WO2013099183A1

WO2013099183A1 - High-strength hot-rolled steel sheet and manufacturing method therefor

Info

Publication number: WO2013099183A1
Application number: PCT/JP2012/008190
Authority: WO
Inventors: 船川　義正; 珠子有賀; 永明森安; 貴幸村田; 浩大和田
Original assignee: Ｊｆｅスチール株式会社
Priority date: 2011-12-28
Filing date: 2012-12-21
Publication date: 2013-07-04
Also published as: TW201333220A; JPWO2013099183A1; TWI506147B; JP5644964B2

Abstract

Provided are: a high-strength hot-rolled steel sheet that combines excellent strength and workability; and a manufacturing method therefor. By mass, said steel sheet contains more than 0.035% and no more than 0.065% carbon, up to 0.2% silicon, up to 0.65% manganese, up to 0.03% phosphorus, up to 0.02% sulfur, up to 0.1% aluminum, up to 0.01% nitrogen, and 0.09-0.25% titanium, with the remainder comprising iron and unavoidable impurities. By area fraction, a ferrite phase constitutes more than 95% of the metal structure, titanium carbide with a mean grain diameter of 5 nm or less is finely dispersed inside the crystal grains of the ferrite phase, and the mean diameter of said ferrite crystal grains is at least 1 µm. This results in a high-strength hot-rolled steel sheet with a tensile strength of at least 780 MPa.

Description

High strength hot rolled steel sheet and method for producing the same

The present invention has a tensile strength of 780 MPa or more and excellent stretch flangeability suitable for structural materials such as automobiles, transportation equipment, and building equipment, and has mechanical properties in the coil. The present invention relates to a high-strength hot-rolled steel sheet, a high-strength hot-rolled plated steel sheet, and a method for manufacturing the same.

In order to reduce CO ₂ emissions from the viewpoint of global environmental conservation, maintaining the strength of automobile bodies while reducing their weight and improving automobile fuel efficiency has always been an important issue in the automobile industry. In order to reduce the weight of the vehicle body while maintaining the strength of the automobile body, it is effective to reduce the thickness of the steel sheet by increasing the strength of the steel sheet used as a material for automobile parts. Therefore, in recent years, high-strength hot-rolled steel sheets (thin steel sheets) having a tensile strength of 780 MPa or more have been actively applied as materials for automobile parts.

Here, since many automotive parts made of thin steel sheets are formed by pressing or burring, etc., steel sheets for automotive parts have high press strength and stretch flange processing in addition to high strength. It is required to have processability such as property. In addition, since automobiles are configured by combining many parts, individual parts are required to have a high degree of dimensional accuracy. Therefore, it is required for steel sheets for automobile parts to obtain parts having excellent dimensional accuracy by pressing or the like.
For the above reasons, when applying high-strength hot-rolled steel sheets to automobile parts, etc., development of high-strength hot-rolled steel sheets that have both strength and workability and that are capable of obtaining parts with excellent dimensional accuracy by pressing etc. Has become essential, and many studies have been made to date, and various techniques have been proposed.

However, in response to today's increasing demand for higher strength, non-uniform mechanical properties such as strength and workability of high-strength steel plates, that is, fluctuations in mechanical properties within the coils of high-strength steel plates, The application of high-strength steel sheets to parts has been hindered. In particular, non-uniformity in the steel strength coil induces fluctuations in the amount of springback that occurs during press processing (that is, the amount of springback that occurs in the pressed part becomes uneven in the pressed part), and the shape of the pressed part Make it unstable. In addition, if the steel plate strength becomes non-uniform, stretch flangeability becomes non-uniform, which may cause press cracks.

Originally, steel has a tensile strength of about 300 MPa class. Therefore, an increase in strength of a steel sheet, for example, an increase in strength due to a complex structure or refinement of crystal grains, causes variations in strength. This variation in strength is caused by temperature fluctuations and fluctuations in production in the longitudinal and width directions of the steel sheet in the rolling direction, and also by structural changes caused by differences in processing conditions. Therefore, in order to provide higher-quality automobile parts, in addition to imparting desired strength and workability to the steel sheet for automobile parts, these mechanical characteristics (particularly strength) do not vary within the coil. It is also extremely important to apply uniformly throughout the entire steel plate.

As a technique for improving the mechanical properties of a steel sheet for automobile parts, for example, in Patent Document 1, a high-strength steel sheet having a tensile strength of 500 MPa or more containing 60% or more of a ferrite structure is deformed after deformation of 20% or more. By including 50% or more of ferrite crystal grains with a structure in which dislocation cell structures aligned in one direction intersect in two or more directions in the region, the amount of springback during high-strength steel sheet pressing is suppressed. Thus, techniques for improving the shape freezing property of high-strength steel sheets have been proposed.

Patent Document 2 specifies the texture, r value, and uniform elongation anisotropy of a high-strength hot-rolled steel sheet, thereby improving the bending workability of the steel sheet and suppressing springback. A technique for improving the shape freezing property of the glass has been proposed. In addition, regarding the steel sheet structure, by ensuring that the ferrite volume fraction is 80% or less, the shape freezing property of the steel sheet is secured, and when the martensite or retained austenite volume fraction is 1% or more and 25% or less, the steel sheet A technique for ensuring the strength and formability of the steel has been proposed. Furthermore, in Patent Document 3, the texture and r value of the steel sheet are defined, the structure having the maximum area ratio in the steel sheet is made of ferrite, and the coarse cementite of the grain boundary is further reduced, thereby improving the bending workability of the steel sheet. Techniques have been proposed for improving and suppressing springback, ensuring the shape freezing property of the steel sheet, and improving stretch flangeability.

Patent Document 4 discloses that a steel sheet composition is a composition in which one or more of Ti, Mo, and W are added, and carbides of 10 nm or less containing these elements are dispersed in ferrite, whereby a high-tensile hot-rolled steel sheet is obtained. A technique for improving the stability of the yield strength in the width direction has been proposed. Patent Document 5 discloses a high-tensile hot-rolled steel sheet having a steel sheet composition in which at least one of Ti, Mo, and W is added, and carbides of 10 nm or less containing these elements are dispersed in ferrite. A technique for improving the strength stability in the coil longitudinal direction has been proposed.

JP 2007-308771 A JP 2004-250743 A JP 2002-363893 A JP 2003-321734 A JP 2003-321735 A

However, in the technique proposed in Patent Document 1, the volume fraction of the hard phase other than ferrite affects the steel plate strength, and the fluctuation of the hard phase volume fraction sensitive to changes in the manufacturing conditions unavoidably changes the steel plate strength. As a result, non-uniform mechanical properties are increased in the coil. Therefore, even if the steel sheet obtained by such a technique is subjected to press working to form a pressed part, press cracking occurs or the shape of the pressed part becomes unstable, which is an industrially feasible technique. hard. In addition, when the steel sheet is pressed into a desired part shape, the interface between the soft ferrite and the hard second phase is likely to be a starting point for cracking during processing, and the workability is not stable. .

In addition, with the technique proposed in Patent Document 2, it is difficult to stably obtain a desired texture throughout the entire length and width of the coil, and martensite and retained austenite are actively used as the steel sheet structure. As a result, the stability of the strength is significantly deteriorated. Therefore, even in the technique proposed in Patent Document 2, it is not possible to obtain a stable shape freezing property, and as in the technique proposed in Patent Document 1, when pressing a steel plate to form a pressed part, Problems such as press cracks and unstable shape of pressed parts are seen. Furthermore, even in the technique proposed in Patent Document 3, since it is extremely difficult to stably obtain a desired texture throughout the entire length and width of the coil, stable strength cannot be obtained. Problems similar to those proposed in 2 are observed.

Moreover, in the technique proposed by patent document 4, as the Example shows, the desired steel plate intensity | strength is ensured by adding 1% or more of Mn which is a solid solution strengthening element to a steel plate. Therefore, according to the technique proposed in Patent Document 4, the strength varies due to the segregation of Mn, and the fluctuation in strength in the width direction cannot be suppressed. Further, if the content of Mn, which is a solid solution strengthening element, is reduced for the purpose of suppressing segregation of Mn, the strength of the steel sheet is lowered, and a tensile strength of 780 MPa or more cannot be obtained.

On the other hand, Patent Document 5 proposes reducing the fluctuation in strength by reducing the Mn content of the steel sheet. However, all of the steel sheets disclosed in the examples of Patent Document 5 contain 1% or more of Mn, which is a solid solution strengthening element, for the purpose of ensuring a desired steel sheet strength. That is, in the technique proposed in Patent Document 5, Mn segregation is still large, and the stability of the tensile strength in the coil longitudinal direction is not guaranteed. Further, the amount of change in strength in the width direction varies depending on the position in the coil longitudinal direction, and there is room for improvement in material uniformity. Furthermore, if the Mn content, which is a solid solution strengthening element, is reduced for the purpose of suppressing segregation of Mn (“segregation” of “Mn”), the strength of the steel sheet decreases, and a tensile strength of 780 MPa or more cannot be obtained.

As described above, even the techniques proposed in Patent Document 4 and Patent Document 5 have not yet solved the problem of strength stability caused by segregation of Mn while maintaining a desired steel plate strength. Also, with these technologies, the high Mn content of the steel sheet induces cracking during press forming of the steel sheet due to Mn segregation, making it difficult to stably ensure excellent stretch flangeability. And sufficient stretch flangeability cannot always be obtained.

As described above, in the above-described conventional technology, the mechanical characteristics inside the coil, particularly the strength variation is large, while the steel sheet is excellent in mechanical properties such as high strength and high workability and high shape freezing property. Since the variation in stretch flangeability is large, it has not yet been a high-strength steel sheet that suppresses variations in strength and workability within the coil. For this reason, it is extremely difficult to industrially mass-produce press parts having desired strength and stable dimensional accuracy with the prior art.
An object of the present invention is to provide a high-strength hot-rolled steel sheet having a small fluctuation in tensile properties inside a coil, excellent stretch flangeability, stable component dimensional accuracy, and a method for producing the same, which has been made under such circumstances. And

In order to solve the above-mentioned problems, the present inventors, in addition to various factors affecting workability such as high strength of hot-rolled steel sheet and stretch flangeability, uniformity of mechanical properties in the coil of the steel sheet, especially steel sheet Various factors affecting the uniformity of strength were studied.
As a means for increasing the strength of a steel sheet to a tensile strength of 780 MPa or more while maintaining workability such as stretch flangeability, for example, there is a means for forming a composite structure by dispersing martensite structure or bainite structure in soft ferrite. It is done. However, when the steel sheet structure is a composite structure, the variation in each phase fraction in the coil and the change in the hardness of each phase overlap, and the steel sheet strength varies greatly within the coil. Therefore, the present inventors have made the metal structure (microstructure) into a single phase, not a high strength by complex structure, in order to suppress the steel plate strength change in the coil and ensure the uniformity of the strength in the coil. Therefore, I thought that it should be oriented to higher strength. In addition, the present inventors focused on the ferrite phase with excellent workability such as stretch flangeability, and examined a means for increasing the strength of the steel sheet after making the metal structure of the steel sheet a ferrite single phase structure. Proceeded.

As means for enhancing the strength of a steel sheet having a ferrite single phase structure as a metal structure, solid solution strengthening by increasing the amount of Mn added or strengthening means by crystal grain refinement can be considered. However, as a result of investigations by the present inventors, it is clear that all of these strengthening means cause fluctuations in the strength of the steel sheet in the coil and stretch flangeability, and are disadvantageous in ensuring the uniformity of the steel sheet strength. became.

If the amount of Mn in the steel is large, Mn segregates, the austenite transformation is delayed and it becomes easier to form a hard phase, or the amount of solid solution strengthening is larger than other parts, resulting in uneven steel plate strength and workability. Invite Therefore, when trying to increase the strength of the steel sheet by adding Mn, the tensile strength in the width direction of the steel sheet becomes non-uniform due to the Mn segregation described above, and the stretch flangeability becomes non-uniform due to this. The present inventors have found that.
Further, when the tensile strength of the steel sheet is made 780 MPa or more by grain size refinement, the crystal grain size of ferrite needs to be less than about 1 μm. However, it is not easy to make the ferrite grain size less than about 1 μm over the entire area of the steel sheet. The crystal grain size of ferrite greatly depends on the steel sheet manufacturing conditions, particularly the cooling conditions after the end of hot rolling. Because the cooling rate tends to be unstable at the end of longitudinal direction of hot coil or top end and bottom end of hot coil and at the end of width direction, the ferrite crystal grain size at the end of steel plate is also coarsened. Easy to coarsening.

For these reasons, the present inventors have further studied a means for increasing the strength of a steel sheet having a metal structure of a ferrite single phase structure regardless of high Mn solid solution strengthening or crystal grain refinement. As a result, suppressing the Mn content in the steel sheet and precipitating fine Ti carbide in the individual ferrite crystal grains forming the ferrite single-phase structure maintains the workability of the steel sheet (stretch flange workability, etc.). However, the present inventors have found that the tensile strength of the steel sheet is 780 MPa or more and that the fluctuation of the steel sheet strength is suppressed, and that it is extremely effective as a means for imparting uniform strength in the longitudinal direction and the width direction of the steel sheet.

In addition, in order to maintain the workability of the steel sheet (stretch flange workability, etc.) while the steel sheet has a tensile strength of 780 MPa or more, it is necessary to sufficiently refine the Ti carbide, and further, Ti carbide Ti It has been found that when the content is more than the C content by atomic ratio, the carbide tends to be coarsened, which may adversely affect the properties of the hot-rolled steel sheet.
On the other hand, if the crystal grain size of ferrite becomes extremely small, it becomes a factor of non-uniformity of steel sheet strength, so it has been found that ferrite average crystal grain size of 1 μm or more is effective in suppressing non-uniformity of steel sheet strength. .

Furthermore, Ti carbide precipitates during the cooling process after hot rolling in the hot-rolled steel sheet manufacturing process, but if the amount of Mn in the steel is large, Mn segregates, and the timing of Ti carbide precipitation is Mn segregated. It became clear that the Mn segregation part became excessively hard compared with the other parts, resulting in non-uniform steel plate strength. In addition, the present inventors have found a new technology that can stabilize the strength by reducing this, which is caused by the addition of 1% or more of Mn, which has been common knowledge in conventional high-strength steel sheets. It was.

Further, Ti carbide is a phase interface precipitation type precipitate that precipitates simultaneously with the austenite → ferrite transformation in the cooling process after the hot rolling in the hot rolled steel sheet manufacturing process. Here, as a result of the study by the present inventors, when the ferrite transformation point of the steel is significantly higher than the coiling temperature in the hot-rolled steel sheet manufacturing process, Ti carbide precipitated simultaneously with the austenite → ferrite transformation reaches the coiling temperature. It became clear that it became coarse during the cooling process and the desired steel plate strength could not be obtained.

Therefore, the present inventors suppress the coarsening of Ti carbide by adjusting the ferrite transformation point of the steel to the same level as the coiling temperature, and fine Ti carbide throughout the longitudinal and width directions of the steel sheet. It came to the idea of making it precipitate uniformly. And in order to adjust the ferrite transformation point of steel to the same level as the coiling temperature, the amount of Mn contained in the steel sheet, the steel sheet manufacturing conditions, especially the cooling rate after the hot rolling and the coiling temperature should be specified. Was found to be important.

The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%
C: more than 0.035% and 0.065% or less, Si: 0.2% or less,
Mn: 0.65% or less, P: 0.03% or less,
S: 0.02% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.09% or more and 0.25% or less, with the balance being composed of Fe and inevitable impurities, with an area ratio of more than 95% being the ferrite phase, and the ferrite phase crystal grains A high-strength hot-rolled steel sheet having a structure in which Ti carbide having an average particle size of 5 nm or less is finely dispersed, the ferrite crystal has an average crystal grain size of 1 μm or more, and a tensile strength of 780 MPa or more.

[2] The high-strength hot-rolled steel sheet according to [1], wherein an atomic ratio of C and Ti contained in the Ti carbide satisfies the following formula (1).
Record
Ti / C <1.0 (1)
(Ti / C: atomic ratio of C and Ti in Ti carbide)

[3] In the above [1] or [2], in addition to the composition, in addition to mass, Cu, Ni, Cr, Co, Mo, Sb, W, As, Pb, Mg, Ca, Sn, Ta, A high-strength hot-rolled steel sheet containing 1% or less in total of any one of Nb, V, REM, Cs, Zr, B, and Hf.

[4] The high-strength hot-rolled steel sheet according to any one of [1] to [3], wherein the steel sheet surface has a plating layer.

[5] The high-strength hot-rolled steel sheet according to [4], wherein the plating layer is a zinc plating layer or a zinc-containing alloy plating layer.

[6] The steel material is subjected to hot rolling consisting of rough rolling and finish rolling. After finishing rolling, the steel material is cooled, wound, and hot rolled steel sheet.
The steel material in mass%,
C: more than 0.035% and 0.065% or less, Si: 0.2% or less,
Mn: 0.65% or less, P: 0.03% or less,
S: 0.02% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.09% or more and 0.25% or less, with the balance being Fe and inevitable impurities,
The finish rolling temperature of the finish rolling is 840 ° C. or more and 1050 ° C. or less, the average cooling rate from the end of finish cooling of the cooling to 750 ° C. is 30 ° C./s or more, and the winding temperature of the winding is 570 ° C. or more and 750 ° C. A method for producing a high-strength hot-rolled steel sheet having a tensile strength of 780 MPa or more, characterized in that the temperature is not higher than ° C.

[7] In the above [6], in addition to the composition, in addition to mass, Cu, Ni, Cr, Co, Mo, Sb, W, As, Pb, Mg, Ca, Sn, Ta, Nb, V, A method for producing a high-strength hot-rolled steel sheet, comprising a total of 1% or less of any one of REM, Cs, Zr, B, and Hf.

According to the present invention, by optimizing the structure without adding a large amount of Mn or alloy elements, the mechanical property variation that could not be achieved conventionally is small, and the 780 MPa class (tensile strength: about 780 to 900 MPa) is high. It is possible to provide a high-strength hot-rolled steel sheet and a method for producing the same, and have a remarkable industrial effect.

FIG. 1 is a diagram showing a schematic shape of Ti carbide.

Hereinafter, the present invention will be described in detail.
First, the structure of the steel sheet of the present invention and the reasons for limiting the carbide will be described.
The hot-rolled steel sheet of the present invention has an area ratio of more than 95% of the ferrite phase, and Ti carbide having an average particle diameter of 5 nm or less is finely dispersed in the ferrite phase crystal grains. It has a tissue with a diameter of 1 μm or more. That is, the steel sheet of the present invention is characterized by having a ferrite single-phase metal structure and increasing the strength of the ferrite phase crystals with fine Ti carbides. Also, by not actively refining ferrite grains, the amount of refinement strengthening is made constant, and further the cause of segregation, that is, the size of carbides that cause strength fluctuations and the amount of precipitation that induces fluctuations in carbides. In order to eliminate the cause of the variation in the timing of precipitation, the amounts of Si and Mn are reduced. Thereby, strength fluctuation can be minimized by keeping the precipitation amount and size of Ti carbide in the steel plate constant, and the shape accuracy of the press-formed product is also improved.

Ferrite phase: Over 95% in area ratio (area ratio relative to the entire metal structure)
In the present invention, it is important that the metal structure of the hot-rolled steel sheet is a ferrite single phase. If the metal structure of a hot-rolled steel sheet is a dual-phase steel sheet containing a hard phase such as martensite or bainite in addition to the ferrite phase, the strength changes depending on the volume fraction of the hard phase, resulting in non-uniform steel sheet strength. To do. Further, in order to ensure the workability of the hot-rolled steel sheet (elongation flange workability, etc.), it is preferable that the metal structure is a ferrite single phase. However, even if the metal structure of the hot-rolled steel sheet is not a complete ferrite single phase, if the ferrite single phase is substantially the ferrite phase, that is, if the area ratio with respect to the entire metal structure is more than 95%, the above effect is sufficiently obtained. Can demonstrate. For this reason, in order to suppress fluctuations in strength, the metal structure is a ferrite phase with an area ratio exceeding 95%. Preferably it is 98% or more.

In the hot-rolled steel sheet of the present invention, examples of phases other than the ferrite phase include cementite, pearlite, bainite phase, martensite phase, residual austenite phase, etc., and these totals are acceptable if the area ratio is less than 5%. Is done. Preferably it is 2% or less. The metal structure here means a structure observed at a magnification of 100 to 5000 using an optical microscope or a scanning electron microscope.

Ti carbide Ti is a strong carbide-forming element, and carbides containing Ti tend to be fine carbides having an extremely small average particle size. Therefore, in the present invention in which the strength of the hot-rolled steel sheet is increased by dispersively precipitating fine carbide in the hot-rolled steel sheet, the fine carbide to be dispersed and precipitated is Ti carbide. Thus, according to the present invention utilizing precipitation strengthening, that is, control is facilitated by increasing the steel sheet strength only by carbide control, and stable strength can be obtained. Here, Ti carbide in the present invention is expressed by a chemical formula of TixMyCz (0 <x ≦ 1, 0 ≦ y <1, 0 <z ≦ 1, M: alloy element other than Ti; x + y ≦ 1). The carbide may contain carbide forming elements such as V and Mo other than Ti. However, y may be substantially zero.

Average particle size of Ti carbide: 5 nm or less The average particle size of Ti carbide that is dispersed and precipitated in the crystal grains of the ferrite phase is extremely important for imparting desired strength (tensile strength: 780 MPa or more) to hot-rolled steel sheets. In the present invention, the average particle size of Ti carbide is 5 nm or less. When fine carbide precipitates in the matrix, the fine carbide acts as a resistance to dislocation movement that occurs when deformation is applied to the steel sheet, thereby strengthening the hot-rolled steel sheet. Here, when the average particle diameter of the fine carbide is set to 5 nm or less, the above action becomes more remarkable. On the other hand, if the average particle diameter of the fine carbide exceeds 5 nm, it becomes difficult to ensure the strength of a 780 MPa grade steel sheet. Therefore, the average particle diameter of Ti carbide is 5 nm or less.

FIG. 1 shows a schematic shape of Ti carbide. In the case where the Ti carbide has a disk shape as shown in FIG. 1, assuming that the diameter of the disk is D and the thickness is t, the average particle diameter d of the Ti carbide is a value calculated by the following equation.
d = (D + t) / 2
On the other hand, when the Ti carbide is elliptical, the arithmetic average of the long axis and the short axis is taken as the average particle diameter of the Ti carbide. When the Ti carbide is spherical, the diameter of the sphere is the average particle diameter of the Ti carbide.

The atomic ratio of Ti and C contained in Ti carbide It is preferable that the atomic ratio of Ti and C contained in Ti carbide satisfies the following formula (1).
Ti / C <1.0 (1)
(Ti / C: atomic ratio of C and Ti in Ti carbide)
It is also effective to control the atomic ratio of Ti / C in Ti carbide in order to reduce the size of Ti carbide. By making Ti / C less than 1, Ti carbide with a size of 5 nm or less is stabilized. Obtained. The coarsening of Ti carbide is limited by the diffusion of Ti in the steel. That is, since the coarsening of Ti carbide is greatly affected by the amount of Ti solid solution, when Ti / C is 1 or more, Ti carbide tends to coarsen and stable strength may not be obtained. For this reason, it is preferable that the Ti / C atomic ratio is less than 1. The Ti / C atomic ratio can be controlled to a desired ratio by adjusting the steel plate composition and the manufacturing conditions of the steel plate. Conventionally, when Ti is added as a main carbide forming element, Ti / C may have exceeded 1.0 because it tends to be excessively added to C.

Average crystal grain size of ferrite phase: 1 μm or more Generally, the strength of a steel sheet improves when the crystal grains are refined. However, in the present invention, in order to stabilize the steel sheet strength, it is necessary to eliminate as much as possible elements that are strength variation factors other than precipitation strengthening. Here, when the ferrite average crystal grain size is less than 1 μm, the amount of fine grain strengthening increases rapidly, the strength greatly depends on the crystal grain size, and the strength becomes unstable. Therefore, in the present invention, the lower limit of the average crystal grain size of the ferrite phase is 1 μm. Preferably, it is 1.5 μm or more. On the other hand, if the average crystal grain size of the ferrite phase exceeds 10 μm, there is a concern about a decrease in toughness. Therefore, the average crystal grain size of the ferrite phase is preferably 10 μm or less.

Next, the reason for limiting the component composition of the hot-rolled steel sheet of the present invention will be described. In addition,% showing the following component composition shall mean the mass% unless there is particular notice.
C: Over 0.035% and below 0.065%
C is an essential element for forming Ti carbide in the steel sheet and increasing the tensile strength to 780 MPa or more. If the C content is 0.035% or less, a tensile strength of 780 MPa class cannot be realized. On the other hand, when the C content exceeds 0.065%, pearlite is easily generated, and the stability of strength deteriorates. In addition, stretch flangeability deteriorates due to the formation of pearlite. Therefore, the C content is more than 0.035% and not more than 0.065%. Preferably they are 0.04% or more and 0.06% or less. In order to precipitate Ti carbide that satisfies the above formula (1), the C content is preferably 0.04% or more and 0.065% or less.

Si: 0.2% or less
In conventional high-strength steel sheets, Si has been added as a solid solution strengthening element that increases strength while not decreasing elongation. However, since Si enhances hardenability and facilitates the formation of hard phases such as martensite phase and bainite phase, it inhibits the formation of a ferrite single phase structure. Therefore, the upper limit of Si content is 0.2%. Preferably, it is 0.1% or less. More preferably, it is 0.05% or less. There is no problem even if the Si content is zero.

Mn: 0.65% or less
Mn, like Si, has been positively added as a solid solution strengthening element in conventional high-strength steel sheets. However, Mn, like Si, enhances hardenability and facilitates the formation of hard phases such as martensite phase and bainite phase, and thus inhibits the formation of a ferrite single phase structure. When hard phases other than ferrite phase are mixed (more than 5% in area ratio), non-uniform steel sheet strength and deterioration of stretch flangeability are caused. Further, when a large amount of Mn is contained, segregation easily occurs, and this segregation partially lowers the ferrite transformation point. Here, precipitation of Ti carbide, which is the strengthening mechanism of the present invention, occurs simultaneously with the austenite → ferrite transformation. However, when the ferrite transformation point becomes non-uniform in the steel sheet as described above, the precipitation amount and size of Ti carbide also become non-uniform, resulting in a deterioration in strength stability. Therefore, the Mn content is set to 0.65% or less. Preferably it is 0.5% or less. There is no problem even if the Mn content is zero.

P: 0.03% or less
When the content of P exceeds 0.03%, segregation becomes prominent and inhibits fine precipitation of Ti carbide. Therefore, the P content is 0.03% or less. Preferably it is 0.02% or less, More preferably, it is 0.01% or less. There is no problem even if the P content is zero.

S: 0.02% or less S forms TiS in steel and causes strength fluctuations. In particular, TiS serves as a base point for fracture during stretch flange processing, thus lowering the tensile strength and causing fluctuations in strength. Therefore, in the present invention, it is preferable to reduce S as much as possible, and to be 0.02% or less. Preferably it is 0.005% or less, More preferably, it is 0.001% or less. There is no problem even if the S content is zero.

Al: 0.1% or less Al is an element that acts as a deoxidizer. In order to acquire such an effect, it is desirable to contain 0.01% or more, but when the content exceeds 0.1%, coarse alumina is formed, and the elongation flange workability deteriorates by becoming a starting point of fracture. Therefore, the Al content is 0.1% or less.

N: 0.01% or less N is a harmful element in the present invention and is preferably reduced as much as possible. N combines with Ti in the steel to form TiN. Here, when the N content exceeds 0.01%, the amount of coarse TiN increases, and the fluctuation of the tensile strength of the steel sheet increases particularly on the low strength side. Therefore, the N content is 0.01% or less. Preferably it is 0.006% or less. There is no problem even if the N content is zero.

Ti: 0.09% to 0.25%
Ti is an indispensable element for forming Ti carbide to increase the strength of steel, and is one of the most important elements in the present invention. When the Ti content is less than 0.09%, the precipitation amount of Ti carbide becomes insufficient, and it becomes difficult to obtain a desired steel plate strength (tensile strength of 780 MPa or more). On the other hand, if the Ti content exceeds 0.25%, solid solution Ti increases and the coarsening of Ti carbide cannot be suppressed, making it difficult to obtain the desired steel sheet strength (tensile strength of 780 MPa or more). Therefore, the Ti content is 0.09% or more and 0.25% or less.
In order to precipitate Ti carbide that satisfies the above formula (1), the Ti content is preferably 0.12% or more and 0.20% or less.

The above is the basic composition in the present invention. In addition to the basic composition described above, Cu, Ni, Cr, Co, Mo, Sb, W, As, Pb, Mg, Ca, Sn, Ta, Nb, V , REM, Cs, Zr, B, and Hf may be contained in total of 1% or less. If these contents are 1% or less in total, the above-described effects of the present invention are not affected. Components other than the above are Fe and inevitable impurities. For example, elements (Cu and the like) mixed from ore and scrap do not need to be reduced as long as they are below the total content. However, since relatively inexpensive Ti is used as the main carbide forming element in the present application, Mo, W, Nb, and V, which have a strong tendency to form carbide, may be added (content of impurities).

The steel sheet of the present invention may have a plating layer on the surface. By forming a plating layer on the surface of the steel sheet, the corrosion resistance of the hot-rolled steel sheet is improved, and a hot-rolled steel sheet suitable for a material for automobile parts exposed to severe corrosive environments can be obtained. Moreover, even if the surface of the steel sheet of the present invention is plated, the steel sheet characteristics of the present invention are not affected at all, and the above-described excellent effects of the present invention are still expressed. The type of the plating layer is not particularly limited, and electroplating or hot dipping may be used. If it is hot dip plating, hot dip galvanization is mentioned as a suitable example. It may be alloyed hot dip galvanized alloyed after plating. In addition, the plating layer of the present invention includes a pretreatment that is advantageous for chemical conversion treatment to disperse a metal or an oxide thereof on the surface.

Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the present invention, hot rolling consisting of rough rolling and finish rolling is applied to a steel material having the above composition, and after finishing rolling, the steel material is cooled, wound, and made into a hot-rolled steel sheet. At this time, the finish rolling temperature of the finish rolling is 840 ° C. or more and 1050 ° C. or less, the average cooling rate from the end of the cooling finish rolling to 750 ° C. is 30 ° C./s or more, and the winding temperature of the winding is 570 It is characterized by the temperature being not lower than ℃ and not higher than 750 ° C.

In the present invention, the melting method of the steel material is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed. Moreover, after melting, it is preferable to use a slab (steel material) by a continuous casting method because of problems such as segregation, but a slab can also be formed by a known casting method such as ingot-bundling rolling or thin slab continuous casting. good. In addition, when hot-rolling the slab after casting, the slab may be rolled after being reheated in a heating furnace, and when the temperature is maintained at a predetermined temperature or higher, direct rolling without heating the slab You may do it.

The steel material obtained as described above is subjected to heating, rough rolling and finish rolling. In the present invention, it is necessary to dissolve carbides in the steel material before rough rolling. In the present invention containing Ti which is a carbide forming element, the heating temperature of the steel material is preferably set to 1150 ° C. or higher. However, if the heating temperature of the steel material is excessively high, the surface is excessively oxidized and TiO ₂ is generated and Ti is consumed. Is preferably 1350 ° C. or lower. In addition, as described above, when the steel material before rough rolling maintains a temperature equal to or higher than a predetermined temperature, and the carbide in the steel material is dissolved, the step of heating the steel material before rough rolling is It can be omitted. The rough rolling conditions are not particularly limited.

Final rolling temperature: 840 ° C. or higher and 1050 ° C. or lower When the final rolling temperature exceeds 1050 ° C., the finally obtained ferrite crystal grains are easily coarsened more than necessary, and the steel sheet strength is significantly reduced. Accordingly, the finish rolling temperature is 1050 ° C. or lower. Preferably it is 980 degrees C or less. However, when the finish rolling temperature becomes extremely low, the average crystal grain size of the finally obtained ferrite becomes less than 1 μm. Furthermore, since there is concern about the generation of grains that have expanded in the rolling direction due to ferrite region rolling, the finish rolling temperature is set to 840 ° C. or higher. Preferably it is 880 degreeC or more.

Average cooling rate from finish rolling to 750 ° C: 30 ° C / s or more Optimizing the cooling rate can be achieved by changing the structure of the hot-rolled steel sheet over the entire length and width of the hot-rolled steel sheet. This is extremely important for obtaining a finely dispersed structure of Ti carbide having an average particle diameter of 5 nm or less in the crystal grains.

When the average cooling rate from the end of finish rolling to 750 ° C. is less than 30 ° C./s after the end of hot rolling, ferrite transformation starts in a temperature range higher than the coiling temperature described later. This makes it extremely difficult to achieve a desired steel sheet strength by uniformly dispersing and precipitating fine Ti carbide having an average particle diameter of 5 nm or less in the ferrite crystal grains in the longitudinal direction and the width direction of the steel sheet. Therefore, the average cooling rate from the end of finish rolling to 750 ° C. or lower is set to 30 ° C./s or higher. Preferably, it is 60 ° C./s or more.

Mn has the effect of shifting the nose of the ferrite transformation of steel to the long time side in the CCT diagram (continuous cooling transformation diagram). Therefore, when the Mn content in the steel is high, even if the cooling rate after the hot rolling is relatively slow (for example, about 10 to 30 ° C / s), the winding is performed before starting the ferrite transformation. It can be cooled to the coiling temperature and wound almost simultaneously with the ferrite transformation. However, as the Mn content in the steel decreases, the nose of the ferrite transformation of the steel in the CCT diagram shifts to the short time side. In other words, in the case of low-Mn steel, in order to cool to the coiling temperature without ferrite transformation after the hot rolling is completed and to wind the ferrite transformation almost simultaneously with the ferrite transformation, from the end of hot rolling to the cooling to the coiling temperature. It is necessary to shorten the time, that is, to increase the cooling rate after the hot rolling is completed.

Therefore, in the present invention, for example, when the Mn content is 0.5% or less, the average cooling rate from the end of finish rolling to 750 ° C. is preferably 50 ° C./s or more, and 100 ° C./s or more. More preferably.
However, if the average cooling rate is excessively large, temperature unevenness in the width direction of the steel sheet becomes remarkable, and mechanical characteristics in the width direction may become non-uniform, so the average cooling rate is 500 ° C./s or less. It is preferable to set it to 300 ° C./s or less.
Further, in order to precipitate Ti carbide satisfying the formula (1), the average cooling rate is preferably set to 60 ° C./s or more and 300 ° C./s or less.

Winding temperature: 570 ° C. or higher and 750 ° C. or lower When the winding temperature is lower than 570 ° C., bainitic ferrite and bainite are generated, making it difficult to make the metal structure substantially a ferrite single phase. Therefore, the coiling temperature is 570 ° C. or higher. On the other hand, when the coiling temperature exceeds 750 ° C., ferrite is easily obtained, but pearlite and coarse Ti carbide are generated and the strength is lowered. Therefore, the coiling temperature is 750 ° C. or lower. In order to more reliably suppress the formation of pearlite and coarse Ti carbide, the temperature is preferably 700 ° C. or lower. In order to precipitate Ti carbide that satisfies the above formula (1), it is more preferable that the coiling temperature be 600 ° C. or higher and 680 ° C. or lower. Further, by setting the coiling temperature to 570 ° C. or more in addition to the finish rolling temperature, the ferrite average crystal grain size can be made 1 μm or more.

As described above, according to the method of the present invention, the area ratio is more than 95% ferrite phase, Ti carbide having an average particle size of 5 nm or less is finely dispersed in the ferrite phase crystal grains, the ferrite phase A hot-rolled steel sheet having a structure with an average crystal grain size of 1 μm or more and a tensile strength of 780 MPa or more is obtained.

In addition, according to the present invention, the desired Ti carbide is finely dispersed in the longitudinal direction and the width direction of the steel sheet, particularly by reducing the Mn content and defining the cooling and winding conditions after completion of hot rolling. Therefore, a steel sheet excellent in material uniformity in which unevenness in strength is suppressed can be obtained. Specifically, for steel sheets with a plate width of 600 to 1600 mm, the difference in strength ΔTS between the tensile strength TSc at the center in the width direction of the steel sheet and the tensile strength TSe at a position 50 mm from the end in the width direction of the steel sheet is 20 MPa or less. A steel plate with a small thickness can be obtained. The steel sheet of the present invention is preferably 780 MPa class from the viewpoint of suppressing ΔTS.

In the present invention, a plated layer may be formed on the surface of the steel sheet by subjecting the hot-rolled steel sheet manufactured as described above to a plating process. The plating process may be either electroplating or hot dipping. For example, a hot dip galvanizing process may be performed as the plating process, or an alloying process may be further performed after the hot dip galvanizing process.

A hot-rolled steel sheet having a thickness of 2.6 mm and a width of 1200 mm was prepared by hot rolling a steel material (slab) having a thickness of 250 mm having the composition shown in Table 1. Table 2 shows the heating temperature, finish rolling temperature, average cooling rate from the end of finish rolling to 750 ° C., and the coiling temperature. Subsequently, after removing the scale of the surface layer by pickling, hot dip galvanizing treatment (plating bath composition: 0.1% Al-Zn, plating bath temperature: 470) is applied to some parts (hot rolled steel sheets No. 16, 17, 18). ° C) and further alloying treatment (alloying temperature: 520 ° C). The amount of plating adhered was 45 g / m ² per side.

Samples were taken from the hot-rolled steel sheet (hot-rolled steel sheet, alloyed hot-dip galvanized steel sheet) obtained as described above, and subjected to structure observation, tensile test, and hole expansion test (hole expanding test). The average crystal grain size, the average grain size of Ti carbide, the tensile strength, the hole expansion rate (stretch flangeability), and the strength difference between the central part and the end part in the width direction of the steel sheet were determined. The test method was as follows. The sampling position of the test piece was 20 m from the tail end (outer end) of the coil (however, excluding test piece sampling for investigating the fluctuation of mechanical characteristics in the coil longitudinal direction, which will be described later).

(1) Microstructure observation A specimen for microstructural observation is collected from the obtained hot-rolled steel sheet, polished on a cross section (L cross section) parallel to the rolling direction, corroded with nital, and optical microscope (magnification 500 times) And the structure | tissue was observed with the scanning electron microscope (magnification 3000 times), the structure | tissue was discriminate | determined, and the area ratio of structure | tissues other than a ferrite phase and a ferrite phase was calculated | required.
Further, a cross section parallel to the rolling direction was mirror-polished, corroded with a nital corrosive solution, ferrite grains were revealed, and the structure was imaged with an optical microscope (magnification: 100 times). In the obtained structure photograph, 10 straight lines were drawn in the rolling direction and the plate thickness direction at intervals of 100 μm or more, and the number of intersections between the grain boundaries and the straight lines was counted. The total length divided by the number of intersections was taken as the line segment length of one ferrite grain, and this was multiplied by 1.13 to determine the ASTM ferrite grain size.
Further, using a thin film collected from the steel sheet, 100 or more Ti carbides were observed with a transmission electron microscope (magnification of 340000 times), and the average particle diameter of Ti carbide was determined according to the above-mentioned rules.
The atomic ratio Ti / C of C and Ti in Ti carbides is obtained by collecting the extraction residue and using EDX (energy dispersive X-ray analysis) to determine the Ti concentration and EELS (electron energy loss spectroscopy) to determine the C concentration. Quantitatively calculated.

(2) Tensile test Tensile test specimens were collected from the obtained hot-rolled steel sheet, subjected to a tensile test in accordance with JIS Z2241, and measured for tensile properties (tensile strength TS). Tensile test specimens were sampled from the center of the width in the longitudinal direction and 50 mm from the end so that the rolling direction and the tensile direction were parallel, and the tensile strength of each test specimen was measured. Further, the difference ΔTS between the tensile strength TSc at the central part of the steel plate width and the tensile strength TSe at the position of 50 mm from the end of the steel plate width was determined. Table 3 shows the tensile strengths TSc and ΔTS at the center in the width direction. The magnitude relationship of tensile strength was all TSc> TSe.

(3) Hole expansion test A 130 mm square sample was cut out from the obtained hot-rolled steel sheet, a 10 mmφ hole was punched in the center with a clearance of 12.5%, and the hole was expanded with a conical punch with a vertex angle of 60 degrees from the punch side. The punch was stopped when a clear crack penetrating the plate thickness occurred, the specimen was taken out, and the diameter of the hole was measured. The difference between the hole diameter after the hole expansion and the hole diameter before the hole expansion was divided by the value before the hole expansion and multiplied by 100 to obtain the hole expansion ratio λ, which was obtained as an index of stretch flangeability.
The obtained results are shown in Table 3.

All of the examples of the present invention have high tensile strength TS: 780 MPa or more and stretch flangeability with a hole expansion ratio λ: more than 90%. The difference in tensile strength ΔTS at the 50 mm position is less than 20 MPa, and it is a hot-rolled steel sheet with small strength fluctuation. On the other hand, in a comparative example that is out of the scope of the present invention, a predetermined high strength cannot be secured, or the hole expansion ratio λ or the steel plate strength uniformity cannot be secured.

No. 4 and No. 7 hot-rolled steel sheets (coils) shown in Table 2 are each 40 m, 100 m, 300 m, 500 m from the longitudinal position shown in Table 4 (coil end (outside end of coil), At 700m), JIS No. 5 tensile test piece and hole expansion test piece were collected from the center in the width direction of the plate, and the tensile test and hole expansion test were carried out by the same method as (2) and (3) above. . Table 4 shows the obtained results. In addition, the difference ΔTS _L in tensile strength at each position in the longitudinal direction with respect to the 40 m position in the longitudinal direction is also shown. A sample having a difference ΔTS _L of less than 20 MPa was evaluated as good (◯).

In any of the coils shown in Table 4, the difference in tensile strength in the longitudinal direction ΔTS _L is less than 20 MPa, and the steel sheet is a hot-rolled steel sheet having a small strength fluctuation.

Claims

% By mass
C: more than 0.035% and 0.065% or less, Si: 0.2% or less,
Mn: 0.65% or less, P: 0.03% or less,
S: 0.02% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.09% or more and 0.25% or less, with the balance being composed of Fe and inevitable impurities, with an area ratio of more than 95% being the ferrite phase, and the ferrite phase crystal grains A high-strength hot-rolled steel sheet having a structure in which Ti carbide having an average particle diameter of 5 nm or less is finely dispersed, the ferrite phase has an average crystal grain diameter of 1 μm or more, and a tensile strength of 780 MPa or more.
The high-strength hot-rolled steel sheet according to claim 1, wherein the atomic ratio of C and Ti contained in the Ti carbide satisfies the following formula (1).
Record
Ti / C <1.0 (1)
(Ti / C: atomic ratio of C and Ti in Ti carbide)
In addition to the above composition, Cu, Ni, Cr, Co, Mo, Sb, W, As, Pb, Mg, Ca, Sn, Ta, Nb, V, REM, Cs, Zr, B, Hf The high-strength hot-rolled steel sheet according to claim 1 or 2, which contains a total of 1% or less of any one of the above.
The high-strength hot-rolled steel sheet according to any one of claims 1 to 3, wherein the steel sheet surface has a plating layer.
The high-strength hot-rolled steel sheet according to claim 4, wherein the plating layer is a zinc plating layer or a zinc-containing alloy plating layer.
The steel material is subjected to hot rolling consisting of rough rolling and finish rolling, and after finishing rolling, cooling, winding, hot rolling steel sheet,
The steel material in mass%,
C: more than 0.035% and 0.065% or less, Si: 0.2% or less,
Mn: 0.65% or less, P: 0.03% or less,
S: 0.02% or less, Al: 0.1% or less,
N: 0.01% or less, Ti: 0.09% or more and 0.25% or less, with the balance being Fe and inevitable impurities,
The finish rolling temperature of the finish rolling is 840 ° C. or more and 1050 ° C. or less, the average cooling rate from the end of the cooling finish rolling to 750 ° C. is 30 ° C./s or more, and the winding temperature of the winding is 570 ° C. or more and 750 ° C. A method for producing a high-strength hot-rolled steel sheet having a tensile strength of 780 MPa or more, characterized in that the temperature is not higher than ° C.
In addition to the above composition, Cu, Ni, Cr, Co, Mo, Sb, W, As, Pb, Mg, Ca, Sn, Ta, Nb, V, REM, Cs, Zr, B, Hf The manufacturing method of the high intensity | strength hot-rolled steel plate of Claim 6 which contains any 1 or more of these in total 1% or less.