in Bainite
Cooling
Cooling

Heating
Transformation
Transformation
Termination
Termination

Temperature
Temperature
Temperature
Temperature
Temperature -

Experiment
Cold-rolled
Chemical
Type of
(T1)
Range
Range
(T2)
Ms

Example
Steel Sheet
Constituent
Steel
° C.
second
° C./second
° C.
° C.

83
ca
AA
CR
786
27
118
355
−45
Example

84
cb
AA
CR
793
61
46
332
−49
Example

85

cc

AA
CR
787
33
79
286
−104
Comparative

Example

86
cd
AA
CR
795
30
57
385
−5
Comparative

Example

87
ce
AB
CR
816
64
19
231
−58
Example

88
cf
AB
CR
790
102
56
209
−44
Example

89

cg

AB
CR
823
67
59
263
−32
Comparative

Example

90
ch
AB
CR
782
35
50
273
−45
Comparative

Example

91
ci
AC
CR
778
46
34
351
−33
Example

92
cj
AC
CR
840
72
61
360
−23
Example

93

ck

AC
CR
845
82
60
267
−56
Comparative

Examplele

94
cl
AC
CR
801
40
59
344
−35
Comparative

Example

95
cm
AD
CR
776
93
52
310
−38
Example

96
cn
AD
CR
784
47
37
307
−54
Example

97
co
AE
CR
854
156
67
253
−43
Example

98
cp
AE
CR
800
79
33
230
−72
Example

99
cq
AF
CR
827
79
53
294
−33
Example

100
cr
AF
CR
778
80
28
214
−66
Example

101
cs
AF
CR
800
61
58
248
−45
Comparative

Example

102

ct

AF
CR
858
54
58
302
−26
Comparative

Example

103
cu
AG
CR
774
58
38
130
−30
Example

104
cv
AG
CR
819
41
50
264
−35
Example

105
cw
AH
CR
834
85
82
277
−41
Example

106
cx
AH
CR
800
203
65
239
−51
Example

107
cy
AI
CR
818
75
53
302
−49
Example

108
cz
AI
CR
877
61
52
300
−47
Example

109
da
AJ
CR
852
349
23
279
−70
Example

110
db
AJ
CR
783
159
60
300
−46
Example

111
dc
AK
CR
762
84
18
229
−46
Example

112
dd
AK
CR
791
107
66
292
−21
Example

113
de
AL
CR
905
95
75
340
−24
Example

114
df
AL
CR
869
41
31
328
−35
Example

115
dg
AM
CR
783
129
106
278
−73
Example

116
dh
AM
CR
840
186
62
299
−39
Example

117
di
AM
CR

1052
47
37
343
−47
Comparative

Example

118
dj
AN
CR
814
67
57
231
−39
Example

119
dk
AN
CR
796
30
69
234
−53
Example

120
dl
AN
CR

703

35
24

340

484

Comparative

Example

121
dm
AO
CR
800
26
57
315
−37
Example

122
dn
AO
CR
855
66
46
311
−45
Example

123
do
AO
CR
830
130
28

380

93
Comparative

Example

124
dp
AP
CR
803
46
33
257
−31
Example

125
dq
AP
CR
821
86
64
253
−27
Example

126
dr
AQ
CR
785
115
33
277
−59
Example

127
ds
AQ
CR
851
264
56
249
−54
Example

TABLE 24

Maintaining
Reheating Process

Process
Average Rate of

Maintaining
Temperature

Time in
Increase in

Total Maintaining

Martensite
Bainite

Time in Bainite

Transformation
Transformation
Reheating Stop
Reheating Stop
Transformation

Temperature
Temperature
Temperature
Temperature -
Temperature

Experiment
Range
Range
(T3)
Bs
Range

Example
Second
° C./second
° C.
° C.
Second

83
16
25
544
−35
7
Example

84
33
21
511
−56
13
Example

85
15
31
537
−37
8
Comparative

Example

86
1
31
532
−40
9
Comparative

Example

87
16
25
425
−30
15
Example

88
28
55
478
28
7
Example

89
31
20
448
−19
12
Comparative

Example

90
19
15

349

−129
14
Comparative

Example

91
24
43
493
−26
7
Example

92
25
21
528
10
9
Example

93
16
37
459
−23
7
Comparative

Example

94
23
3
467
−50

43

Comparative

Example

95
22
29
493
−27
9
Example

96
26
32
482
−48
10
Example

97
20
18
504
10
14
Example

98
63
26
451
−50
14
Example

99
27
23
534
−11
14
Example

100
11
18
514
−5
22
Example

101

2031
24
491
−24
13
Comparative

Example

102
26
25
493
−58
13
Comparative

Example

103
34
17
457
−17
27
Example

104
42
77
470
−72
8
Example

105
29
25
488
2
9
Example

106
30
45
418
−52
7
Example

107
21
30
509
−14
9
Example

108
8
37
526
6
8
Example

109
52
36
378
−67
7
Example

110
24
19
442
0
7
Example

111
21
31
419
−67
18
Example

112
21
29
476
−32
10
Example

113
29
24
573
−24
13
Example

114
24
18
509
−89
20
Example

115
41
50
540
18
5
Example

116
26
39
482
−49
8
Example

117
19
14
572
18
16
Comparative

Example

118
76
41
437
−38
9
Example

119
34
29
498
8
10
Example

120
0
32
471
193
0
Comparative

Example

121
23
14
500
−47
17
Example

122
8
46
520
−28
8
Example

123
0
30
478
−28
39
Comparative

Example

124
31
26
487
−30
16
Example

125
23
30
465
−41
11
Example

126
21
41
544
−71
15
Example

127
9
20
533
−51
19
Example

TABLE 25

Third Cooling Process
Fourth Cooling

Maintaining Time in
Process
Bainite
Martensite

Bainite
Average Cooling
Transformation Start
Transformation Start

Experiment
Transformation
Rate
Temperature (Bs)
Temperature (Ms)

Example
Second
° C./second
° C.
° C.

83
135
3
579
400
Example

84
149
9
567
381
Example

85
236
4
574
390
Comparative

Example

86
130
11
572
390
Comparative

Example

87
461
9
455
289
Example

88
524
8
450
253
Example

89
411
12
467
295
Comparative

Example

90
590
4
478
318
Comparative

Example

91
403
7
519
384
Example

92
65
5
518
383
Example

93
577
13
482
323
Comparative

Example

94
558
6
517
379
Comparative

Example

95
193
6
520
348
Example

96
232
4
530
361
Example

97
130
11
494
296
Example

98
218
12
501
302
Example

99
173
4
545
327
Example

100
295
5
519
280
Example

101
156
13
515
293
Comparative

Example

102
146
12
551
328
Comparative

Example

103
218
6
474
160
Example

104
275
9
542
299
Example

105
50
6
486
318
Example

106
171
9
470
290
Example

107
463
11
523
351
Example

108
484
4
520
347
Example

109
606
8
445
349
Example

110
535
7
442
346
Example

111
233
13
486
275
Example

112
264
13
508
313
Example

113
115
3
597
364
Example

114
241
8
598
363
Example

115
236
12
522
351
Example

116
92
7
531
338
Example

117
163
11
554
390
Comparative

Example

118
136
8
475
270
Example

119
152
10
490
287
Example

120
163

—

278

−144
Comparative

Example

121
164
9
547
352
Example

122
75
6
548
356
Example

123
244
3
506
287
Comparative

Example

124
399
6
517
288
Example

125
382
11
506
280
Example

126
276
5
615
336
Example

127
205
9
584
303
Example

TABLE 26

Micro Structure Observation Results

Average

Volume Fraction
Crystal

Experiment
Cold-rolled
Chemical
Type of
F
B
BF
B + BF
TM
M
Retained γ
Others
Grain

Example
Steel Sheet
Constituent
Steel
%
%
%
%
%
%
%
%
μn

83
ca
AA
CR
12
19
24
43
40
0
3
2
3.5
Example

84
cb
AA
CR
31
26
14
40
27
0
2
0
5.5
Example

85
cc
AA
CR
23
18
2
20

56

1
0
0
4.8
Comparative

Example

86
cd
AA
CR
26

32

34

66
4
0
4
0
5.1
Comparative

Example

87
ce
AB
CR
36
10
13
23
30
3
7
1
4.9
Example

88
cf
AB
CR
45
24
8
32
19
0
3
1
6.1
Example

89
cg
AB
CR
33
27
5
32
19

13
3
0
4.8
Comparative

Example

90
ch
AB
CR
21
28
8
36
34
3
5
1
3.4
Comparative

Example

91
ci
AC
CR
19
13
31
44
28
2
4
3
4.2
Example

92
cj
AC
CR
25
37
6
43
31
0
0
1
5.1
Example

93
ck
AC
CR

68

7
8
15
15
0
1
1
7.1
Comparative

Example

94
cl
AC
CR
27
35
2
37
33
3
0
0
5.7
Comparative

Example

95
cm
AD
CR
31
33
0
33
32
2
0
2
4.2
Example

96
cn
AD
CR
22
27
13
40
34
0
3
1
4.0
Example

97
co
AE
CR
38
17
10
27
30
0
5
0
7.1
Example

98
cp
AE
CR
32
26
2
28
40
0
0
0
7.2
Example

99
cq
AF
CR
26
36
8
44
23
3
4
0
3.7
Example

100
cr
AF
CR
42
4
9
13
33
0
11
1
8.0
Example

101
cs
AF
CR
40
0
0
0
27
0
0

33
5.3
Comparative

Example

102
ct
AF
CR
12

33

36

69
17
1
2
0
3.1
Comparative

Example

103
cu
AG
CR
48
0
25
25
13
0
14
0
6.0
Example

104
cv
AG
CR
16
11
33
44
20
2
17
1
4.4
Example

105
cw
AH
CR
27
11
12
23
43
1
6
0
6.3
Example

106
cx
AH
CR
41
7
13
20
31
0
7
1
6.5
Example

107
cy
AI
CR
22
34
9
43
29
1
3
2
4.7
Example

108
cz
AI
CR
23
11
25
36
33
1
6
1
6.7
Example

109
da
AJ
CR
23
22
7
29
47
0
1
0
5.6
Example

110
db
AJ
CR
26
17
21
38
31
0
4
1
6.4
Example

111
dc
AK
CR
37
10
23
33
19
0
11
0
7.0
Example

112
dd
AK
CR
21
2
41
43
18
3
15
0
7.5
Example

113
de
AL
CR
25
42
0
42
30
0
2
1
5.4
Example

114
df
AL
CR
26
8
33
41
24
1
8
0
4.8
Example

115
dg
AM
CR
43
19
0
19
38
0
0
0
6.7
Example

116
dh
AM
CR
42
7
28
35
17
0
5
1
7.5
Example

117
di
AM
CR
1
26
18
44
45
4
4
2
1.3
Comparative

Example

118
dj
AN
CR
28
30
0
30
37
0
2
3
5.5
Example

119
dk
AN
CR
18
35
3
38
43
0
0
1
3.2
Example

120
dl
AN
CR

78

0
0
0
0
3
3

16

16.9
Comparative

Example

121
dm
AO
CR
15
0
44
44
29
3
8
1
6.7
Example

122
dn
AO
CR
12
9
33
42
37
0
9
0
4.4
Example

123
do
AO
CR
45
27
16
43
2
3
5
2
9.8
Comparative

Example

124
dp
AP
CR
28
12
30
42
19
1
9
1
4.4
Example

125
dq
AP
CR
32
5
36
41
15
0
11
1
6.8
Example

126
dr
AQ
CR
32
27
8
35
33
0
0
0
5.9
Example

127
ds
AQ
CR
45
5
16
21
23
1
10
0
6.1
Example

TABLE 27

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment
H2
H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

83
121
513
4.23
−0.89
18
7
2.42
1.53
0.89
952
23
67
Example

84
120
541
4.51
−0.60
21
3
2.49
1.46
1.03
1080
23
61
Example

85
117
524
4.50

−0.05

33
1
2.10
1.89

0.21

1144
11
23
Comparative

Example

86
123
542
4.39

−0.21

28

0

2.40
1.77
0.63
944
16
17
Comparative

Example

87
137
534
3.91
−0.57
16
5
3.18
2.16
1.02
1527
13
35
Example

88
128
459
3.58
−0.44
19
4
3.00
2.12
0.88
1349
15
48
Example

89
125
602
4.81

−0.34

23
3
2.71
2.44

0.27

1427
13
22
Comparative

Example

90
131
566
4.34

−0.30

17

0

3.15
2.04
1.11
1260
18
28
Comparative

Example

91
121
584
4.82
−0.91
17
7
3.24
1.88
1.36
1090
22
65
Example

92
136
372
2.74
−0.78
21
3
3.16
1.95
1.21
1085
16
66
Example

93
121
430
3.55
0.13
35

0

2.74
2.51

0.23

917
22
15
Comparative

Example

94
121
581
4.79

−0.26

19
0
3.09
2.25
0.84
1027
22
25
Comparative

Example

95
132
680
5.15
−0.56
18
5
2.52
1.60
0.92
1066
26
65
Example

96
139
721
5.20
−0.64
19
4
2.97
1.48
1.49
1091
24
57
Example

97
123
646
5.25
−0.58
23
5
3.27
1.59
1.68
1129
22
63
Example

98
129
484
3.76
−0.68
18
8
3.05
1.94
1.11
1403
15
52
Example

99
124
613
4.94
−0.72
20
4
2.03
1.49
0.54
1124
21
47
Example

100
111
438
3.94
−0.46
18
7
2.20
1.43
0.77
1376
16
37
Example

TABLE

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment
H2
H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

101
112
456
4.06

−0.14

27

0

2.37
1.51
0.86
1228
18
17
Comparative

Example

102
121
510
4.21

−0.29

30
1
1.86
1.68

0.18

1306
9

22

Comparative

Example

103
108
476
4.40
−0.44
23
4
2.69
1.31
1.38
1398
18
44
Example

104
114
465
4.08
−0.57
18
7
2.58
1.56
1.02
1532
15
42
Example

105
136
518
3.82
−0.65
16
5
2.98
1.76
1.22
1081
20
53
Example

106
131
655
5.00
−0.58
22
2
3.05
1.93
1.12
1135
23
48
Example

107
139
569
4.11
−0.86
18
7
2.83
1.41
1.42
1098
20
77
Example

108
140
725
5.17
−0.79
20
6
2.53
1.60
0.93
1404
18
48
Example

109
153
572
3.74
−0.63
18
7
3.65
2.36
1.29
1131
16
51
Example

110
153
773
5.04
−0.95
19
6
3.49
2.38
1.11
1250
21
64
Example

111
129
661
5.11
−0.45
21
2
1.90
1.23
0.67
1332
22
44
Example

112
130
491
3.77
−0.66
21
3
1.71
1.15
0.56
1450
15
35
Example

113
106
465
4.37
−0.59
17
4
3.50
1.80
1.70
1280
18
48
Example

114
112
515
4.59
−0.84
17
7
2.92
2.15
0.77
1237
19
59
Example

115
120
624
5.19
−0.45
22
5
2.84
1.69
1.15
1194
22
55
Example

116
115
422
3.66
−0.50
18
4
2.74
1.51
1.23
1011
20
55
Example

117
304
419

1.38

−0.32

23
3
2.86
1.76
1.10
1056
11
26
Comparative

Example

118
138
648
4.68
−0.61
20
3
2.44
1.43
1.01
1319
18
43
Example

119
136
491
3.61
−1.01
21
6
2.58
1.71
0.87
1455
14
49
Example

120
129
615
4.77
0.21
32

0

2.50
1.65
0.85
733
13
16
Comparative

Example

121
126
507
4.03
−0.46
23
2
2.59
1.39
1.20
1113
19
44
Example

122
125
459
3.66
−0.58
18
8
2.50
1.21
1.29
1311
15
52
Example

123
127
522
4.11

−0.24

29

0

2.36
1.33
1.03
1005
18
31
Comparative

Example

124
109
408
3.74
−0.62
19
8
1.78
1.11
0.67
1129
18
65
Example

125
112
552
4.95
−0.72
17
7
1.73
1.12
0.61
1380
18
57
Example

126
89
375
4.20
−0.57
18
6
3.29
2.13
1.16
1278
16
46
Example

127
95
517
5.42
−0.49
24
1
2.83
2.27
0.56
1351
20
36
Example

TABLE 29

Slab
Ar₃
Finish
Cooling

Heating
Transformation
Rolling
Rate After
Winding

Experiment
Chemical
Temperature
Point
Temperature
Rolling
Temperature

Example
Constituent
° C.
° C.
° C.
° C./second
° C.

dt
AA
1205
707
903
35
642

du
AA
1200
707
918
30
635

dv
AA
1220
707
897
31
628

dw
AB
1210
648
915
29
626

dx
AB
1215
648
907
36
618

dy
AC
1230
669
926
29
623

dz
AC
1235
669
890
31
646

Cold-

Volume

rolled

Cooling Rate
Fraction of

Rolling
Sheet

Experiment
Left Side of
After Winding
Austenite
Bs
Reduction
Thickness

Example
Equation (1)
° C./hour
% by volume
° C.
%
mm

dt
22.6
15
81
576
0
3.0
Example

du
19.7
12
83
574
0
3.0
Example

dv
20.4
13
88
572
0
3.0
Example

dw
18.9
13
85
482
0
2.3
Example

dx
15.9
14
86
483
0
2.3
Example

dy
17.9
13
77
518
0
4.0
Example

dz
28.2
15
86
521
0
4.0
Example

TABLE 30

First Cooling
Second Cooling Process

Process
Average

Maintaining
Cooling Rate

Maximum
Time in Ferrite
in Bainite
Cooling
Cooling

Heating
Transformation
Transformation
Termination
Termination

Temperature
Temperature
Temperature
Temperature
Temperature -

Experiment
Hot-rolled
Chemical
Type of
(T1)
Range
Range
(T2)
Ms

Example
Steel Sheet
Constituent
Steel
° C.
Second
° C./second
° C.
° C.

128
dt
AA
HR
838
32
58
339
−49

129
du
AA
HR
843
52
55
343
−29

130
dv
AA
HR-GA
837
38
44
332
−60

131
dw
AB
HR
873
49
52
249
−76

132
dx
AB
HR-GA
863
45
48
280
−39

133
dy
AC
HR
840
53
62
344
−28

134
dz
AC
HR-GI
822
46
50
320
−51

Maintaining

Process
Reheating Process

Maintaining
Average Rate of

Time in
Temperature

Total Maintaining

Martensite
Increase in Bainite
Reheating
Reheating
Time in Bainite

Transformation
Transformation
Stop
Stop
Transformation

Temperature
Temperature
Temperature
Temperature -
Temperature

Experiment
Range
Range
(T3)
Bs
Range

Example
Second
° C./second
° C.
° C.
Second

128
17
49
480
−90
5
Example

129
8
35
498
−65
7
Example

130
10
37
478
−96
6
Example

131
14
45
501
24
6
Example

132
10
40
493
20
7
Example

133
14
40
499
−12
5
Example

134
15
25
479
−31
7
Example

TABLE 31

Third Cooling

Process

Maintaining

Time in
Fourth

Bainite
Cooling

Alloying Conditions

Transformation
Process
Bainite
Martensite

Alloying

Temperature
Average
Transformation
Transformation Start

Temperature
Maintaining

Experiment
Range
Cooling Rate
Start Rate (Bs)
Temperature (Ms)
Plating Bath
(Tg)
Time

Example
Second
° C./second
° C.
° C.
Position
° C.
Second

128
432
12
570
388
—
—
—
Example

129
330
11
563
372
—
—
—
Example

130
350
10
574
392
After
505
25
Example

Annealing

131
252
11
477
325
—
—
—
Example

132
143
10
473
319
Reheating
493
21
Example

Process

133
338
8
511
372
—
—
—
Example

134
433
11
510
371
After
—
—
Example

Annealing

TABLE 32

Micro Structure Observation Results

Average

Volume Fraction
Crystal

Experiment
Hot-Rolled
Chemical
Type of
F
B
BF
B + BF
TM
M
Retained γ
Others
Grain

Example
Steel Sheet
Constituent
Steel
%
%
%
%
%
%
%
%
μm

128
dt
AA
HR
27
25
15
40
29
1
3
0
7.5
Example

129
du
AA
HR
38
13
28
41
16
0
5
0
8.7
Example

130
dv
AA
HR-GA
20
36
0
36
42
0
2
0
6.4
Example

131
dw
AB
HR
15
15
22
37
43
0
5
0
6.3
Example

132
dx
AB
HR-GA
19
37
6
43
33
2
3
0
5.7
Example

133
dy
AC
HR
33
28
12
40
25
1
0
1
6.3
Example

134
dz
AC
HR-GI
34
25
13
38
23
0
4
1
5.7
Example

TABLE 33

Mn Segregation

Difference

between
Material

Maximum
Quality

Hardness Measurement Results

Value and
Measurement

f
f
Maximum
Minimum
Minimum
Results

Experiment
H2
H98

(Maximum)
(Minimum)
Concentration
Concentration
Value
TS
EL
λ

Example
Hv
Hv
H98/H2
K*
%
%
mass %
mass %
mass %
MPa
%
%

128
108
441
4.08
−0.62
13
2
2.39
1.71
0.68
980
19
56
Example

129
103
442
4.29
−0.57
15
2
2.41
1.79
0.62
924
24
59
Example

130
105
412
3.92
−0.67
12
3
2.41
1.65
0.76
963
21
52
Example

131
115
510
4.43
−0.64
17
2
2.97
2.15
0.82
1418
13
34
Example

132
122
495
4.06
−0.58
13
4
3.00
2.16
0.84
1305
15
39
Example

133
101
396
3.92
−0.48
15
2
3.06
2.12
0.94
1019
18
44
Example

134
104
426
4.10
−0.66
15
2
2.98
2.18
0.80
1107
18
45
Example

Number	Date	Country	Kind
2010-208329	Sep 2010	JP	national
2010-208330	Sep 2010	JP	national

Number	Name	Date	Kind
20030196735	Sugiura et al.	Oct 2003	A1
20110008647	Azuma et al.	Jan 2011	A1
20110024004	Azuma et al.	Feb 2011	A1
20110139305	Bettacchini	Jun 2011	A1

Number	Date	Country
1343262	Apr 2002	CN
101646794	Feb 2010	CN
1607489	Dec 2005	EP
59-219473	Dec 1984	JP
2001-192768	Jul 2001	JP
2004-068050	Mar 2004	JP
2004-263270	Sep 2004	JP
2007-009317	Jan 2007	JP
2007-211334	Aug 2007	JP
2007-231369	Sep 2007	JP
2007-302918	Nov 2007	JP
2007-327098	Dec 2007	JP
2008-266778	Nov 2008	JP
2009-084648	Apr 2009	JP
2009084648	Apr 2009	JP
2009-270171	Nov 2009	JP
2010-065307	Mar 2010	JP
10-2004-0075971	Aug 2004	KR
WO 2009096344	Aug 2009	WO
WO 2009119751	Oct 2009	WO
WO 2009125874	Oct 2009	WO
WO 2011062151	May 2011	WO

High-strength steel sheet and high-strength zinc-coated steel sheet which have excellent ductility and stretch-flangeability and manufacturing method thereof

Information

Patent Number

Date Filed

Date Issued

Inventors

Original Assignees

Examiners

Agents

CPC

Field of Search

US

CPC

International Classifications

Abstract

Description

Claims

Priority Claims (2)

PCT Information

US Referenced Citations (4)

Foreign Referenced Citations (22)

Non-Patent Literature Citations (4)

Related Publications (1)

Entry
International Search Report issued in PCT/JP2011/071222, mailed on Dec. 20, 2011.
Notice of Allowance dated Aug. 23, 2013 issued in Korean Patent Application No. 10-2013-7006419.
Office Action dated Sep. 29, 2013 issued in Chinese Patent Application No. 201180044334.3.
An extended European Search Report for Appl. No. EP11825267.5 dated Jun. 24, 2014.