JP2011012308A - High-yield-ratio type hot-rolled steel plate superior in burring property and manufacturing method therefor - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high-yield-ratio type hot-rolled steel plate having superior burring properties while obtaining a tensile strength of a 370-490 MPa grade by a strength grade, and to provide a manufacturing method therefor.SOLUTION: The high-yield-ratio type hot-rolled steel plate with high burring properties includes components in a predetermined range, such a Ti content (wt.%) as to satisfy mathematical expression (1), Si and Mn of which the total amount is limited according to the amount of Ti, and the balance Fe with unavoidable impurities; has a microstructure in which pro-eutectoid ferrite occupies 90% or more by an area rate, and an average crystal grain size is 5-12 μm; has an elongation rate of 1.2-3; includes precipitates of TiC or NbC in the crystal grains of the microstructure, of which the average particle diameter is 1.5-3 nm and of which the density is 1×10to 5×10pieces/cm.

Description

  The present invention relates to a hot-rolled steel sheet and a method for producing the hot-rolled steel sheet, and more particularly, to a high yield ratio hot-rolled steel sheet excellent in burring suitable for obtaining a tensile strength of 370 to 490 MPa as a strength grade and a method for producing the same. .

  In recent years, there has been an increasing demand for steel plate thickness reduction for the purpose of reducing the weight of the vehicle body in order to improve the fuel efficiency of automobiles, etc., and how to maintain the material properties such as strength and press formability of the steel plate on the premise of thinning. It is a challenge to do.

  Application of high-strength steel sheets has been studied as a means for reducing the thickness of steel sheets to solve such problems. However, increasing the strength of a steel sheet that exceeds a strength grade of 540 MPa in terms of tensile strength generally results in deterioration of material properties such as press formability (workability). This is not an effective solution because it requires a lot of burden and special consideration.

  Here, the design strength of the press part constituting the automobile is based on the yield strength of the steel plate, not the tensile strength of the steel plate, which is an index for determining the strength grade of the steel plate. For this reason, in order to achieve the thinning that solves the above-mentioned problems, it is not always necessary to increase the tensile strength and rank up the strength grade of the steel sheet, and the yield strength is maintained while keeping the tensile strength of the steel sheet as it is. A means for increasing, that is, a means for increasing the yield ratio, which is the ratio between the tensile strength and the yield strength, can be considered.

  Up to now, it has been said that an increase in yield ratio leads to deterioration of uniform elongation and deteriorates press formability, but this is a view on molding methods that depend on uniform elongation such as stretch forming. It does not apply to molding methods that do not depend on uniform elongation, such as bending molding, drawing molding, and stretch flange molding. In particular, as the press forming characteristics required for inner plate members, structural members, and suspension member steel plates to which hot-rolled steel plates are applied, stretch flangeability (burring properties) is important, and high yield ratio steel plates should be used. If the mold and press process are optimized based on the premise, even if there is a deterioration of uniform elongation, it does not matter.

  For this reason, it can be said that increasing the yield ratio of the steel sheet is an effective solution as a means for reducing the wall thickness to solve the above-described problems. In particular, increasing the yield ratio of steel sheets does not accompany a significant increase in the strength grade of steel sheets, which is an index for determining the unit price of steel materials, and can reduce the amount of steel materials used by reducing the thickness of the steel. This makes it possible to reduce the weight of the vehicle body without increasing costs.

  In order to realize a high yield ratio of a steel sheet, there are two means: (1) refinement of crystal grains and (2) application of precipitation strengthening. Both introduce dislocation motion obstacles directly related to the yielding phenomenon of metallic materials. The obstacle here refers to a crystal grain boundary in the former and a precipitate in the latter.

  At present, addition of Nb is known as a specific means for crystal grain refinement that can be industrially realized (see, for example, Non-Patent Document 1). By adding Nb, grain growth and recrystallization of austenite grains during hot rolling are suppressed, and further, the ferrite grain size after the γ → α transformation can be reduced, thereby yielding the steel sheet. The ratio can be improved.

  On the other hand, as an example of applying precipitation strengthening, the steel structure of the hot-rolled steel sheet after hot rolling is made into a uniform structure of a single phase of ferrite, and further, the ferrite is precipitation strengthened with a composite precipitate of Ti and Mo, and high burring. Technology that realizes high properties, high yield strength, and high strength is disclosed (for example, see Patent Documents 1 and 2).

  Incidentally, conventionally, it is not common to add precipitation strengthening elements such as Ti in order to obtain a tensile strength of 370 to 490 MPa as a strength grade, and solid solution strengthening using C-Si-Mn based components. Alternatively, a means for obtaining a desired tensile strength within this range is usually taken by strengthening the structure.

JP 2002-322540 A JP 2002-322541 A

"Iron and Steel Vol.57, No.11" published in 1971, published by Japan Iron and Steel Institute, pp S624

  However, when the yield ratio is increased by adding Nb as in Non-Patent Document 1, the formation of a transformation texture having a large in-plane anisotropy is caused by suppressing recrystallization by Nb. The problem fluctuates depending on the rolling direction, and sufficient uniformity cannot be obtained. Furthermore, since the in-plane anisotropy is large, there is a problem that the burring property is lowered.

  In addition, when a high yield ratio is achieved by precipitation strengthening with Ti or the like as in Patent Documents 1 and 2, in order to sufficiently develop the strengthening ability by precipitation strengthening, phase interface precipitation in the γ → α transformation after winding Therefore, it was necessary to finely disperse the precipitate. For this reason, in this case, the temperature of the hot-rolled steel sheet at the time of winding is set to around 600 ° C., which is the precipitation temperature range of Ti, etc., and the microstructure of the hot-rolled steel sheet is also in this temperature range around 600 ° C. In order to become austenite, a large amount of austenite former alloys such as Mo and Mn had to be added. As a result, the tensile strength became 590 MPa or more, and it was not possible to obtain a hot rolled steel sheet having a tensile strength of 370 to 490 MPa with the strength grade intended by the present invention.

  Further, the solid solution strengthening as described above has a small strengthening ability compared with the strengthening ability of the structure strengthening and precipitation strengthening, and the solid solution strengthening elements such as Mo and V are expensive in terms of cost. It could not be said to be a preferable means.

  Moreover, when using the structure strengthening (quenching strengthening) as described above, it is essential to increase the cooling rate after hot rolling and take up at a lower temperature in order to utilize a phase that transforms at a low temperature. However, in order to obtain a low-temperature transformation phase, the aim of the coiling temperature must be set to 400 to 450 ° C., which is a transition boiling region where local temperature irregularities are likely to occur, and the accuracy is poor. Material variations and yield reduction were caused. In addition, since the generation of the low temperature transformation phase lowers the yield ratio rather by introducing movable dislocation due to transformation strain, the high yield ratio while obtaining the tensile strength of 370 to 490 MPa class with the intended strength grade of the present invention. It was difficult to get.

  Therefore, the present invention has been devised in view of the above-mentioned problems, and the object is to obtain a high yield with excellent burring while obtaining a tensile strength of 370 to 490 MPa class as a strength grade. It is to provide a specific hot rolled steel sheet and a method for producing the same.

  In order to solve the above-mentioned problems, the present invention takes into account the manufacturing process of a steel sheet having a tensile strength of 370 to 490 MPa as a strength grade, which is produced on an industrial scale by a production facility that is currently normally employed. It was made. In the present invention, by adding a precipitation strengthening element such as Ti, a high yield ratio that has not been obtained so far can be obtained, and as a result, a design strength equivalent to the case where a steel sheet of a half to one grade in strength grade is applied. The first object of the present invention is to provide a high yield ratio hot-rolled steel sheet with excellent burring properties that can be easily formed even with parts that require severe stretch flange processing due to excellent burring properties. In addition, the present invention can stably produce such a hot-rolled steel sheet at a low cost, and without adding a large amount of an austenite former alloy such as Mo and Mn at the time of production, further, the transition boiling region is around 400 ° C. The second object is to provide a method for producing a high yield ratio hot rolled steel sheet having excellent burring properties, which is excellent in yield, operability, and other manufacturability because it can avoid the coiling temperature. is there.

・・・（１）

・・・（２） That is, the high yield ratio hot rolled steel sheet having excellent burring properties according to the first invention is in mass%, C: 0.03 to 0.07%, Si: 0.005 to 1.8%, Mn: 0.1 to 1.9%, P ≦ 0.05% (not including 0%), S ≦ 0.005% (not including 0%), Al: 0.001 to 0.1%, N ≦ 0.005%, Nb (excluding 0%): 0.002 to 0.008%, S content (% by mass) [S], N content (% by mass) [N] In the following formula (1), Ti is contained in an amount (mass%) represented by [Ti], Si content (mass%) is [Si], and Mn content (mass%) is [ Mn], the following formula (2) is satisfied, the balance is a steel plate made of Fe and inevitable impurities, and 90% by volume or more of the microstructure is pro-eutectoid ferrite. Yes, the average crystal grain size is 5 μm to 12 μm, the degree of spread is 1.2 to 3, and the average grain size of the precipitate made of TiC in the crystal grains of the microstructure is 1.5 to 3 nm. In addition, the density is 1 × 10 ^{16 to} 5 × 10 ¹⁷ pieces / cm ³ .

... (1)

... (2)

  The high yield ratio hot-rolled steel sheet having excellent burring properties according to the second invention is the same as that in the first invention, further in mass%, Ca: 0.0005-0.005%, REM: 0.0005-0. 0.02% of any one or two kinds are contained.

  The high yield ratio hot-rolled steel sheet having excellent burring properties according to the third invention is characterized in that, in the first or second invention, the surface is galvanized.

  The method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to the fourth invention, after heating the steel slab containing the component according to any one of the first to third to 1000 ° C. or higher. Rough rolling is performed, and finish rolling in which the total reduction ratio of the final stage and the preceding stage is 30 to 45% is performed in a temperature range of 880 ° C. or more, and 1.5 to 3.5 seconds after finishing rolling. After air cooling, it is cooled to a temperature range of 750 to 620 ° C. at a cooling rate of 20 to 50 ° C./sec, then air cooled for 1 to 5 seconds, and further cooled to a temperature range of 620 to 480 ° C. at 2 to 10 ° C./sec. It is characterized by winding after cooling at a speed.

・・・（３） The method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to the fifth invention is the fourth invention, wherein the C content (% by mass) is [C] when the steel slab is heated. In this case, heating is performed to a temperature SRT (° C.) or higher that satisfies the following formula (3).

... (3)

  The method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to the sixth aspect of the invention is the fourth or fifth aspect, wherein the finish rolling is performed at a rolling start temperature of 1050 ° C. or higher. It is characterized by that.

  A method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to a seventh aspect of the present invention is the method according to any one of the fourth to sixth aspects, wherein a rough bar obtained by rough rolling the steel slab is obtained. The heating is performed from the end of the rough rolling to the start of the finish rolling and / or during the finish rolling.

  According to an eighth aspect of the present invention, there is provided a method for producing a high yield ratio hot rolled steel sheet having excellent burring properties. In any one of the fourth to seventh aspects of the present invention, the hot rolled steel sheet obtained after the winding is galvanized. It is characterized by being dipped in and galvanizing the surface.

  The method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to the ninth invention is characterized in that, in the eighth invention, the hot rolled steel sheet is galvanized and then alloyed.

  According to the first to ninth inventions, a hot-rolled steel sheet having a tensile strength of 370 to 490 MPa as a strength grade, a high yield ratio, excellent burring properties, and a small in-plane anisotropy is obtained. Is possible. In this way, it is possible to obtain a tensile strength of 370 to 490 MPa class with a strength grade while obtaining a high yield ratio and excellent burring properties. Equivalent design strength can be expected, and excellent burring makes it possible to easily form parts that require severe stretch flange processing. For this reason, it can be said that the present invention has a high industrial value.

  According to the fourth to ninth inventions, a high yield ratio is obtained by precipitation strengthening with Ti, although a large amount of austenite former alloy such as Mo and Mn is not added. It is possible to obtain a hot-rolled steel sheet having a tensile strength of 370 to 490 MPa with a strength grade that could not be achieved. In addition, since it is not necessary to set the winding temperature to around 400 ° C., which is the transition boiling region, the ease of manufacturing such as yield and operability is excellent.

It is a figure which shows the relationship of the yield ratio with respect to effective Ti amount (Ti *). It is a figure which shows the measurement result of yield ratio and (DELTA) r in the relationship between the average crystal grain diameter of a microstructure, and a degree of elongation. It is a figure which shows the relationship of the yield ratio with respect to the average particle diameter and density of a precipitate.

  Hereinafter, as a form for carrying out the present invention, a high yield ratio hot rolled steel sheet excellent in burring properties and a method for producing the same will be described in detail.

  In addition, this invention aims at obtaining the steel plate of the range of 370-490MPa grade by the strength grade about tensile strength, and specifically, it is 370MPa or more and less than 540MPa in tensile strength including a product tolerance. The aim is to obtain a range of steel sheets. Further, according to the present invention, the yield ratio is 85% or more, the hole expansion ratio which is an index indicating burring property is 120% or more, and | Δr | which is an index indicating in-plane anisotropy is 0.3 or less The purpose is to obtain a steel plate.

  First, the basic research results that led to the completion of the present invention will be described.

  The inventors have conducted extensive research to obtain a hot-rolled steel sheet that improves the yield ratio and further has excellent burring properties while suppressing in-plane anisotropy with a tensile strength in the range of 370 MPa to less than 540 MPa. As a result, we came to the conclusion that the following points are important.

  That is, from the viewpoint of obtaining a hot-rolled steel sheet having the above-mentioned properties, first, the yield ratio is improved by precipitation strengthening by addition of Ti and refinement by addition of Nb. Defining the upper limit of Nb that promotes the property, and secondly, reducing alloy elements such as Ti, Mn, and Si as much as possible in order to keep the tensile strength in the range of 370 MPa or more and less than 540 MPa, The average grain size and elongation of the microstructure to obtain the yield ratio and | Δr | to achieve, and fourth, to sufficiently exhibit the precipitation strengthening ability by TiC, It was found that these four points are important to define the range of the average particle size and density of the precipitate made of TiC.

  Therefore, the inventors have conducted detailed investigations and experiments on chemical components, microstructures, and precipitate conditions that can achieve a high level of tensile strength, yield ratio, and burring properties based on the above new findings. Carried out.

・・・（４） FIG. 1 is a diagram showing the relationship of the yield ratio with respect to the mass fraction of Ti * expressed by the following mathematical formula (4) obtained as a result. Note that Ti * means an effective amount of Ti that can be obtained by considering stoichiometry and can contribute to precipitation strengthening as TiC by combining with C. [Ti], [N] in the following formula (4) ] And [S] mean the contents in mass% of Ti, N, and S in the steel sheet, respectively.

... (4)

In order to improve the yield ratio of the hot-rolled steel sheet to 85% or more, which is the object of the present invention, it is necessary to deposit a sufficient amount of TiC precipitates in the ferrite. As shown in FIG. 1, the effective amount of Ti (Ti *) that binds to C even after coarse precipitates such as TiN, TiS, or Ti ₄ C ₂ S ₂ precipitated in a relatively high-temperature austenite region are precipitated. It was found that 0.02% or more is necessary.

・・・（１） As a result of conducting further investigations and experiments based on this fact, in order to obtain the intended tensile strength and yield ratio of the present invention, Ti is an amount represented by [Ti] in the following formula (1). It has been found that it is sufficient to contain in the steel sheet. In addition, the lower limit value of Ti in the following mathematical formula (1) is a numerical value obtained by developing the above mathematical formula (4), and is a numerical value that should be satisfied in improving the yield ratio. Further, the upper limit value of Ti is a numerical value that should be satisfied when the tensile strength is reduced to fall within the range of 370 MPa to less than 540 MPa.

... (1)

・・・（２） Moreover, it turned out that Si and Mn should be contained in the steel sheet so that the following mathematical formula (2) is satisfied when the tensile strength falls within the range of 370 MPa or more and less than 540 MPa on the assumption of Ti addition. did. The lower limit value of [Si] + [Mn] in the following formula (2) is a numerical value that should be satisfied when the tensile strength is increased to a range of 370 MPa to less than 540 MPa, and the upper limit value reduces the tensile strength. It is a numerical value that should be satisfied when the range is 370 MPa or more and less than 540 MPa. In addition, [Si] and [Mn] in the following numerical formula (2) mean the contents in mass% with respect to Si and Mn in the steel sheet, respectively.

... (2)

  FIG. 2 is a graph showing the relationship between the yield ratio and | Δr | with respect to the average crystal grain size and elongation of the microstructure obtained as a result of detailed investigation and experiment. As shown in this figure, the yield ratio tends to decrease with the coarsening of the grain size of the microstructure, and the tendency becomes stronger when the average crystal grain size exceeds 10 μm, and the average crystal grain size becomes 12 μm. It has been found that the yield ratio falls below the target value of 85%. In addition, as shown in FIG. 2, | Δr | tends to decrease with an increase in the degree of extension of the microstructure, and when the degree of extension exceeds 3, | Δr | becomes more than 0.3, which is the target value. I understood that.

FIG. 3 is a diagram showing the relationship of the yield ratio to the average grain size and density of precipitates made of TiC in the crystal grains of the microstructure obtained as a result of detailed investigation and experiment. As shown in this figure, the yield ratio tends to decrease as the density of the precipitates decreases, and it is found that the target value of 85% is interrupted when the density is less than 1 × 10 ¹⁶ pieces / cm ^3. It was. Further, it was found that it is important that the average particle size of the precipitates is 1.5 to 3 nm in order to obtain a yield ratio of 85%. This is because when the average particle size is 1.5 nm, the precipitation state is sub-aging and a sufficient number of precipitates are not deposited, and the density is less than 1 × 10 ¹⁶ pieces / cm ³ . When the average particle size is 3 nm, the precipitation state is overaging, and the number of precipitates is reduced by Ostwald growth, and the density becomes less than 1 × 10 ¹⁶ particles / cm ³ .

As a result of further detailed investigations and experiments based on the above idea, as a result of hot rolling with the characteristics that ferrite is sufficiently precipitation strengthened by TiC, the yield ratio is high, the in-plane anisotropy is small, and the burring property is excellent. In order to obtain a steel sheet, 90% by area or more of the microstructure is polygonal pro-eutectoid ferrite, the average crystal grain size of the microstructure is 5 μm to 12 μm, and the extension of the microstructure is 1. The average particle size of precipitates made of TiC in the crystal grains of the microstructure is 1.5 to 3 nm, and the density thereof is 1 × 10 ^{16 to} 5 × 10 ¹⁷ pieces / cm ³ . It turned out to be necessary.

  In addition, the inventors have found that the following points are important from the viewpoint of obtaining a method for producing a hot-rolled steel sheet having the target properties, including the above-described component ranges. That is, since the temperature of the γ → α transformation point increases due to the reduction of the austenite former alloy such as Mn, the γ → α transformation after winding at around 600 ° C. and the precipitation of precipitates at the γ / α phase interface. At the same time, it becomes impossible to promote and a hot-rolled steel sheet with a high density of precipitates cannot be obtained, so a new hot-rolled steel sheet with a high density of precipitates can be obtained. It is important to build a simple process (utilization of proeutectoid ferrite).

  If the transformation point temperature of the γ → α transformation is around 600 ° C. as in the prior art, the solubility product of TiC at γ and the solubility product at α when the winding temperature is around 600 ° C. Precipitation nuclei are preferentially generated at the γ / α phase interface using the difference as the driving force, and the subsequent precipitation nuclei do not grow because the temperature is about 600 ° C., and the density is reduced by coarsening the precipitates. Never do.

  On the other hand, when the transformation point temperature of the γ → α transformation is high due to the suppression of the addition of austenite former alloy such as Mn as a component of the steel plate as in the present invention, from hot rolling to winding The γ → α transformation starts during the cooling. In this case, if it is uniformly cooled from hot rolling to winding, TiC is preferentially nucleated at the γ / α phase interface at a higher temperature than before, and the precipitated TiC is higher than before. Precipitation nuclei grow even if cooling is performed for a short time, and the precipitates become coarser, resulting in a lower density, resulting in a hot-rolled steel sheet with a high density of precipitates. It will not be possible.

  Therefore, we investigated a cooling control process that can effectively utilize precipitation strengthening even when γ → α transformation occurs during cooling from hot rolling to winding. However, it is important not to use phase interface precipitation at the γ / α transformation interface, but to nucleate TiC as homogeneously as possible from ferrite in which Ti is dissolved in supersaturation as much as possible, and not to grow the precipitation nuclei. I came to think that there was.

  Specifically, in order to nucleate TiC as homogeneously as possible from ferrite in which Ti is supersaturated, it has been found that it is important not to diffusely transform austenite during the γ → α transformation but to perform a massive transformation. did. That is, after hot rolling, it is important to quickly cool to a temperature range in which massive transformation is promoted before γ → α transformation, and to sufficiently austenite to transform massively by air cooling in that temperature range for a predetermined time. As a result, it has been found that it is possible to obtain ferrite in which Ti is supersaturated and to uniformly nucleate TiC in the ferrite phase. Furthermore, in order not to grow the precipitation nuclei, it has been found that it is important to cool at a predetermined cooling rate to a temperature range in which nuclei growth becomes difficult after air cooling in a temperature range in which massive transformation is promoted. As a result, a hot-rolled steel sheet having a high TiC precipitation state can be obtained, and the precipitation strengthening ability of TiC can be sufficiently exhibited.

Based on the above findings, further investigations and experiments have revealed that the microstructure of the numerical conditions such as the intended hot-rolled steel sheet, hot rolling to obtain precipitates, and the subsequent cooling conditions are concrete. Specifically, finish rolling in which the total reduction ratio of the final stage and the preceding stage is 30 to 45% is performed in a temperature range of 880 ° C. or higher, and air cooling is performed for 1.5 to 3.5 seconds after finishing rolling. After cooling to a temperature range of 750 to 620 ° C. at a cooling rate of 20 to 50 ° C./sec, air cooling is performed for 1 to 5 seconds in the temperature range, and further cooling to 2 to 10 ° C./sec to a temperature range of 620 to 480 ° C. It was determined that it was necessary to wind up after cooling at speed.

  The present invention has been devised based on the knowledge as described above, and has constituent elements as described below. Hereinafter, the reason for limitation of each component will be described.

  First, the reasons for limiting the chemical components of the present invention will be described in detail. In the following, the description regarding the mass fraction in the composition is simply expressed as%.

  C is one of the most important elements in the present invention. If the C content exceeds 0.07%, the carbide that becomes the starting point of stretch flange cracking increases and burring properties deteriorate, so the C content is set to 0.07% or less. Further, if the C content is less than 0.03%, TiC sufficient to improve the yield ratio by precipitation strengthening cannot be obtained, so the C content is set to 0.03% or more.

  If the content of Si is less than 0.005%, precipitation during cooling of iron carbide, which becomes the starting point of stretch flange cracking, increases and burring properties deteriorate, so 0.005% or more should be added. . Further, if the content of Si exceeds 1.8%, the tensile strength increases excessively due to solid solution strengthening, and the tensile strength becomes 540 MPa or more, so 1.8% or less is added. And Si is added so as to satisfy the following formula (2).

  Since Mn contributes to the tensile strength of the steel sheet as a solid solution strengthening element, it is added according to the required tensile strength. If the content of Mn is less than 0.1%, the effect is lost and the necessary tensile strength cannot be obtained, so 0.1% or more is added. Further, if the content of Mn is more than 1.9%, the tensile strength increases excessively due to solid solution strengthening or quenching strengthening, and the tensile strength becomes 540 MPa or more, so 1.9% or less. Added. Mn is added so as to satisfy the following formula (2). In addition, when the element which suppresses the generation | occurrence | production of the hot crack by S other than Mn is not fully added, it is desirable to add the quantity of Mn of [Mn] / [S]> = 20 by mass%.

・・・（２） Since Si and Mn are both solid solution strengthening elements, they are added so as to satisfy the following formula (2) from the viewpoint of obtaining a tensile strength of 370 MPa or more and less than 540 MPa as described above. When [Si] + [Mn] is less than (1.38-18.5 × [Ti])%, the tensile strength is less than 370 MPa, and exceeds (3.23-18.5 × [Ti])%. In this case, the tensile strength becomes 540 MPa or more, and the intended tensile strength of the present invention cannot be obtained. Further, Si and Mn are required to be not more than the upper limit value in the following formula (2) from the viewpoint of controlling the γ → α transformation point temperature during cooling after completion of rolling. When [Si] + [Mn] is less than (3.23-18.5 × [Ti])%, the hardenability is increased and the Ar ₃ transformation point temperature becomes low. As a result, after the end of rolling, Precipitation due to TiC does not occur during the air-cooled holding performed in the third-stage cooling step, and the effect of precipitation strengthening may not be obtained. These numerical conditions are obtained based on experience.

... (2)

The lower the content of P, the more desirable it is, and if it exceeds 0.05%, the workability and weldability including burring properties are adversely affected, so the content is made 0.05% or less. However, in consideration of burring properties and weldability, the P content is preferably 0.02% or less. Note that the P content is never 0%.

  S not only causes cracking during hot rolling, but if it is too much, it produces A-based inclusions that degrade burring properties. Further, S reduces the effective Ti amount effective for fixing desired C by forming precipitates with Ti at a higher temperature than C. For this reason, S should be reduced as much as possible, but is acceptable if the content is 0.005% or less. When higher burring properties are required, the S content is preferably 0.003 or less. Note that the S content is never 0%.

  Al needs to be added by 0.001% or more for deoxidation of molten steel. If Al is added excessively, the cost increases, so the content is made 0.1% or less. Further, if the content of Al exceeds 0.06%, nonmetallic inclusions are increased and the elongation is deteriorated, so the content is desirably 0.06% or less.

  N, like S, forms precipitates with Ti at a higher temperature than C and reduces the effective Ti amount. Therefore, it should be reduced as much as possible, but if the content is 0.005% or less, it is an acceptable range. Note that the N content is never 0%.

  Nb is added in an amount of 0.002% or more because Nb has the effect of suppressing the grain growth of the ferrite grains after transformation due to the solution drag phenomenon and improving the yield ratio by refining. However, if the content exceeds 0.008%, recrystallization of austenite during hot rolling is suppressed, and ferrite obtained by transformation from the unrecrystallized austenite forms a texture. Since large in-plane anisotropy is caused in the hot-rolled steel sheet obtained later, the content is made 0.008% or less.

  Ti is one of the most important elements in the present invention. Ti not only contributes to the increase in the tensile strength of the steel sheet by precipitation strengthening, but also has the effect of improving the yield strength by improving the yield strength. Furthermore, Ti has the effect of suppressing the precipitation of coarse carbides such as cementite, which is the starting point of cracks, and improving the burring property.

・・・（１） Here, in order to improve the yield ratio, Ti *, which is the effective Ti amount, is required to be 0.02 or more as described above. Therefore, the addition amount of Ti is a numerical value equal to or greater than the lower limit in the following formula (1). It is necessary to. Further, when Ti is added in an amount of 0.07% or more, the strengthening ability due to precipitation strengthening is excessively increased, and the ferrite grains are excessively refined, so that the tensile strength may exceed 540 MPa. The effect is small compared with Nb, but there is an effect of suppressing recrystallization in austenite, and the in-plane anisotropy may be increased by the transformation texture from the non-recrystallized austenite. Less than 0.07%.

... (1)

  The above is the reason for limiting the basic components of the present invention. In the present invention, Ca, REM, B, Mo, V, Cr, Cu, Ni, Zr, Sn, Co, Zn, W are used as necessary. And one or two or more of the group consisting of Mg may be contained.

  Ca and REM are elements that are detrimental by changing the form of non-metallic inclusions that become the starting point of destruction and degrade workability. However, even if Ca and REM are both added in an amount of less than 0.0005%, there is no effect, and if Ca is added in excess of 0.005% and REM in excess of 0.02%, the effect is saturated. For this reason, it is preferable to add 0.0005 to 0.005% of Ca and 0.0005 to 0.02% of REM.

  B, Mo, V, Cr, Cu, Ni, Zr, Sn, Co, Zn, W and Mg may be contained in a total of 1% or less, but if over 1%, the strength may increase excessively. It is desirable to reduce as much as possible because it may affect the transformation behavior during cooling after rolling. Further, Sn is preferably 0.05% or less because wrinkles may occur during hot rolling. Further, Mo is an austenite former alloy, and its tensile strength may exceed 540 MPa when added in a large amount. Therefore, it is desirable to add less than 0.05%.

  Next, the reasons for limiting the numerical conditions related to the microstructure and precipitates of the steel sheet in the present invention will be described in detail.

  The microstructure must be proeutectoid ferrite with an area fraction of 90% or more. This is because in the manufacturing process of hot-rolled steel sheet, in order to obtain a high yield ratio by precipitating TiC homogeneously in a sufficient amount in a short time, the γ → α transformation is performed during cooling from hot rolling to winding. This is because 90% by area or more of the microstructure becomes polygonal shaped pro-eutectoid ferrite as a result. Moreover, when 90 area% or more of the microstructure is pro-eutectoid ferrite, excellent burring properties can be obtained. When the pro-eutectoid ferrite in the microstructure is less than 90% in area fraction, the area of ferrite, bainite and pearlite, which is sub-aging of precipitation strengthening transformed at a relatively low temperature after coiling other than pro-eutectoid ferrite. The rate increases and the yield ratio decreases and the burring property deteriorates.

Further, the microstructure is a ferrite in which most of the polygonal-shaped pro-eutectoid ferrite is transformed at a low temperature. In addition, the microstructure is allowed to contain 3% or less of the total amount of bainite, pearlite, retained austenite (γ _r ), MA (martensite-austenite constituent), etc.

  Proeutectoid ferrite has the property that the change in crystal orientation within the grain is extremely small, and the microstructure of ferrite transformed at a low temperature other than proeutectoid ferrite and other bainite changes in crystal orientation within the grain. Is relatively large. In addition, proeutectoid ferrite and ferrite transformed at low temperature have the property of being difficult to distinguish when observed with an optical microscope. For this reason, in order to discriminate pro-eutectoid ferrite in the microstructure, it is desirable to use the KAM method described later, whereby the pro-eutectoid ferrite and ferrite or bainite transformed at a low temperature other than pro-eutectoid ferrite are used. Discrimination becomes easy.

  The average crystal grain size of the microstructure needs to be 5 μm or more and 12 μm or less. If the average crystal grain size of the microstructure exceeds 12 μm, the grain size is excessively coarsened, so the yield ratio is reduced and the intended yield ratio cannot be obtained. In addition, if the average crystal grain size of the microstructure is less than 5 μm, the crystal grains are excessively refined, the tensile strength becomes 540 MPa or more and the desired tensile strength cannot be obtained, and the uniform elongation is reduced. Degradation of press formability including burring properties becomes significant.

The degree of extension of the microstructure needs to be 1.2-3. If the degree of extension is more than 3, ferrite grains in which the processed structure remains are often obtained, the elongation deteriorates, and the increase in in-plane anisotropy becomes remarkable even if the processed structure does not remain. Since the in-plane anisotropy increases as the degree of extension increases, the degree of extension is preferably 3 or less. On the other hand, in order to promote the transformation of pro-eutectoid ferrite in the component range of the present invention, when the final rolling which is the requirement of the present invention and the final rolling of the preceding stage is 30 to 45%, the elongation is 1. 2 or more. This is presumably because the grain growth in the thickness direction of ferrite during cooling after hot rolling or after winding was preferentially suppressed by the pinning effect of TiC.

  The area fraction of pro-eutectoid ferrite in the microstructure here is the sum of all pro-eutectoid ferrite areas in the measurement field divided by the field area in the measurement field. It can be determined by law. The average crystal grain size of the microstructure is obtained from a histogram showing the crystal grain size distribution of each grain in the measurement visual field. Specifically, a histogram with a section width of 1.0 μm is created for the crystal grain size of each grain in the measurement field of view, and the sum of the number ratios of the sections is added to the center value of each section. Is obtained. Here, the crystal grain size is a circle-equivalent diameter. Further, the area fraction of the microstructure and the average crystal grain size referred to here mean those for a cross section parallel to the rolling direction and the thickness direction of the steel plate.

・・・ （５） Further, the degree of extension is a numerical value indicating the degree to which the grains in the steel sheet containing proeutectoid ferrite crystal grains are extended by hot rolling, and is represented by e in the following formula (5). N ₁ in the following mathematical formula (5) means the number of crystal grains cut by virtual lines of a certain length extending in the plate thickness direction in a cross section parallel to the rolling direction and the plate thickness direction of the steel plate, n2 is in its cross-section, which means the number of the cut grain by imaginary line extending in the rolling direction line segments of the same length was determined n _1.

(5)

  As for the precipitate, TiC precipitated in the grains of the microstructure needs to satisfy the following conditions.

These precipitates need to have a density of 1 × 10 ^{16 to} 5 × 10 ¹⁷ pieces / cm ³ . This is because when the density is less than 1 × 10 ¹⁶ pieces / cm ^{3, the} precipitate interval that becomes a barrier for dislocation motion becomes too large, and the Orowan stress that is inversely proportional to the precipitate interval is reduced. This is because the desired yield ratio cannot be obtained. Further, if the density exceeds 5 × 10 ¹⁷ pieces / cm ³ , the precipitate size becomes so small that it cannot contribute to precipitation strengthening, so that sufficient precipitation strengthening cannot be obtained, and the intended tensile strength is obtained. This is because the yield ratio may not be obtained.

These precipitates need to have an average particle size of 1.5 to 3 nm. This is because if the deposition state is less than 1.5 nm, the precipitation state is sub-aging and a sufficient number of precipitates are not deposited, and the density is less than 1 × 10 ¹⁶ pieces / cm ³ , so that the intended yield ratio is This is because it cannot be obtained. If it exceeds 3 nm, the precipitation state is over-aged, and the number of precipitates is reduced by Ostwald growth, and the density becomes less than 1 × 10 ¹⁶ / cm ³ , so that the intended yield ratio is obtained. This is because there is not.

In addition, what is necessary is just to measure the average particle diameter and density of the precipitate which consists of TiC using a three-dimensional atom probe. Specifically, a needle-shaped sample is prepared through processing by a focused ion beam processing method as necessary while performing electropolishing after cutting out a sample to be measured. Next, a plurality of two-dimensional distribution images of the needle-shaped sample atoms are acquired in the depth direction of the needle-shaped sample using a three-dimensional atom probe, and the obtained two-dimensional distribution images are reconstructed in real space. Find the three-dimensional distribution image of atoms at. In order to obtain the average particle size and density of the precipitates in the microstructure grains as described later, it is only necessary to measure portions where the grain boundaries of the microstructure are not observed in the three-dimensional distribution image. .

  The average particle size of the precipitates can be obtained by arbitrarily measuring the diameters of 30 or more TiC precipitates and then arithmetically averaging them. In addition, the diameter of a precipitate here means the number of constituent atoms of each precipitate observed in a three-dimensional distribution image obtained by a three-dimensional atom probe, and has the same lattice constant as TiC and is measured. Assuming a spherical precipitate consisting of the number of constituent atoms of each precipitate, the diameter of the assumed precipitate is defined. The density of precipitates is obtained by dividing the number of TiC precipitates observed in the obtained three-dimensional image distribution image by the volume of the three-dimensional image distribution image. When observing a TiC crystal with a three-dimensional atom probe, Nb atoms may be included in the TiC crystal, so that such Nb atoms are the same as Ti atoms so as to obtain the diameter and density of the precipitate. It may be.

  Next, the reason for limitation of each manufacturing process in the manufacturing method of the hot rolled steel sheet according to the present invention will be described in detail.

  In obtaining the steel slab used as the object of hot rolling in this invention, the manufacturing process preceding hot rolling is not specifically limited. That is, following smelting by a blast furnace, a converter, an electric furnace, etc., the obtained molten steel is subjected to various secondary scouring to adjust the components so as to have the target component content as described above, followed by normal continuous casting, In addition to casting by the ingot method, a steel slab may be obtained by casting by a method such as thin slab casting. Scrap may be used as a raw material for the billet. When a slab as a steel slab is obtained by continuous casting, it may be sent directly to a hot rolling mill at a high temperature, or may be hot rolled after being cooled to room temperature and then reheated in a heating furnace.

・・・（３） When hot-rolling a steel slab, this steel slab is heated in a heating furnace. In obtaining the target tensile strength, yield ratio, λ, and anisotropic hot-rolled steel sheet in the present invention, the lower limit value of the heating temperature is not particularly limited. For example, heating is performed in a temperature range of 1000 ° C. or higher. do it. However, in order to obtain a hot-rolled steel sheet having such characteristics stably, it is preferable to heat at a temperature that satisfies the slab reheating temperature SRT (° C.) in the following formula (3) during this heating. . The following formula (3) is obtained by applying the solubility product formula of TiC in austenite (KJ Irvine, FB Pickering and T. Gladman: JISI, 205, (1967), p161). The solution temperature of TiC is indicated by the solubility product of TiC. If the temperature is lower than the slab reheating temperature, the coarse carbide of Ti generated during the slab production is not sufficiently dissolved, and the effect of precipitation strengthening by TiC may not be obtained in the subsequent cooling step. In addition, the retention time above the slab reheating temperature is preferably 20 minutes or more in order to promote sufficient dissolution of the coarse carbide of Ti. In addition, when the slab reheating temperature is 1400 ° C. or higher, the scale-off amount increases and the yield decreases, so the slab reheating temperature is preferably less than 1400 ° C. In addition, when the slab reheating temperature is less than 1100 ° C., the scale-off amount is small, and inclusions on the slab surface layer may not be removed together with the scale by subsequent descaling. Therefore, the slab reheating temperature is preferably 1100 ° C. or higher. .

... (3)

After heating the steel slab, hot rolling is performed on the steel slab extracted from the heating furnace. At the time of hot rolling, finish rolling is performed after roughly rolling the heated steel slab. There are no particular limitations on the rolling start temperature or rolling end temperature of rough rolling.

  After the completion of the rough rolling, finish rolling is performed in which the obtained rough bar is continuously rolled by a plurality of rolling mills. In finish rolling, it is preferable to make the rolling start temperature as high as possible together with the rolling end temperature (FT). This is because if TiC coarsely precipitates in austenite due to work-induced precipitation during finish rolling, the effect of precipitation strengthening by TiC may not be obtained in the subsequent cooling step. In particular, when the rolling start temperature of finish rolling is less than 1050 ° C., recrystallization does not proceed sufficiently, and there is a risk of increasing in-plane anisotropy due to a transformation texture that is γ → α transformed from unrecrystallized austenite grains. . Further, if the rolling start temperature of finish rolling is less than 1050 ° C., it matches the precipitation nose of TiC in the austenite region, and TiC is easily coarsened unless the rolling time in that temperature region is excessively shortened. Since it is required to tighten the restrictions, it is desirable that the rolling start temperature of finish rolling is 1050 ° C. or higher. Furthermore, in order to avoid the material deterioration in the width direction due to the temperature drop of the edge portion, it is desirable that the rolling start temperature of finish rolling is 1100 ° C. or higher. On the other hand, when the rolling start temperature of finish rolling exceeds 1150 ° C., scale is generated, and scale system defects such as scales and spindle scales may occur.

  In finish rolling, since it is necessary to promote precipitation of pro-eutectoid ferrite in the cooling step after completion of hot rolling in order to obtain the target microstructure of the present invention in the component system as described above, the final stage and It is necessary to perform rolling at a total rolling reduction of 30 to 45% in the preceding stage. If the total rolling reduction is less than 30%, sufficient pro-eutectoid ferrite cannot be obtained during cooling, and therefore precipitation strengthening during cooling does not proceed sufficiently, so that a high yield ratio cannot be obtained. On the other hand, when the total rolling reduction is greater than 45%, it is difficult to control the sheet passability and the plate shape, the plate thickness accuracy and the flatness may be deteriorated, and the austenite grains elongated by rolling undergo γ → α transformation. Since the ferrite grains are also elongated, the degree of elongation is increased and the in-plane anisotropy becomes remarkable. Therefore, the total reduction ratio of the last stage and the preceding stage is set to 30 to 45%.

  The finishing temperature (FT) in finish rolling is 880 ° C. or higher. When the rolling finish temperature (FT) in finish rolling is less than 880 ° C., even if Nb is virtually not added, recrystallization does not proceed sufficiently due to the weak recrystallization suppressing effect of Ti, and unrecrystallized austenite The texture formed by the γ → α transformation from the austenite is formed, and the ferrite grains transformed from the austenite grains elongated by rolling are also elongated, so that the degree of elongation increases and in-plane anisotropy may appear in the product plate. is there. There is no particular upper limit for the finish rolling temperature (FT) of finish rolling, but if it exceeds 980 ° C, recovery of dislocations that become ferrite nuclei is promoted, and the area fraction of proeutectoid ferrite may decrease. The finish rolling temperature (FT) of finish rolling is desirably 980 ° C. or lower.

  In order to make the rolling start temperature and the rolling end temperature in finish rolling as high as possible, a rough bar or rolled material is used between the end of rough rolling and the start of finish rolling and / or during finish rolling as necessary. Is preferably heated. Thereby, the finish temperature of finish rolling can be stably kept within the above-mentioned range. Any heating apparatus may be used in this case, but transverse induction heating is desirable because transverse induction heating can equalize the thickness in the thickness direction.

  Further, at the time of hot rolling, a subsequent rough bar may be joined to a preceding rough bar obtained by rough rolling, and finish rolling may be performed continuously. At that time, the coarse bar may be wound once in a coil shape, stored in a cover having a heat retaining function as necessary, and rewound again before joining.

A cooling process is performed after finishing rolling. The cooling process in the manufacturing method according to the present invention has four stages as follows. In the first cooling step, the coarse bar obtained after finishing rolling is air-cooled for 1.5 to 3.5 seconds. If this air cooling time is less than 1.5 seconds, recrystallization does not proceed sufficiently, and the ferrite grains that have undergone γ → α transformation from the austenite grains that have been elongated by rolling are also elongated, resulting in an increase in the degree of spreading. A texture due to transformation from recrystallized austenite is formed, and in-plane anisotropy may appear in the product plate. On the other hand, when the air cooling time exceeds 3.5 seconds, the precipitation of coarse TiC in austenite progresses, and the precipitation strengthening ability of ferrite in the subsequent third cooling step decreases, and Crystal grains become coarse and a high yield ratio cannot be obtained.

  In the second-stage cooling process performed after the first-stage cooling process, cooling is performed at a cooling rate of 20 to 50 ° C./sec. When the cooling rate is less than 20 ° C./sec, diffusion transformation proceeds during the γ → α transformation, and TiC is heterogeneously nucleated at the γ / α phase interface, and TiC is coarsened and reduced in density. In addition, the crystal grains of pro-eutectoid ferrite are also coarsened, and a high yield ratio cannot be obtained. When the cooling rate exceeds 50 ° C./sec, it is difficult to stop the cooling to the air cooling temperature range to be air-cooled in the next third-stage cooling process because of the cooling control. Transformation occurs, proeutectoid ferrite cannot be obtained, and furthermore, TiC does not precipitate sufficiently, so that a high yield ratio cannot be obtained.

  In the second-stage cooling process, cooling is performed to a temperature range of 750 to 620 ° C. that can promote massive transformation during the γ → α transformation. When this temperature is higher than 750 ° C., the transformation of pro-eutectoid ferrite is not sufficiently promoted in the cooling process of the next third stage, pearlite is generated and the burring property is lowered, and TiC has a γ / α phase interface. In this case, nucleation is inhomogeneously and the density is lowered, so that a high yield ratio cannot be obtained. On the other hand, if the temperature is higher than 750 ° C., the grain size of the pro-eutectoid ferrite advances and the crystal grain size becomes coarse, resulting in a decrease in yield ratio. Moreover, if this temperature is less than 620 ° C., bainite transformation occurs, proeutectoid ferrite cannot be obtained, and furthermore, TiC is not sufficiently precipitated, so that a high yield ratio cannot be obtained.

  In the third-stage cooling process performed after the second-stage cooling process, air cooling is performed for 1 to 5 seconds using the temperature range of 750 to 620 ° C. after the cooling in the previous second-stage cooling process as the air cooling start temperature. . This process is a fine and homogeneous TiC effective for the promotion of γ → α transformation by massive transformation to obtain proeutectoid ferrite and precipitation strengthening with the difference in solubility product between austenite and ferrite of TiC as precipitation driving force. It is an important process that promotes the precipitation of If the air cooling time in this step is less than 1 second, the transformation of pro-eutectoid ferrite is not sufficiently promoted, and furthermore, TiC is not sufficiently precipitated, so that a high yield ratio cannot be obtained. On the other hand, if the air cooling time exceeds 5 seconds, not only pearlite is generated and the burring property may be deteriorated, but also precipitation is overaged, the precipitation strengthening ability is lowered, and a high yield ratio cannot be obtained.

  In the fourth-stage cooling process performed after the third-stage cooling process, the cooling is performed to a temperature range of 620 to 480 ° C. that is a temperature range when winding is performed after the completion of the fourth-stage cooling process. When the coiling temperature is over 620 ° C., TiC that is finely and uniformly precipitated during the cooling process in the third stage becomes over-aged, the density is lowered by Ostwald growth, the precipitation strengthening ability is reduced, and a high yield ratio is obtained. I can't. On the other hand, when the coiling temperature is less than 480 ° C., the introduction of movable dislocations due to the increase in the low temperature transformation phase acts, and the yield ratio decreases. The temperature range below 480 ° C is a so-called transition boiling region in which local temperature unevenness is likely to occur, the winding rate is poor, the tolerance with respect to the target temperature is increased, the material variation is increased and the yield is increased. There is a concern that will lead to a decline.

In the cooling process at the fourth stage, cooling is performed at a cooling rate of 2 to 10 ° C./sec to the above-described temperature range where winding is performed. If the cooling rate is less than 2 ° C./sec, pearlite may be generated and the burring property may deteriorate, and the precipitated TiC may grow and become coarse, which may not contribute to precipitation strengthening. is there. If the cooling rate exceeds 10 ° C./sec, there is a concern that the tolerance becomes large with respect to the target temperature of the winding temperature, and the material variation becomes large.

  After completion of the winding process, pickling may be performed as necessary, and then, in-line or off-line, skin pass rolling with a reduction rate of 10% or less or cold rolling to a reduction rate of about 40% may be performed. In order to improve the ductility by correcting the shape of the steel sheet or introducing movable dislocations, it is desirable to perform skin pass rolling of 0.1% or more and 2% or less.

  Moreover, the hot-rolled steel sheet to which the present invention is applied may be subjected to a heat treatment in a hot dipping line in any case after casting, after hot rolling, and after cooling. You may make it perform a surface treatment separately. By applying the plating in the hot dipping line, the corrosion resistance of the hot rolled steel sheet is improved.

  Moreover, when galvanizing the hot-rolled steel sheet after pickling, the obtained steel sheet may be immersed in a galvanizing bath and alloyed as necessary. By applying the alloying treatment, the hot-rolled steel sheet has improved weldability with respect to various types of welding such as spot welding in addition to the improvement in corrosion resistance.

  In addition, the hot-rolled steel sheet in the present invention may be cooled after hot rolling obtained as described above, or may be subjected to heat treatment in a hot dipping line as described above, The surface treatment may be performed as described above.

  Next, the present invention will be further described with reference to examples.

  First, A to N steel slabs having chemical components shown in Table 1 were obtained. These steel slabs were obtained by continuous casting after secondary refining of molten steel obtained by melting in a converter. The obtained steel slab was directly fed or reheated after continuous casting, and was cooled to a sheet thickness of 1.2 to 5.5 mm by finish rolling following rough rolling, and then cooled and wound up. In addition, the display about the chemical composition in Table 1 is the mass%. The balance in the chemical composition shown in Table 1 is Fe and inevitable impurities. Moreover, the underline in Table 1 shows that it is outside the scope of the present invention.

  Table 2 shows details of manufacturing conditions from heating to winding of the steel slab. Here, “heating temperature” in Table 2 is the slab reheating temperature, “holding time” is the holding time at the slab reheating temperature, “rough bar heating” is from the end of rough rolling to the start of finish rolling and / or Or, whether or not the rough bar or rolled material is heated during finish rolling, “finish rolling start temperature” is the rolling start temperature of finish rolling, and “total rolling reduction” is the total rolling of the final stage and the preceding stage in finish rolling "Finish rolling finish temperature" is the finish rolling finish temperature, and "Time to start cooling" is the time from the finish rolling finish in the first stage cooling process to the start of the second stage cooling process. , “Cooling rate to air cooling zone” means the average cooling rate in the second stage cooling process, “Air cooling zone temperature” means the temperature at which air cooling is performed in the third stage cooling process, “Air cooling time” The cooling time in the third stage cooling process is defined as “cooling up to winding”. The degree "the average cooling rate at 4-stage cooling process, the" coiling temperature "indicates the coiling temperature. Steel No. 18 was galvanized. Furthermore, about the steel number 23, the alloying process was performed after galvanization. Under the present circumstances, the temperature immersed in a galvanization tank was 400-500 degreeC, and the alloying temperature was 500-650 degreeC. Moreover, the underline in Table 2 shows that it is out of the scope of the present invention or out of the preferred range.

  The microstructure and material properties of the steel sheet thus obtained were measured and evaluated as follows.

  The microstructure was investigated by polishing a sample cut out so that a cross section parallel to the rolling direction and the plate thickness direction was obtained at a position of 1/4 W or 3/4 W of the steel plate width W, and etched using a Nital reagent. The observation was carried out with a photograph of the visual field at a position of 1/4 t or 1/2 t below the surface layer 0.2 mm and the plate thickness t observed at a magnification of 200 to 500 times using an optical microscope.

For the measurement of the average crystal grain size, the degree of elongation, and the area fraction of pro-eutectoid ferrite, EBSP-OIM (Electron Back) was obtained from the microsample obtained from the sample cut above.
Scatter Diffraction Pattern-Orientation
Measurement was performed using an Image Microscopy method. The sample was polished with a colloidal silica abrasive for 30 to 60 minutes, and EBSP measurement was performed under the measurement conditions of 400 times magnification, 160 μm × 256 μm area, and measurement step of 0.5 μm.

  The EBSP-OIM method irradiates a highly inclined sample with a scanning electron microscope (SEM: Scanning Electron Microscope), irradiates an electron beam, takes a back-scattered Kikuchi pattern formed with a high-sensitivity camera, and performs computer image processing. By doing so, it is composed of a device and software for measuring the crystal orientation of the irradiation point in a short waiting time. The EBSP method can quantitatively analyze the microstructure and crystal orientation of the surface of the bulk sample, and the analysis area is an area that can be observed with an SEM. Depending on the resolution of the SEM, analysis can be performed with a minimum resolution of 20 nm. The analysis takes several hours and is performed by mapping tens of thousands of points to be analyzed in a grid at equal intervals. With polycrystalline materials, the crystal orientation distribution and crystal grain size in the sample can be seen.

  In this example, when determining the area fraction of the microstructure, the average crystal grain size and the elongation, the orientation difference of the crystal grains is generally recognized as a grain boundary in the EBSP-OIM method. The angle was defined as 15 °, which is a threshold value of the tilt grain boundary, and obtained based on an image in which grains were mapped so as to be visible.

  Further, the area fraction of pro-eutectoid ferrite was determined by a Kernel Average Misoration (KAM) method generally used together with the EBSP-OIM method. The KAM method is a method used to evaluate the azimuth fluctuation and distortion amount between pixels measured by the EBSP-OIM method.

  In the KAM method, six pixels adjacent to each other in a regular hexagon shape (first approximation) in measurement data, or 12 pixels (second approximation) further outside the six pixels, or the 12 Arithmetic average of the azimuth differences between the pixels of the 18 pixels (third approximation) further outside the pixels is arithmetically averaged, and a calculation is performed on each pixel with the obtained average value as the value of the center pixel. In the KAM method, when the azimuth difference between adjacent pixels is greater than or equal to a predetermined value, this azimuth difference is determined as a grain boundary, and the above calculation is repeatedly performed so as not to exceed the grain boundary. This makes it possible to create a map that represents the orientation change within the grain. This created map represents a strain distribution based on local orientation changes in the grains.

  In the analysis conditions in the present embodiment, the average value in the third approximation is obtained for each pixel based on the measurement data measured by the EBSP-OIM method, and the azimuth difference at this average value is 5 ° or less. Was displayed. Then, after determining that the set of pixels whose azimuth difference in the average value is 1 ° or less is one pro-eutectoid ferrite, the total area of all pro-eutectoid ferrite obtained in the measurement field of view is summed Divided by the visual field area of the measurement visual field was defined as the area fraction of pro-eutectoid ferrite.

  Tensile strength and yield ratio were obtained by processing the hot-rolled steel sheet obtained as described above into No. 5 test piece described in JIS Z 2201, and conducting the tensile test according to the test method described in JIS Z 2241. Based on the data obtained in this way.

  The burring property was evaluated by the hole expansion value obtained according to the hole expansion test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996.

・・・ （６） The in-plane anisotropy was evaluated by calculating | Δr |, which is an index of the in-plane anisotropy. | Δr | is obtained based on JIS Z2254, and is calculated from the Rankford values r ₀ , r ₄₅ , and r ₉₀ for each direction as shown in the following formula (6). R ₀ is a Rankford value in the 0 ° direction relative to the rolling direction, r ₄₅ is a Rankford value in the 45 ° direction relative to the rolling direction, and r ₉₀ is a 90 ° direction relative to the rolling direction. Rankford value. In order to obtain the Rankford value in each direction, the test piece used in the above-described tensile test was used.

(6)

  Table 3 shows details of the microstructure and material properties of the steel materials obtained by these measuring methods. Here, the microstructure other than the pro-eutectoid ferrite in the steel number described only as “F” in “Microstructure” of Table 3 is ferrite formed at a relatively low temperature. In addition, “F + P” and “F + B” in “Microstructure” in Table 3 indicate that the history passed through cooling the pearlite or bainite region in the processed CCT diagram (continuous cooling transformation diagram). The “proeutectic α fraction”, “average crystal grain size”, and “extensibility” are values obtained by the EBSP-OIM method described above. “YP”, “TS”, “El”, and “YR” are the yield strength, tensile strength, total elongation, and yield ratio obtained from the above-described tensile test, respectively. “Λ” is a hole expansion value obtained from the above-described hole expansion test. “| Δr |” is obtained from the tensile test described above. Moreover, the underline in Table 3 shows that it is outside the range of this invention or a preferable range.

In accordance with the present invention are steel Nos. 1, 2, 3, 4, 19, 20, 21, 22, 23, 24, all containing a predetermined amount of steel components that the present invention should satisfy. 90% by area or more of the microstructure is pro-eutectoid ferrite, the average crystal grain size is 5 μm to 12 μm, the elongation is 1.2 to 3, and from the TiC in the crystal grains of the microstructure A high yield ratio type excellent in burring characteristics, characterized in that the average crystal grain size of the resulting precipitate is 1.5 to 3 nm and the density thereof is 1 × 10 ^{16 to} 5 × 10 ¹⁷ pieces / cm ^3. A hot-rolled steel sheet is obtained. Therefore, the steels of these steel numbers have an objective tensile strength, yield ratio, hole expansion value and in-plane anisotropy index | Δr | of 370 to 540 MPa, 85% or more, 120 respectively. % Or more and 0.3 or less.

  Steels other than the above are outside the scope of the present invention for the following reasons. That is, in Steel No. 5, since the total reduction ratio of the final stage and the preceding stage is too low, the area fraction of pro-eutectoid ferrite becomes too low, and the target yield ratio is not obtained. Steel No. 6 has a finish rolling temperature that is too low, so that the degree of elongation becomes too large, and the target yield ratio and λ are not obtained, which exceeds the upper limit of in-plane anisotropy. In Steel No. 7, the time for the first cooling step is too short, so the degree of elongation becomes too large, the target λ is not obtained, and exceeds the upper limit of anisotropy. In Steel No. 8, the time for the first stage cooling process is too long, so the average crystal grain size becomes too large, and the target yield ratio is not obtained. In Steel No. 9, the cooling rate in the second stage cooling process is too slow, so the average crystal grain size becomes too large, and the target yield ratio is not obtained. In Steel No. 10, the cooling rate in the second-stage cooling process is too fast, so the area fraction of pro-eutectoid ferrite is too low, and the target yield ratio is not obtained. In steel No. 11, the temperature in the third stage cooling process is too high, so that pearlite is generated, and further, the area fraction of pro-eutectoid ferrite becomes too low, the average crystal grain size becomes too large, and is targeted. Yield ratio, λ is not obtained. In Steel No. 12, since the temperature in the third-stage cooling process is too low, the area fraction of pro-eutectoid ferrite becomes too low, and the target yield ratio is not obtained. In Steel No. 13, the temperature in the third-stage cooling process is too high, so the precipitates are reduced in density, and the average grain size and density are out of range, and the target yield ratio is not obtained. In Steel No. 14, since there is no air cooling time in the third stage cooling process, the area fraction of pro-eutectoid ferrite becomes too low, and the target yield ratio is not obtained. In Steel No. 15, the air cooling time in the third stage cooling process is too long, so pearlite is generated, and the area fraction of proeutectoid ferrite is too low, and the target yield ratio, λ, is obtained. Absent. In steel No. 16, the cooling rate in the fourth stage cooling process is too slow, so pearlite is generated, and the area fraction of proeutectoid ferrite is too low, and the target yield ratio, λ, is obtained. Absent. Steel No. 17 has a winding temperature that is too high, so the average crystal grain size becomes too large, and the target yield ratio is not obtained. In Steel No. 18, since the coiling temperature is too low, the area fraction of pro-eutectoid ferrite becomes too low, and the target yield ratio is not obtained. In Steel No. 25, the Nb content exceeds the range of the present invention, so that the degree of expansion becomes too large and exceeds the target upper limit of anisotropy. In Steel No. 26, since the C content exceeds the upper limit of the range of the present invention, pearlite is generated, the area fraction of pro-eutectoid ferrite becomes too small, and the target λ is not obtained. In Steel No. 27, since the C content is below the lower limit of the range of the present invention, the average crystal grain size becomes too large, the target yield ratio is not obtained, and the C content is low. From this, the precipitation of TiC is reduced, the pinning effect is suppressed, and the degree of expansion is below the range of the present invention. In Steel No. 28, since the Ti content is below the lower limit of the range of the present invention, the target yield ratio is not obtained with little contribution of precipitation strengthening. In Steel No. 29, since the Ti content exceeds the upper limit of the range of the present invention, the average crystal grain size becomes too small, and the tensile strength exceeds 540 MPa. In Steel No. 30, the content of [Si] + [Mn] exceeds the upper limit of the range of the present invention, so the tensile strength exceeds 540 MPa, and the target λ is not obtained. In Steel No. 31, the content of [Si] + [Mn] is below the lower limit of the range of the present invention, so the tensile strength is below 370 MPa.

Claims (9)

% By mass
C: 0.03-0.07%,
Si: 0.005 to 1.8%
Mn: 0.1 to 1.9%
P ≦ 0.05% (excluding 0%),
S ≦ 0.005% (excluding 0%),
Al: 0.001 to 0.1%,
N ≦ 0.005% (excluding 0%),
Nb: 0.002 to 0.008%
When the S content (% by mass) is [S] and the N content (% by mass) is [N], the amount (% by mass) represented by [Ti] in the following formula (1) ) Ti,
When the Si content (% by mass) is [Si] and the Mn content (% by mass) is [Mn], the following formula (2) is satisfied:
The balance is a steel plate made of Fe and inevitable impurities,
90% by area or more of the microstructure is pro-eutectoid ferrite, the average crystal grain size is 5 μm to 12 μm, and the elongation is 1.2 to 3,
The average grain size of the precipitate made of TiC in the crystal grains of the microstructure is 1.5 to 3 nm, and the density thereof is 1 × 10 ^{16 to} 5 × 10 ¹⁷ pieces / cm ^3. High yield ratio hot rolled steel sheet with excellent burring properties.

... (1)

... (2) Furthermore, in mass%,
Ca: 0.0005 to 0.005%,
REM: 0.0005 to 0.02%,
The high yield ratio hot-rolled steel sheet having excellent burring properties according to claim 1, wherein one or two of these are contained. The high yield ratio hot rolled steel sheet having excellent burring properties according to claim 1 or 2, wherein the surface is galvanized. After the steel slab containing the component according to any one of claims 1 to 3 is heated to 1000 ° C or higher, rough rolling is performed, and finish rolling in which the total rolling reduction of the final stage and the preceding stage is 30 to 45% The rolling end temperature is set to a temperature range of 880 ° C. or higher, air cooled for 1.5 to 3.5 seconds after finishing rolling, and then cooled to a temperature range of 750 to 620 ° C. at a cooling rate of 20 to 50 ° C./sec. After that, it is air-cooled for 1 to 5 seconds, and further cooled to a temperature range of 620 to 480 ° C. at a cooling rate of 2 to 10 ° C./sec. A method of manufacturing a steel sheet. The heating of the steel slab is performed at a temperature SRT (° C) or more that satisfies the following formula (3) when the C content (% by mass) is [C]. A method for producing a high yield ratio hot rolled steel sheet with excellent burring properties.

... (3) 6. The method for producing a high yield ratio hot rolled steel sheet having excellent burring properties according to claim 4, wherein the finish rolling is performed at a rolling start temperature of 1050 ° C. or higher. The rough bar obtained by rough rolling the steel slab is heated from the end of the rough rolling to the start of the finish rolling and / or during the finish rolling. The manufacturing method of the high yield ratio type | mold hot-rolled steel plate excellent in burring property of 1 item | term. The high yield yielding excellent burring property according to any one of claims 4 to 7, wherein the hot-rolled steel sheet obtained after the winding is immersed in a galvanizing bath and the surface thereof is galvanized. A method for producing a specific hot-rolled steel sheet. 9. The method for producing a high yield ratio hot rolled steel sheet with excellent burring properties according to claim 8, wherein the hot rolled steel sheet is galvanized and then alloyed.

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