JP5402847B2 - High-strength hot-rolled steel sheet excellent in burring properties and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength hot-rolled steel sheet excellent in burring properties and a method for producing the same.

  In recent years, for the purpose of improving the fuel efficiency of automobiles, in order to reduce the weight of various members, thinning by increasing the strength of steel plates such as iron alloys and the use of light metals such as Al alloys have been promoted. However, compared with heavy metals such as steel, light metals such as Al alloys have the advantage of high specific strength, but have the disadvantage of being extremely expensive, so their application is limited to special applications. Therefore, in order to promote the weight reduction of various members at a lower cost and in a wider range, it is necessary to reduce the thickness by increasing the strength of the steel sheet.

  Higher strength of steel sheets generally involves deterioration of material properties such as formability (workability), so how to increase strength without deteriorating material properties is important in the development of high strength steel plates. Become. In particular, steel plates used as automobile members such as inner plate members, structural members, and suspension members have stretch flangeability, burring workability, ductility, fatigue durability, impact resistance, corrosion resistance, etc., depending on their applications. Therefore, it is important to show how these material characteristics and high strength are balanced in a high dimension.

  For example, steel plates used for automobile members such as structural members and suspension members that account for approximately 20% of the weight of the vehicle body are mainly subjected to stretch flange processing and burring processing after blanking and punching by shearing and punching processing. Therefore, good hole expansibility (λ value) is required.

  In addition, steel sheets used for such members may have flaws and microcracks near the end surfaces formed by shearing and punching, and fatigue cracks may originate from these flaws and microcracks. It is feared that the fatigue endurance is reduced. In addition, there is a need to improve low temperature toughness in order to ensure impact resistance particularly in cold regions, even if the member is not easily destroyed even after being impacted by a collision or the like after being mounted on a vehicle as a part after molding. This low temperature toughness is specified by vTrs (Charpy fracture surface transition temperature) and the like. For this reason, in order to improve fatigue durability and impact resistance on the end face of the steel material, it is necessary not to cause wrinkles or microcracks and to consider the impact resistance itself.

  As wrinkles and minute cracks generated on these end faces, for example, as shown in Patent Document 1, a crack that occurs parallel to the thickness direction of the end faces is known. Hereinafter, this crack is referred to as “peeling”.

  Furthermore, as a steel plate used for automobile members such as a seat rail, a seat belt buckle, and a wheel disk, a high-strength steel plate excellent in design and high formability is required.

  Thus, in order to achieve both high strength and various material properties such as formability in particular, the steel structure is 90% or more ferrite, and the remainder is bainite, so that high strength, ductility, and hole expandability are achieved. A method for producing a compatible steel sheet is disclosed. (For example, see Patent Document 2.)

  However, the steel sheet manufactured by applying the technique disclosed in Patent Document 2 has verified that peeling occurs after punching, and there is a concern that fatigue cracks progress and fatigue durability deteriorates.

  Moreover, the technique of the high-tensile-strength hot-rolled steel sheet which is excellent in the stretch flangeability while being high intensity | strength by adding Mo and refine | miniaturizing a precipitate is disclosed. (For example, Patent Documents 3 and 4)

  However, the steel sheet to which the techniques disclosed in Patent Documents 3 and 4 described above are applied has a problem that the manufacturing cost is high because it is essential to add 0.07% or more of Mo which is an expensive alloy element. . In addition, the techniques disclosed in Patent Documents 3 and 4 do not disclose any technique for suppressing wrinkles and microcracks on end surfaces formed by shearing or punching.

  In addition, research to improve low-temperature toughness for steel sheets has been progressing in the past, but while maintaining high strength, it is possible to improve hole expansibility while suppressing peeling, and to achieve both low-temperature toughness. Further, the high-strength hot-rolled steel sheet having excellent burring properties is not particularly disclosed even with Patent Documents 1 to 4 described above.

WO / 2008/123366 JP-A-6-293910 JP 2002-322540 A JP 2002-322541 A

Metal Material Fatigue Design Handbook, Edited by Japan Society for Materials Science, page 84 “Refining Limit of Impurity Elements in Mass Production Scale” Japan Iron and Steel Institute High Temperature Scouring Process Group Scouring Forum Japan Society for the Promotion of Science, 19th Committee, Reaction Process Study Group, March 1996, pages 184-187

  Therefore, the present invention has been devised in view of the above-described problems, and the object of the present invention is to be applied to a member that requires high workability and hole expansibility while having high strength. High strength hot-rolled steel sheet with high burring properties, which is high in low-temperature toughness, and is a steel plate grade of 780 MPa class or higher, which is excellent in resistance to wrinkles and microcracks at the end face of a member formed by shearing or punching. And it aims at providing the manufacturing method which can manufacture the steel plate stably cheaply.

  In order to solve the problems as described above, the present inventors have invented a high-strength hot-rolled steel sheet having excellent burring properties shown below.

1)% by mass
C: 0.02 to 0.06%,
Si: 0.01 to 2.0%,
Mn: 0.7 to 2.3%
P: 0.1% or less,
S: 0.03% or less,
N: 0.02% or less Al: 0.001-1%,
Nb: 0.005 to 0.05%,
Ti: 0.03-0.17%,
B: 0.0002 to 0.002% is contained,
When Nb content is [Nb], Ti content is [Ti], N content is [N], S content is [S], C content is [C], and B content is [B]. Satisfies the following formula,
0.004 ≦ [C] +12/11 [B] −12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]),
[C] -12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]) ≦ 0.012
The balance consists of Fe and inevitable impurities,
The total grain boundary number density of solute C and solute B is more than 4.5 / nm ^{2 and} 12 / nm ² or less, and the cementite grain size precipitated at the grain boundaries in the steel sheet is 2 μm or less. Yes, the average grain size at the center of the plate thickness is 9 μm or less, and the {211} random strength ratio at the center of the plate thickness is 2 or less,
A high-strength hot-rolled steel sheet having excellent burring characteristics, wherein the average grain size of precipitates containing TiC in crystal grains is 3 nm or less and the density is 1 × 10 ¹⁶ pieces / cm ³ or more.

2) Furthermore, in mass%,
Cu: 0.02 to 1.2%,
Ni: 0.01 to 0.6%,
Mo: 0.01 to 1%,
V: 0.01-0.2%
Cr: 0.01-1%,
A high-strength hot-rolled steel sheet having excellent burring properties as described in 1), which comprises any one or more of the following.

3) Furthermore, in mass%,
Ca: 0.0005 to 0.005%,
REM: 0.0005 to 0.02%,
A high-strength hot-rolled steel sheet having excellent burring properties as described in 1) or 2), which comprises any one or two of the following.

4) Interspersed on a straight line in the rolling direction on a cross section having the width direction as a normal line and within 50 μm of each other and consisting of a collection of inclusions having an equivalent circle diameter of 3 μm or more and having a length of 30 μm or more 1 mm in cross section of the length of the object group and the length of the inclusion with a circle equivalent diameter of 3 μm or more, which is present at a position where there is no other inclusion within 50 μm on the straight line in the rolling direction, and is extended to 30 μm or more in the rolling direction ^The high-strength hot-rolled steel sheet having excellent burring properties according to any one of 1) to 3), wherein the total length per ² is 0.25 mm or less.

5) A steel slab having the component according to any one of 1) to 3) is heated to a temperature SRTmin (° C.) or higher and 1260 ° C. or lower satisfying the following formula:
SRTmin = 6670 / {2.26-log ([Nb] × [C])}-273
Furthermore, rough rolling is performed at a temperature of 1080 ° C. or higher and 1150 ° C. or lower, and the cumulative rolling reduction of the final stage of the rough rolling and the preceding stage is 40% or more and 65% or less,
Then, finish rolling is started at 1050 ° C. or more within 150 seconds, the final pass rolling reduction is over 15% and 25% or less, and finish rolling finish temperature FT is 848 + 2167 × [Nb] + 40353 × [B] ≦ FT ≦ 955 + 1389 × [Nb]
Finish rolling is finished in the temperature range to be cooled at a cooling rate of 15 ° C./sec or more, and the winding temperature CT is 8.12 × exp (4863 / (FT + 273)) with respect to the finish rolling finish temperature FT. A method for producing a high-strength hot-rolled steel sheet excellent in burring properties, characterized by satisfying the relationship of C ≦ CT ≦ 640 ° C. and winding up at 480 ° C. or more and less than 560 ° C.

  The present invention relates to a high-strength hot-rolled steel sheet excellent in burring properties and a method for producing the same, and by using these steel sheets, it can be easily applied to members that require strict workability and hole expansibility. In addition, these steel plates are highly resistant to high strength steel plates of 780 MPa class or higher that have excellent resistance to flaws and microcracks on the end faces of members formed by shearing and punching, and also have excellent low temperature toughness. Can be manufactured. For this reason, it can be said that this invention is an invention with high industrial value.

It is a figure which shows the presence or absence of the peeling in the relationship between the grain boundary segregation density of the solid solution C and the solid solution B, and coiling temperature. It is a figure which shows the relationship between a grain-boundary cementite particle size and a hole expansion value. It is a figure which shows the relationship between coiling temperature and a grain-boundary cementite particle size. It is a figure which shows the relationship between the X-ray random intensity ratio of a {211} surface parallel to the rolling surface of a steel plate, and a hole expansion rate. It is a figure which shows the relationship between average grain size and vTrs which is a parameter | index of low temperature toughness. It is a figure which shows the relationship between the size and density of a precipitate containing TiC, and tensile strength. It is a figure which shows the range of the finish rolling end temperature and coiling temperature of the range of this invention in the average particle diameter and density of the precipitate containing TiC. It is a figure which shows the relationship between finish rolling completion | finish temperature FT, the X-ray random intensity ratio of {211} surface, and an average crystal grain size.

  Hereinafter, as a form for carrying out the present invention, a high-strength hot-rolled steel sheet (hereinafter simply referred to as a hot-rolled steel sheet) excellent in burring properties will be described in detail. Hereinafter, mass% in the composition is simply referred to as%.

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

  First, slabs of steel components shown in Table 1 are melted, and in the manufacturing process of hot-rolled steel sheets, the hot rolling steel sheets are changed by changing the finishing rolling finish temperature and the coiling temperature, which have a great influence on the material of the hot-rolled steel sheets. Manufactured. Specifically, after hot rolling under conditions where the heating temperature is 1260 ° C. and the finish rolling end temperature is 750 to 1000 ° C., the steel is cooled at an average cooling rate of about 40 ° C./sec and wound at a temperature of 0 to 750 ° C. And a hot-rolled steel sheet having a thickness of 2.9 mm was manufactured, and various investigations were performed.

  In the following investigation, samples cut from 1/4 or 3/4 position of the steel plate width were used unless otherwise specified.

In the present specification, 1 ^* in the table is [C] +12/11 [B] -12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48 / 32 [S]), 2 ^* is [C] -12 / 48 × ([Ti] +48/93 [Nb] -48/14 [N] -48/32 [S]). In the formula, Nb content is [Nb], Ti content is [Ti], N content is [N], S content is [S], C content is [C], and B content is [B. ]. Here, steel A is steel to which Ti, Nb, and B are not added, steel B is steel to which B is not added, and steel C is steel to which B is added.

  The present inventors first examined conditions for suppressing peeling. According to the study by the present inventors, it has been clarified that the number of grain boundaries of solute C and solute B influences the occurrence of peeling. It is also known that the grain boundary number density of solute C and solute B is affected by the coiling temperature.

  Accordingly, the obtained hot-rolled steel sheet was examined for the presence or absence of fracture surface cracks in the relationship between the coiling temperature and the grain boundary segregation density of solute C and solute B.

  Here, in this investigation, the number density of the grain boundary of peeling, solute C, and solute B was evaluated according to the following method.

  The presence or absence of peeling was punched with a clearance of 20% by the same method as the hole expansion test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996, and the punched surface was visually confirmed.

  A three-dimensional atom probe method was used to measure the solid solution C and solid solution B existing in the grain boundaries and grains. The position sensitive atom probe (PoSAP) developed by A. Cerezo et al. At Oxford University in 1988 incorporates a position sensitive detector into the detector of the atom probe. It is a device that can simultaneously measure the flight time and position of atoms that have reached the detector without using an aperture for analysis. Using this device, not only can all the constituent elements in the alloy existing on the sample surface be displayed as a two-dimensional map with spatial resolution at the atomic level, but also the sample surface is evaporated one atomic layer at a time using the field evaporation phenomenon. By doing so, it is possible to display and analyze as a three-dimensional map by expanding the two-dimensional map in the depth direction. For grain boundary observation, an FIB (focused ion beam) apparatus / Hitachi FB2000A is used to produce a PoSAP needle-like sample including the grain boundary part, and the cut sample is arbitrarily formed into a needle shape by electrolytic polishing. The grain boundary portion was made to be the tip of the needle with a shape scanning beam. Taking advantage of the contrast between crystal grains having different orientations due to the channeling phenomenon of SIM (scanning ion microscope), the sample was observed and the grain boundary was identified and cut with an ion beam. The apparatus used as the three-dimensional atom probe is OTAP manufactured by CAMECA, and the measurement conditions are a sample position temperature of about 70 K, a probe total voltage of 10 to 15 kV, and a pulse ratio of 25%. Each sample was measured three times at the grain boundary and within the grain, and the average value was taken as the representative value.

The value obtained by removing background noise and the like from the measured value was defined as the atomic density per unit grain interface area, and this was defined as the grain boundary number density (grain boundary segregation density) (pieces / nm ² ). Therefore, the solid solution C existing at the grain boundary is exactly the C atom existing at the grain boundary, and the solid solution B existing at the grain boundary is the B atom existing exactly at the grain boundary.

  The total grain boundary number density of the solid solution C and the solid solution B in the present invention is defined as the density per unit area of the grain boundary of the solid solution C and the solid solution B existing at the grain boundary. This value is a value obtained by adding the measured values of solute C and solute B together.

  Since the atomic map shows the distribution of atoms in three dimensions, it can be confirmed that the number of C atoms and B atoms is large at the grain boundary positions. In addition, if it is a precipitate, it can be specified by the positional relationship between the number of atoms and other atoms (for example, Ti).

  The survey results are shown in FIG. FIG. 1 shows the presence or absence of peeling in the relationship between the total grain boundary density of solute C and solute B and the coiling temperature (CT). In FIG. 1, an unfilled plot (referred to as “Open” in the figure) indicates that no peeling occurred, and a filled plot (referred to as “Solid” in the figure) peels off. Has occurred.

From FIG. 1, it was found that when the grain boundary number density of solute C and solute B is more than 4.5 / nm ² , peeling can be suppressed. The reason why peeling occurred at 4.5 particles / nm ² or less is presumed to be because the grain boundary strength was relatively lowered as compared with that in the grains.

Next, in relation to the coiling temperature, the steel A to which Ti and Nb are not added has a grain boundary number density of solid solution C and solid solution B exceeding 4.5 / nm ² at any coiling temperature. Yes, peeling does not occur, but with steels B and C added with Ti and Nb, the grain boundary number density of solute C and solute B becomes 4.5 pieces / nm ² or less when the coiling temperature increases, and peeling occurs. did.

  This is because, in Steel A, Ti and Nb are not added, so that precipitation of TiC or the like does not occur even when the coiling temperature is high, and the grain boundary number density of solute C and solute B remains high. On the other hand, in steels B and C, when the coiling temperature increases, the solid solution C segregated at the grain boundaries mainly precipitates in the grains as TiC after winding, and the grain boundary number density of the solid solution C is low. Presumed to have decreased.

The grain boundary number density of more than 4.5 / nm ² is obtained with steel C up to a higher coiling temperature than steel B because B is added even if C precipitates in the grains as TiC. This is because the solid solution B segregates at the grain boundaries, thereby making it possible to compensate for the decrease of the solid solution C at the grain boundaries.

  Next, the present inventors examined conditions for improving hole expansibility. As a result of various investigations on the obtained steel sheet, it was found that the grain boundary cementite particle size and the texture were particularly affected by the hole expandability.

  Here, in this investigation, the hole expansion value, grain boundary cementite particle size, and texture were evaluated according to the following methods.

  A test piece having a length in the rolling direction of 150 mm and a length in the plate width direction of 150 mm was produced from the test steel, and the hole expansion value was evaluated according to the hole expansion test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996. In evaluating the hole expansion property, ten test pieces were manufactured from one test steel, and the measured values obtained by performing the hole expansion test on each of the manufactured test pieces were arithmetically averaged to obtain the hole expansion value. .

  The grain boundary cementite grain size precipitated at the grain boundaries was obtained by taking a transmission electron microscope sample from the 1/4 thickness of the sample cut from the 1/4 W or 3/4 W position of the steel plate width of the test steel. The observation was made with a transmission electron microscope equipped with a field emission gun (FEG) having an acceleration voltage of 200 kV. The precipitates observed at the grain boundaries were confirmed to be cementite by analyzing the diffraction pattern. In this study, the grain boundary cementite particle size is defined as an average value calculated from the measured value by measuring the grain boundary cementite as a circle-equivalent particle size by image processing or the like of all grain boundary cementite observed in one field of view.

  The texture was examined by X-ray diffraction. A specimen cut to 30 mmφ from the 1/4 W or 3/4 W position of the plate width was ground on a three-sided finish, and then the sample was subjected to X-ray diffraction by removing strain by chemical polishing or electrolytic polishing. The crystal orientation represented by {hkl} <uvw> indicates that the normal direction of the plate surface is parallel to <hkl> and the rolling direction is parallel to <uvw>. Measurement of crystal orientation by X-ray follows, for example, the method described in pages 274 to 296 of the new edition of Karity X-ray diffraction (published in 1986, translated by Gentaro Matsumura, Agne Co., Ltd.).

  The survey results are shown below. First, FIG. 2 shows the relationship between the grain boundary cementite particle size and the hole expansion value.

  In addition, it is recognized from FIG. 2 that there is a correlation between the hole expansion value and the grain boundary cementite particle size. When the grain boundary cementite particle size is smaller, the hole expansion value is improved, and when the grain boundary cementite particle size is 2 μm or less. It was newly discovered that a hole expansion value of 80% or more can be obtained.

  The reason why the hole expansion ratio is improved as the grain size of cementite existing at the grain boundary is smaller is considered to be as follows.

  First, it is considered that stretch flange processing and burring workability represented by the hole expansion value are affected by voids that are the starting points of cracks generated during punching or shearing.

  This void is considered to be generated when the cementite phase precipitated in the mother phase grain boundary is somewhat larger than the mother phase grain, because the mother phase grain near the interface of the mother phase grain receives excessive stress.

  However, when the grain boundary cementite particle size is small, the cementite grains are relatively small relative to the parent phase grains, so that the stress concentration is not mechanically concentrated and voids are less likely to be generated, which is thought to improve the hole expansion value. .

  FIG. 3 shows the relationship between the coiling temperature and the grain boundary cementite particle size.

  In either case, the grain boundary cementite particle size increases as the coiling temperature increases, but when the temperature exceeds a certain temperature, the grain boundary cementite particle size tends to decrease rapidly. In particular, in steels B and C containing Ti and Nb, the grain boundary cementite particle size is remarkably reduced, and when the coiling temperature is 480 ° C. or higher, it becomes 2 μm or less. This is considered as follows.

It is considered that there is a nose zone in the precipitation temperature of cementite in the α phase. It is known that this is expressed by a balance between nucleation with the driving force of supersaturation of C in the α phase and Fe ₃ C grain growth controlled by diffusion of C and Fe. When the coiling temperature is lower than this nose, the degree of supersaturation of C is large and the driving force for nucleation is large, but because of the low temperature, almost no diffusion is possible, and precipitation of cementite is suppressed not only at grain boundaries and within grains. Even if deposited, the size is small. On the other hand, when the coiling temperature is higher than the nose zone temperature, the solubility of C increases and the driving force for nucleation decreases, but the diffusion distance increases and the density decreases but the size tends to become coarser. Show. However, when elements that form carbides such as Ti and Nb are included, the precipitation nose in the α phase of Ti and Nb is on the higher temperature side than that of cementite. Both the precipitation amount and the size of the decrease.

  Next, the results of investigation on the influence of texture are shown. The relationship between the hole expansion value and the texture was arranged for the test steels of each steel number in which the finish rolling end temperature (FT) was changed at a constant coiling temperature. As a result, it has been found that the texture having a predetermined crystal orientation has an influence on the hole expandability of the hot rolled steel sheet obtained from the test steel. That is, as shown in FIG. 4, in a hot-rolled steel sheet, as the X-ray random intensity ratio (α {211} plane strength) of the {211} plane parallel to the rolled surface increases, the hole expandability deteriorates. found.

  The X-ray random intensity ratio here is the intensity ratio of the X-ray intensity of the hot-rolled steel sheet sample to be measured to the X-ray intensity of the powder sample having a random orientation distribution in the X-ray diffraction measurement. This means that the larger the random strength ratio, the greater the amount of texture having crystal planes in a predetermined orientation parallel to the plate surface in the steel plate.

As shown in FIG. 4, it can be seen that the smaller the {211} plane X-ray random intensity ratio, the better the hole expansion value. The reason why the hole expansion rate deteriorates when the α {211} plane strength is high is explained as follows. In a hot-rolled steel sheet, the anisotropy of the steel material increases due to the large number of α {211} surfaces. In particular, if the plastic strain ratios (r values) in the rolling direction and the 45 ° direction and the 90 ° direction (sheet width direction) with respect to the rolling direction are defined as r ₀ , r ₄₅ , and r ₉₀ , r _{0 in} this case. And r ₄₅ and r ₉₀ are increased, and r ₉₀ is greatly decreased. This is thought to be due to the fact that, during hole-opening forming, the reduction in plate thickness increases at the end surface in the rolling direction that receives tensile strain in the plate width direction, and high stress is generated on the end surface, so that cracks are easily generated and propagated. In addition, in Steel A, the hole expansion value is low despite the {211} plane X-ray random intensity ratio being 2 or less because the grain boundary cementite particle size is large.

  As shown in FIG. 8, it was discovered that the α {211} plane X-ray random intensity ratio decreases as the finish rolling finish temperature (FT) in the hot rolling process increases.

  These reasons are considered as follows. It is known that the {211} plane X-ray random intensity ratio represents the degree of accumulation of transformation textures from unrecrystallized austenite after hot rolling. From this, when the finish rolling finish temperature is high, the austenite recrystallization after finish finish finish is promoted, and this is considered to reduce the X-ray random intensity ratio of the {211} plane. .

  Next, the present inventors investigated low temperature toughness. The low temperature toughness was evaluated by vTrs (Charpy fracture surface transition temperature) obtained by the V-notch Charpy impact test. Here, in the V-notch Charpy impact test, a test piece was prepared based on JISZ2202, a Charpy impact test was performed on the test piece according to JISZ2242, and vTrs was measured.

In addition, since the average crystal grain size of the structure is greatly influenced by the low temperature toughness, the average crystal grain size at the center of the plate thickness was also measured. First, an EBSP-OIM ^™ (Electron Back Scatter Diffraction Pattern-Orientation Image Microscope) was used to measure the crystal grain size and microstructure from the cut microsample. 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 ^TM method irradiates a sample with high sensitivity in a scanning electron microscope (SEM) by irradiating an electron beam, photographing a Kikuchi pattern formed by backscattering with a high-sensitivity camera, and processing the computer image. It consists of a device and software that measure the crystal orientation of a point in a short 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 the present invention, the grain is visualized from an image mapped by defining the orientation difference of the crystal grain as 15 ° which is a threshold value of a large tilt grain boundary generally recognized as a crystal grain boundary, and the average crystal grain size Asked. Here, the “average crystal grain size” is a value obtained by EBSP-OIM ^™ .

  FIG. 5 shows the relationship between the average crystal grain size and vTrs. The lower the average crystal grain size, the lower the temperature of vTrs and the better the toughness. In the present invention, it can be seen that the average crystal grain size is 9 μm or less and vTrs is −20 ° C. or less, which is the target, and it can be used in cold regions.

  The average grain size directly related to low temperature toughness becomes finer as the finish rolling finish temperature is lower, and the low temperature toughness is improved. However, the X-ray random strength of the {211} plane, which is one of the dominant factors of the above-mentioned hole expansion value The ratio has an inverse correlation with the average grain size with respect to the finish rolling temperature, and no technology has been shown so far to achieve both the hole expansion value and the low temperature toughness.

  In order to make the average crystal grain size and the {211} plane X-ray random intensity ratio appropriate values and balance both the low temperature toughness and the hole expansion value in a more detailed study, the present inventors have described below. Thus, it has been newly found that the Nb amount, rough rolling temperature, final rolling step of rough rolling and the rolling reduction of the preceding stage, finish rolling end temperature, and final rolling final pass rolling reduction may be controlled.

  Next, the present inventors investigated the conditions for obtaining high strength. In the case of a hot-rolled steel sheet, the strength is generally increased by precipitation strengthening. In the case of a steel sheet containing Ti, the influence of precipitates containing TiC is large.

  In this investigation, among the steel components shown in Table 1 described above, for the test steel C, the relationship between the average particle diameter and density of precipitates containing TiC and the tensile strength was investigated. The manufacturing conditions are as described above.

  Here, the measurement of the TiC precipitate size and the TiC precipitate density was performed by the three-dimensional atom probe measurement method as follows.

  First, a needle-like sample is prepared from a sample to be measured by cutting and electrolytic polishing using a focused ion beam processing method in combination with an electrolytic polishing method as necessary. In the three-dimensional atom probe measurement, the accumulated data can be reconstructed and obtained as an actual distribution image of atoms in real space. The number density of TiC precipitates is determined from the volume of the three-dimensional distribution image of TiC precipitates and the number of TiC precipitates.

  The size of the TiC precipitate is the diameter calculated from the observed number of constituent atoms of the TiC precipitate and the lattice constant of TiC, assuming that the precipitate is spherical. The diameter of 30 or more TiC precipitates is arbitrarily measured, and the average value is obtained.

  Moreover, the tensile test of the hot-rolled sheet was performed according to the test method described in JIS Z 2241 by first processing the specimen into a No. 5 test piece described in JIS Z 2201.

FIG. 6 shows the relationship between the average particle diameter, density, and tensile strength of the precipitate containing TiC. If the components are constant, the average particle size and density of the precipitate containing TiC are almost inversely related. The number in the plot indicates the tensile strength (MPa). In order to obtain a tensile strength of 780 MPa or more, the average particle size of the precipitate containing TiC is 3 nm or less and the density is 1 × 10 ¹⁶ particles / cm 3. It turns out that it is necessary to be ³ or more.

  In other words, compared to the prior art, the average particle size of the precipitate containing TiC can be reduced while maintaining the tensile strength high, and the precipitate density can be controlled within a range that can be increased. It becomes.

  In the lower left plot of FIG. 6, the finish rolling finish temperature and the coiling temperature are both low, and precipitates containing TiC were hardly observed and the strength was low.

  FIG. 7 shows the range of the finish rolling end temperature and the coiling temperature within the range of the present invention in terms of the average particle diameter and density of the precipitate containing TiC.

In FIG. 7, “◯” means that the average particle diameter of the precipitate containing TiC is 3 nm or less and the density is 1 × 10 ¹⁶ particles / cm ³ or more, which means the scope of the present invention. “♦” means more than 3 nm or less than 1 × 10 ¹⁶ pieces / cm ³ . Further, “■” indicates a case where a precipitate containing TiC was not observed.

From FIG. 7, the average particle size of the precipitate containing TiC in the range of CT ≧ 8.12 × exp (4863 / (FT + 273)) and CT ≦ 650 ° C. is 3 nm or less and the density is 1 × 10 ¹⁶ particles / cm ^3. It became clear that this was the case. Moreover, if it was in this range, it was confirmed that the strength of the component of the present invention described later was 780 MPa or more and the precipitation strengthening ability was sufficiently exhibited. This mechanism for improving the strength is based on precipitation strengthening of TiC.

  Based on the knowledge obtained by the basic research as described above, the present inventors have intensively studied a hot-rolled steel sheet having a well-balanced tensile strength and hole-expandability and a manufacturing method thereof. The present inventors have come up with a hot-rolled steel sheet and a method for producing the same, which satisfy conditions.

  First, the reasons for limiting chemical components in the present invention will be described.

C: 0.02 to 0.06%
C segregates at the grain boundaries and has the effect of suppressing peeling at the end face formed by shearing or punching, and forms precipitates in the steel sheet by combining with Nb, Ti, etc. It contributes to strength improvement. In addition, it is an element that generates iron-based carbides such as cementite (Fe ₃ C), which is the starting point of cracks when expanding holes. If the C content is less than 0.02%, the strength cannot be improved by the precipitation strengthening and the effect of suppressing peeling cannot be obtained. If the C content exceeds 0.06%, the cementite becomes the starting point of cracking when the hole is expanded. Iron-based carbides such as (Fe ₃ C) increase and the hole expansion value deteriorates. For this reason, the C content is limited to a range of 0.02% to 0.06%. Further, considering the improvement in ductility as well as the improvement in strength, the C content is preferably 0.03 to 0.05%.

Si: 0.01 to 2.0%
Si is an element that contributes to an increase in the strength of the base metal, and also has a role as a deoxidizer for molten steel. The Si content exhibits the above effect when added in an amount of 0.01% or more, but even if added over 2.0%, the effect contributing to the increase in strength is saturated. For this reason, Si content was limited to the range of 0.01% or more and 2.0% or less. Addition of over 0.1% Si suppresses the precipitation of cementitious carbides such as cementite in the material structure as the content increases, and promotes the precipitation of fine carbide precipitates of Nb and Ti. It contributes to the improvement of strength and the ability to expand holes. Moreover, if this Si exceeds 1.0%, the effect of suppressing precipitation of iron-based carbide will be saturated. Therefore, the desirable range of Si content is more than 0.1 to 1.0%.

Mn: 0.7 to 2.3%
Mn is an element that contributes to strength improvement by solid solution strengthening and quenching strengthening. If the Mn content is less than 0.7%, this effect cannot be obtained, and even if added over 2.3%, this effect is saturated. For this reason, Mn content was limited to the range of 0.7% or more and 2.3% or less. In addition, when elements other than Mn are not sufficiently added to suppress the occurrence of hot cracking due to S, the Mn content ([Mn]) and the S content ([S]) are in mass% and [Mn It is desirable to add an amount of Mn such that] / [S] ≧ 20. Furthermore, Mn is an element that expands the austenite temperature to the low temperature side with an increase in the content thereof, improves the hardenability, and facilitates the formation of a continuous cooling transformation structure having excellent burring properties. Since this effect is hardly exhibited when the Mn content is less than 1.0%, it is desirable to add 1.0% or more. Further, if added over 1.6%, the austenite region temperature becomes too low, and it becomes difficult to obtain Nb and Ti carbides that are finely precipitated by ferrite transformation. Therefore, it is preferably 1.0% or more and 1.6% or less.

P: 0.1% or less P is an impurity contained in the hot metal, and is an element that segregates at the grain boundary and decreases toughness as the content increases. For this reason, the P content is preferably as low as possible. If the P content exceeds 0.1%, the workability and weldability are adversely affected. In particular, in consideration of hole expandability and weldability, the P content is preferably 0.02% or less.

S: 0.03% or less S is an impurity contained in the hot metal, and if the content is too large, not only causes cracking during hot rolling, but also an A-based inclusion that degrades hole expandability. It is an element to be generated. For this reason, the S content should be reduced as much as possible, but if it is 0.03% or less, it is an acceptable range, so it is 0.03% or less. However, the S content when a certain degree of hole expansibility is required is preferably 0.01% or less, more preferably 0.005% or less.

Al: 0.001 to 1%
Al needs to be added in an amount of 0.001% or more for molten steel deoxidation in the steel refining process, but the upper limit is set to 1% because of an increase in cost. Moreover, if adding too much Al, nonmetallic inclusions are increased and ductility and toughness are deteriorated, so 0.06% or less is desirable. More desirably, it is 0.04% or less. Moreover, in order to acquire the effect which suppresses precipitation of iron-type carbides, such as cementite, in material structure like Si, it is desirable to make it contain 0.016% or more. Therefore, it is more desirably 0.016% or more and 0.04% or less.

Nb: 0.005 to 0.05%
Nb is one of the most important elements in the present invention. Nb is finely precipitated as carbide during cooling after rolling or after winding, and improves strength by precipitation strengthening. Furthermore, Nb fixes C as a carbide and suppresses the formation of cementite, which is harmful to burring properties. Nb also exhibits a function of reducing the average crystal grain size of the steel sheet, thereby contributing to improvement of low temperature toughness. In order to obtain these effects, it is necessary to add at least 0.005% or more of Nb, and a more desirable addition amount is more than 0.01%. Further, by setting the lower limit of the Nb addition amount as high as 0.005%, the crystal grain size can be refined, and the degree of freedom in setting the rolling temperature can be improved without adversely affecting the low temperature toughness. On the other hand, when this Nb is added in excess of 0.05%, the temperature of the non-recrystallized region in the hot rolling process is expanded, and the non-recrystallized state that increases the X-ray random intensity ratio of the {211} plane Many rolling textures remain after the hot rolling process. For this reason, the Nb content is limited to 0.005% or more and 0.05% or less. A desirable content of Nb is 0.01% or more and 0.02% or less.

N: 0.02% or less N forms precipitates with Ti and Nb at a higher temperature than C, fixes Ti and reduces Ti and Nb effective for precipitation strengthening, thereby causing a decrease in tensile strength . Therefore, the N content should be reduced as much as possible, but is acceptable if it is 0.02% or less. Further, Ti and Nb nitrides precipitated at high temperatures are likely to be coarse and serve as a starting point for brittle fracture, thereby reducing low-temperature toughness. Further, from the viewpoint of aging resistance, 0.005% or less is more desirable.

Ti: 0.03-0.17%,
Ti is one of the most important elements in the present invention. During the cooling after rolling or during the γ → α transformation after winding, fine precipitates are formed as carbides, and the strength is improved by precipitation strengthening. Further, Ti fixes C as a carbide to TiC, and suppresses generation of cementite that is harmful to burring properties. In addition to this, Ti suppresses the precipitation of MnS forming stretched inclusions by precipitating as TiS during the heating of the steel slab in the hot rolling process, and reduces the total length M of the inclusions in the rolling direction. It is an element to be made. In order to obtain these effects, at least 0.03% or more of Ti should be added, and a more desirable content is 0.1% or more. On the other hand, even if added over 0.17%, these effects are saturated. For this reason, the Ti content is limited to 0.03% or more and 0.17% or less. The Ti content is more desirably 0.1% or more and 0.15% or less.

B: 0.0002 to 0.002%
B, like C, segregates at the grain boundary and is an element effective for increasing the grain boundary strength. That is, by causing segregation at the grain boundary as solid solution B together with solid solution C, it effectively works to prevent peeling. In order to exert such an effect, it is necessary to set the grain boundary number density of solute C and solute B in the range of 4.5 / nm ^{2 to} 12 / nm ^2. Even if it precipitates inside, the segregation of B at the grain boundary makes it possible to compensate for the decrease in the grain boundary of C. In order to compensate for the decrease in the grain boundary of C, if B is not added at least 0.0002%, the function of preventing peeling (breaking surface cracking) together with solid solution C cannot be exhibited. Further, when B is added in excess of 0.002%, it is an element that suppresses recrystallization of austenite during hot rolling in the same way as Nb, and strengthens the γ → α transformation texture from unrecrystallized austenite. . If the X-ray random intensity ratio of the {211} plane increases with this texture index, the hole expandability deteriorates. For this reason, content of B is made into 0.0002% or more and 0.002% or less. Further, B has an effect of improving the hardenability and facilitating the formation of a continuous cooling transformation structure which is a microstructure preferable for burring properties. In order to acquire the effect, it is desirable to contain 0.001% or more. On the other hand, B is an element in which slab cracking is a concern in the cooling step after continuous casting. From this viewpoint, its content is preferably 0.0015% or less. That is, it is desirably 0.001% or more and 0.0015% or less.

0.004 ≦ [C] +12/11 [B] −12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]) ... (1)
The right side of the formula (1) is an indicator of the amount of C and the amount of B that can be a solid solution C and a solid solution B affecting the occurrence of peeling, and is a front part defined as [C] +12/11 [B] And the rear part defined as 12/48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]) as a negative value.

  The former part is obtained by adding the C content to the C content in consideration of the B content and the atomic weight of B and C. On the other hand, the latter part represents the amount of C that can be precipitated as TiC, and is the amount of Ti that can be combined with C to precipitate TiC, multiplied by 12/48 times the atomic weight ratio of C and Ti. The amount obtained by subtracting the rear portion from the front portion is the amount of C and B that can be dissolved C and B.

  Ti is combined with N and S before being combined with C and precipitated as TiC and consumed as TiN and TiS. Further, when Nb is contained, Nb substitutes for Ti, bonds with C and N, and precipitates as TiNb (CN). Based on the above concept, the latter equation is arranged in consideration of the atomic weight of each element, and represents the amount of solid solution Ti that can be combined with C to become TiC.

  In the coiling temperature range defined in the present invention, if the right side of the mathematical formula (1) is less than 0.004 under the condition that addition of B is essential, the remaining solid solution C that does not bind to Ti is too small. Since it has been found that peeling cannot be prevented after punching, Formula (1) is defined.

[C] -12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]) ≦ 0.012 (2)
The left side of the above formula (2) indicates the amount of C that can remain as solute C after TiC precipitation. When this is 0.012 or less, the remaining C is an appropriate amount, so that the cementite particle size is 2 μm or less. However, when it exceeds 0.012, it has been found that the cementite particle size exceeds 2 μm and the hole expanding property is deteriorated. Therefore, Formula (2) is defined.

  The above is the reason for limiting the basic components of the present invention. In the present invention, Cu, Ni, Mo, V, Cr, Ca, and REM may be contained as necessary. The reasons for limiting the components of each element will be described below.

  Cu, Ni, Mo, V, and Cr are elements that have the effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and any one or two or more of these may be added. However, Cu content is less than 0.02%, Ni content is less than 0.01%, Mo content is less than 0.01%, V content is less than 0.01%, and Cr content is 0.01%. If it is less than the above, the above effect cannot be obtained sufficiently. Also, Cu content is over 1.2%, Ni content is over 0.6%, Mo content is over 1%, V content is over 0.2%, Cr content is over 1% Even so, the above effect is saturated and the economy is reduced. Therefore, when Cu, Ni, Mo, V, and Cr are included as necessary, the Cu content is 0.02% or more and 1.2% or less, and the Ni content is 0.01% or more and 0.6% or less. The Mo content is preferably 0.01% to 1%, the V content is 0.01% to 0.2%, and the Cr content is preferably 0.01% to 1%.

Ca and REM (rare earth elements) are elements that improve the workability by controlling the form of non-metallic inclusions that become the starting point of destruction and cause the workability to deteriorate. Even if the Ca and REM contents are added to less than 0.0005%, the above effects are not exhibited. Further, even if the Ca content exceeds 0.005% and the REM content exceeds 0.02%, the above effects are saturated and the economic efficiency is lowered. Therefore, it is desirable to add an amount of 0.0005% to 0.005% and a REM content of 0.0005% to 0.02%.
The hot-rolled steel sheet containing these as main components may contain Zr, Sn, Co, Zn, W, and Mg in total of 1% or less. However, Sn is preferably 0.05% or less because wrinkles may occur during hot rolling.

Next, metallurgical factors such as the microstructure in the hot rolled steel sheet to which the present invention is applied will be described in detail.
Since it is necessary to improve the grain boundary strength in order to suppress peeling that occurs during punching or shearing, the amount of solid solution C and solid solution B in the vicinity of the grain boundary that contributes to the improvement of the grain boundary strength as described above. Limit. When the grain boundary segregation density of the solute C and B is 4.5 pieces / nm ² or less, the above-described effect is not sufficiently exhibited, while when it exceeds 12 pieces / nm ² , the effect is saturated. Therefore, the grain boundary segregation density of the solid solution C and / or the solid solution B is 4.5 / nm ^{2 and} 12 / nm ² or less. In order to improve the grain boundary strength and more effectively suppress the peeling that occurs during punching or shearing, the lower limit of the grain boundary number density of the solid solution C and solid solution B is 5 pieces / nm ² or more. Desirably, the range is more preferably 6 / nm ² or more.

In the present invention, the grain boundary segregation density of solid solution C and solid solution B is the sum of the grain boundary segregation densities of solid solution C and solid solution B.
Stretch flange workability and burring workability typified by hole expansion values are affected by voids that are the starting points of cracks that occur during punching or shearing. Voids are generated when the cementite phase precipitated at the parent phase grain boundary has a certain size with respect to the parent phase grain, because the parent phase grain near the interface of the parent phase grain receives excessive stress concentration. However, when the cementite particle size is 2 μm or less, the cementite particles are relatively small with respect to the parent phase particles, the stress concentration is not mechanically concentrated, and voids are less likely to be generated, so that the hole expandability is improved. Therefore, the grain boundary cementite particle size is limited to 2 μm or less.

  Regarding the limitation of the X-ray random strength ratio of the {211} plane parallel to the rolling surface, as described above, it is found that the larger the X-ray random strength ratio of the {211} plane, the worse the hole expandability, and the tensile strength is 780 MPa. In order to obtain a target hole expansion value of 80% or more in the grade steel sheet grade, the X-ray random intensity ratio (α {211} plane strength) of the {211} plane parallel to the rolling surface is 2 or less. Turned out to be necessary.

The mechanism by which the hole expansion rate deteriorates when the α {211} plane strength is high is that the plastic strain ratio (r) in the 90 ° direction (sheet width direction) with respect to the rolling direction due to the large α {211} plane as described above. will be the value) r ₉₀ is greatly reduced, thickness reduction in the rolling direction end surface for receiving the tensile strain in the plate width direction during hole expansion molding is increased, cracks generated high stress is generated in the end face, easily propagated It is thought to be.

  Regarding the limitation of the average crystal grain size, a ductile-brittle transition temperature (vTrs) of at least −20 ° C. is obtained with a low temperature toughness evaluated by a V-notch Charpy impact test method defined by JISZ2242 using a test piece based on JISZ2202. For this purpose, the average crystal grain size must be 9 μm or less. This is because when the average crystal grain size exceeds 9 μm, the grain boundary that hinders cleavage fracture decreases, and vTrs may become a problem when used in cold regions, and may exceed −20 ° C., so the average crystal grain size is 9 μm. The following.

  The microstructure of the parent phase of the hot rolled steel sheet to which the present invention is applied is preferably a continuous cooling transformation structure (Zw). In addition, the microstructure of the matrix of the hot rolled steel sheet to which the present invention is applied has a volume fraction of polygonal ferrite (PF) of 20% or less in order to achieve both workability and ductility represented by uniform elongation. May be included. Incidentally, the volume fraction of the microstructure refers to the area fraction in the measurement visual field.

Here, the continuous cooling transformation structure (Zw) in the present invention is the Japan Iron and Steel Institute Basic Research Group, Bainite Research Group / Edition; Recent Research on Bainite Structure and Transformation Behavior of Low Carbon Steels-Final Report of Bainite Research Group -As described in (1994 Japan Iron and Steel Institute), the transformation structure is in the intermediate stage between the microstructure including polygonal ferrite and pearlite formed by the diffusion mechanism and the martensite formed by the non-diffusion and shear mechanism. A microstructure defined as That is, the continuous cooling transformation structure (Z _w ) is mainly a basic ferritic (α ° _B ), a granular bainitic ferrite (α _B ), as described in the above-mentioned references 125 to 127 as an optical microscope observation structure, It is composed of quasi-polygonal ferrite (α _q ), and is further defined as a microstructure containing a small amount of retained austenite (γ _r ) and martensite-austenite (MA). Note that α _q does not reveal an internal structure by etching like polygonal ferrite (PF), but the shape is ash and is clearly distinguished from PF. Here, α _q is a grain whose ratio (lq / dq) satisfies lq / dq ≧ 3.5 when the perimeter length lq of the target crystal grain and its equivalent circle diameter is dq. The continuous cooling transformation structure (Zw) in the present invention is defined as a microstructure containing one or more of α ° _B , α _B , α _q , γ _r and MA. A small amount of γ _r and MA is 3% or less in total.

This continuously cooled transformation structure (Zw) may be difficult to distinguish by optical microscope observation during etching using a Nital reagent. In that case, it is determined by using the EBSP-OIM ^TM.

The ^{EBSP-OIM TM (Electron Back Scatter} Diffraction Pattern-Orientation Image Microscopy) method, irradiated with an electron beam at a high slope sample in a scanning electron microscope (Scaninng Electron Microscope), was formed by backscatter Kikuchi pattern Is composed of a device and software that measures the crystal orientation of the irradiation point in a short time by taking a picture with a high-sensitivity camera and processing the computer image. The EBSP method can quantitatively analyze the microstructure and crystal orientation of the bulk sample surface, and the analysis area can be analyzed up to a minimum resolution of 20 nm as long as it is within the region that can be observed with the SEM, depending on the resolution of the SEM. The analysis by the EBSP-OIM ^TM method is performed by mapping several tens of thousands of points in an area of the grid to be analyzed over several hours. For polycrystalline materials, the crystal orientation distribution and crystal grain size in the sample can be seen. In the present invention, an image that can be discriminated from an image mapped with the azimuth difference of each packet as 15 ° may be conveniently defined as a continuous cooling transformation structure (Zw).

Furthermore, inclusions (hereinafter referred to as “stretched inclusions”) having a length in the rolling direction of 30 μm or more, and even if the length of one inclusion in the rolling direction is 30 μm or less, at intervals of 50 μm or less on the straight line in the rolling direction. It has been found that if there is a collection of inclusions arranged in a length of 30 μm or more (hereinafter referred to as inclusion group), the hole expansion value may be lowered. Therefore, it is preferable that the sum M of the lengths in the rolling direction per unit area of the stretched inclusions and the inclusion group is 0.25 mm / mm ² .

Here, the rolling direction length L1 of the inclusion group and the rolling direction length L2 of the extension inclusion are L1 (mm) for each inclusion group and extension inclusion for each visual field according to the following formula (3). And L2 (mm) are summed to obtain L (mm), and based on the obtained L, a numerical value M (mm / mm ² ) is obtained according to the following formula (4), and the obtained M is expressed as a unit area (1 mm ² ) It was defined as the sum M of the rolling inclusion lengths per hit. In the following formula (3), L1 _i and L2 _i are lengths in the rolling direction of each inclusion group and each stretched inclusion in one visual field, respectively, and S is the area (mm) of the observed visual field. ² ).

  In the above-described investigation of inclusions, a test piece was collected from a 300 mm portion at the center in the plate width direction, a cross section having the normal direction in the plate width direction (hereinafter referred to as L cross section) was mirror-polished, and an optical microscope Is used to observe the L cross section at a magnification of × 400.

  Next, the reason for limitation of the manufacturing method of the hot rolled steel sheet to which the present invention is applied will be described in detail below.

  In this invention, the manufacturing method of the steel slab which has the component mentioned above performed prior to a hot rolling process is not specifically limited. That is, as a method for producing a steel slab having the above-described components, following the smelting process using a blast furnace, converter, electric furnace, etc., the components are adjusted so that the desired component content is obtained in various secondary scouring processes, Next, the casting process may be performed by a method such as thin continuous slab casting, in addition to normal continuous casting or ingot casting. In addition, you may use a scrap for a raw material. When a slab is obtained by continuous casting, it may be sent directly to a hot rolling mill with a high-temperature slab, or may be hot-rolled after being reheated in a heating furnace after being cooled to room temperature.

  The slab obtained by the manufacturing method described above is heated in the heating furnace at a temperature equal to or higher than the minimum slab reheating temperature (= SRTmin) calculated based on the following formula (5) in the slab heating step before the hot rolling step. To do.

  SRTmin = 6670 / {2.26-log ([Nb] × [C])}-273 (5)

  Here, the content (%) of Nb is [Nb], the content (%) of C is [C], and SRTmin is the solution temperature of Nb and Ti carbonitrides from the product of Nb and C. It is what I have sought. The conditions for obtaining a TiNbCN composite prayer are determined by the amount of Ti. That is, when Ti is small, TiN alone does not precipitate.

  The reason why the tensile strength of the steel sheet is remarkably improved when the temperature is equal to or higher than the temperature SRTmin that satisfies the above formula (5) is as follows.

  That is, in order to obtain the target tensile strength, it is necessary to effectively utilize precipitation strengthening by Nb and Ti. These Nb and Ti are precipitated as coarse carbonitrides such as TiN, NbC, TiC, and NbTi (CN) in the slab pieces before heating. TiC is also almost dissolved at the solution temperature of Nb.

  This is because it is present in the slab as a composite precipitate of TiNbCN, and it has been found that the solution temperature is much lower than that of single Ti. In addition, in the case of Ti alone in the conventional knowledge, the solution formation becomes very high temperature, but in order to effectively obtain precipitation strengthening by Nb and Ti, these coarse carbonitrides are contained in the base material in the slab heating process. It is necessary to make a sufficient amount of solid solution. Most Nb and Ti carbonitrides dissolve at the solution temperature of Nb. Therefore, it was found that in the slab heating step, it is necessary to heat the slab to the solution temperature of Nb (= SRTmin) in order to obtain the target tensile strength.

  The literature values for normal solubility products are found in TiN, TiC, and NbN—NbC, respectively. Since TiN precipitation occurs at high temperatures, it has been considered difficult to dissolve by low-temperature heating as in the present invention. However, as described above, the inventors have found that most of TiC is dissolved substantially only by solution of NbC.

By observing a precipitate that seems to be a TiNb (CN) composite precipitate by observation with a transmission electron microscope replica, it was found that Ti, Nb, C, The concentration ratio of N is changing. That is, the concentration ratio of Ti and N is high in the central portion, whereas Nb and C are high in the shell portion. This is an MC type precipitate of TiNb (CN), and if it is NbC, Nb coordinates to Msite and C coordinates to Csite, but Nb is replaced by Ti depending on the temperature, This is because C is replaced with N. The same applies to TiN. Even if Nb is a temperature at which NbC completely dissolves, Nb is contained in TiN at a site fraction of 10 to 30%. Therefore, strictly speaking, Nb completely dissolves above the temperature at which TiN completely dissolves. However, in a component system in which the amount of Ti added is relatively small, this solution temperature may be used as the substantial lower limit temperature for dissolving Nb precipitates.

The same applies to TiC, in which Ti is coordinated to Msite, but at a low temperature, it is substituted with Nb at a certain rate. Therefore, the solution temperature of the TiNbCN composite precipitate may be a substantial solution temperature of TiC.

When the heating temperature is lower than SRTmin, the Nb and Ti carbonitrides are not sufficiently dissolved in the base material. In this case, the effect of improving the strength using precipitation strengthening cannot be obtained because Nb and Ti are finely precipitated as carbide during cooling after rolling or after winding. Therefore, the heating temperature in the slab heating process is set to be equal to or higher than SRTmin calculated by the above formula (5).
Further, if the heating temperature in the slab heating process is higher than 1260 ° C., the yield decreases due to the scale-off, so the heating temperature is set to 1260 ° C. or lower. Therefore, the heating temperature in this slab heating step is limited to the minimum slab reheating temperature calculated based on the above formula (5) and not higher than 1260 ° C. In addition, when the heating temperature is lower than 1080 ° C., the operation efficiency is remarkably impaired in the schedule, and therefore, the heating temperature is desirably 1080 ° C. or higher.

Further, the heating time in the slab heating step is not particularly defined, but in order to sufficiently dissolve the Nb and Ti carbonitrides, it is desirable to hold for 30 minutes or more after reaching the heating temperature described above. However, this is not the case when the cast slab is directly fed and rolled at a high temperature.
After the slab heating step, a rough bar is obtained by starting a rough rolling step for performing rough rolling on the slab extracted from the heating furnace without waiting. This rough rolling process is performed at a temperature of 1080 ° C. or higher and 1150 ° C. or lower for the reason described below. That is, when the rough rolling finish temperature is less than 1080 ° C., the X-ray random intensity ratio becomes large, and the hole expandability deteriorates. In addition, hot deformation resistance in rough rolling is increased, and there is a risk that the rough rolling operation may be hindered. On the other hand, when the rough rolling finish temperature is higher than 1150 ° C., not only will the average crystal grain size increase and vTrs decrease, but the secondary scale generated during the rough rolling will grow too much, which will be performed later. It may be difficult to remove the scale by descaling or finish rolling. Furthermore, when the rough rolling finish temperature is higher than 1150 ° C., inclusions may be stretched to cause deterioration of hole expansion property. Furthermore, if the rolling reduction ratio of the final stage of rough rolling and the preceding stage is less than 40%, the average crystal grain size is also increased, which causes a decrease in vTrs. On the other hand, if it exceeds 65%, the X-ray random intensity ratio increases, and the hole expandability deteriorates. Further, if it exceeds 65%, the inclusions may be stretched and cause the hole expandability to deteriorate.

  In addition, about the rough bar obtained after completion | finish of a rough rolling process, it joins each rough bar between a rough rolling process and a finish rolling process, and performs endless rolling which performs a finish rolling process continuously. Also good. 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.

  In addition, during the hot rolling process, it may be desired to control the variation in temperature in the rolling direction, the plate width direction, and the plate thickness direction of the rough bar to be small. In this case, if necessary, between the rough rolling mill in the rough rolling process and the finish rolling mill in the finish rolling process, or between each stand in the final rolling process, the rolling direction of the rough bar, the plate width direction, the plate The coarse bar may be heated by a heating device capable of controlling temperature variations in the thickness direction. Various heating means such as gas heating, energizing heating, induction heating, etc. can be considered as the heating device method, but if the variation in temperature in the rolling direction, plate width direction and plate thickness direction of the coarse bar can be controlled to be small. Any known means may be used. In addition, as a method of the heating device, an induction heating method with a good temperature control response industrially is preferable. If a plurality of transverse type induction heating devices that can be shifted in the plate width direction are installed even by the induction heating method, It is more preferable because the temperature distribution in the plate width direction can be arbitrarily controlled according to the width. Furthermore, as a heating apparatus, an apparatus constituted by a combination with a transverse induction heating apparatus and a solenoid induction heating apparatus that excels in overall plate width heating is most preferable.

  When temperature control is performed using these heating devices, it may be necessary to control the amount of heating by the heating device. In this case, since the temperature inside the coarse bar cannot be measured, the previously measured data such as the charging slab temperature, the slab in-furnace time, the heating furnace atmosphere temperature, the heating furnace extraction temperature, and the table roller transport time are used. Thus, it is desirable to estimate the temperature distribution in the rolling direction, the plate width direction, and the plate thickness direction when the coarse bar arrives at the heating device, and to control the heating amount by these heating devices.

  In addition, control of the heating amount by the induction heating apparatus is controlled as follows, for example. As a characteristic of the induction heating device (transverse induction heating device), when an alternating current is passed through the coil, a magnetic field is generated inside the coil. Then, an eddy current in the direction opposite to the coil current is generated in the circumferential direction perpendicular to the magnetic flux by the electromagnetic induction action in the conductor placed therein, and the conductor is heated by the Joule heat. Eddy currents are generated most strongly on the inner surface of the coil and decrease exponentially toward the inner side (this phenomenon is called the skin effect). Therefore, the smaller the frequency, the greater the current penetration depth, and a uniform heating pattern is obtained in the thickness direction. Conversely, the greater the frequency, the smaller the current penetration depth, and the overheating with the surface layer peaking in the thickness direction. It is known that a small heating pattern can be obtained. Therefore, with the transverse induction heating device, the heating of the rough bars in the rolling direction and the plate width direction can be performed in the same manner as in the past, and the heating in the plate thickness direction can be performed by changing the frequency of the transverse induction heating device. By varying the penetration depth and operating the heating temperature pattern in the thickness direction, the temperature distribution can be made uniform. In this case, it is preferable to use a variable frequency induction heating device, but the frequency may be changed by adjusting a condenser. In addition, in the control of the heating amount by the induction heating device, the distribution of the respective heating amounts may be changed so that a plurality of inductors having different frequencies are arranged to obtain a necessary thickness direction heating pattern. Furthermore, since the frequency varies when the air gap with the material to be heated is changed in the control of the heating amount by the induction heating device, the air gap may be changed to obtain a desired frequency and heating pattern.

  Further, the maximum height Ry of the steel sheet surface after finish rolling is desirably 15 μm (15 μm Ry, l2.5 mm, ln12.5 mm) or less. This is apparent from the fact that, as described in Non-Patent Document 1, the fatigue strength of a hot-rolled or pickled steel sheet correlates with the maximum height Ry of the steel sheet surface. In order to obtain this surface roughness, it is desirable to satisfy the condition of high-pressure water collision pressure P × flow rate L ≧ 0.003 on the steel plate surface in descaling. Further, the subsequent finish rolling is desirably performed within 5 seconds in order to prevent the scale from being generated again after descaling.

  After the rough rolling process is finished, the finish rolling process is started. Here, the time from the end of the rough rolling process to the start of the finish rolling process is 150 seconds or less.

When the time from the end of the rough rolling process to the start of the finish rolling process exceeds 150 seconds, Ti and Nb precipitate as coarse TiC and NbC carbides in the austenite in the coarse bar. Ti and Nb are elements that finely precipitate in ferrite during subsequent cooling or winding, and contribute to strength by precipitation strengthening. Therefore, when Ti and Nb are precipitated as carbides at this stage and the solid solution Ti and Nb are reduced, The strength improvement of the hot rolled steel sheet cannot be expected. In addition, due to the precipitation of coarse TiC and NbC, the absolute amount of solute C is insufficient in the hot coil which is one form as the final product of the hot rolled steel sheet, so the grain boundary segregation density of solute C and solute B is low. It becomes 4.5 pieces / nm ² or less, and peeling easily occurs.

In the finish rolling step, the finish rolling start temperature is set to 1050 ° C. or higher. When the finish rolling start temperature is less than 1050 ° C., the rolling temperature given to the rough bar to be rolled in each finish rolling pass tends to be lowered. Further, when the finish rolling start temperature is lower than 1050 ° C., the texture composed of a predetermined crystal orientation is affected. That is, when the finish rolling start temperature is less than 1050 ° C., the X-ray random intensity ratio is increased, and the hole expanding property is deteriorated. In this temperature range, with the decrease in the solid solubility limit of Nb and Ti, coarse TiC and NbC are likely to precipitate in austenite during finish rolling. Due to the precipitation of coarse TiC and NbC, the absolute amount of solute C is insufficient in a hot coil which is one form as a final product of a hot-rolled steel sheet, so the grain boundary segregation density of solute C and solute B is 4. Peeling is likely to occur at 5 pieces / nm ² or less.

  Thus, when coarse TiC and NbC are deposited in the finish rolling process, the strength of the steel sheet cannot be improved for the reasons described above, and peeling easily occurs. Therefore, in the finish rolling process, the finish rolling start temperature is set to 1050 ° C. or higher.

  The upper limit of the finish rolling start temperature is not particularly specified, but if it is 1150 ° C. or more, blisters that become the starting point of scale-like spindle scale defects occur between the steel plate and the surface scale before finish rolling and between passes. Therefore, it is desirable that the temperature is lower than 1150 ° C.

  Also, in the finish rolling process, if the rolling reduction of the final pass is 15% or less, austenite cannot be refined by recrystallization, and the average crystal grain size in the final product plate is 9 μm or less. Therefore, the desired low temperature toughness cannot be obtained.

  On the other hand, when the rolling reduction of the final pass exceeds 25%, the dislocation density inside the hot-rolled steel sheet increases more than necessary due to the introduction of excessive strain. After the finish rolling process, the region having a high dislocation density has a high strain energy, so it easily transforms into a ferrite structure, and the ferrite fraction exceeds 20%. Since the ferrite formed by such a transformation precipitates without dissolving so much carbon, the carbon contained in the mother layer tends to concentrate on the interface between austenite and ferrite, and coarse NbC, TiC is likely to precipitate.

  Thus, when coarse TiC and NbC are deposited in the finish rolling process, the strength of the steel sheet cannot be improved for the reasons described above, and peeling easily occurs. Therefore, the rolling reduction of the final pass in the finish rolling process is limited to more than 15% and 25% or less.

  Furthermore, the finish rolling finish temperature affects the texture and the average crystal grain size, but the affected temperature range varies depending on the contents of Nb and B and is clarified by the following formula (6). .

  848 + 2167 × [Nb] + 40353 × [B] ≦ FT ≦ 955 + 1389 × [Nb] (6)

  When the finish rolling finish temperature is less than the lower limit of the FT equation (6), which is a function of the Nb and B contents, the X-ray random intensity ratio (α {211} plane of the {211} plane parallel to the rolled plane (Strength) becomes more than 2, and the hole expandability deteriorates. As shown in FIG. 8, it is suggested that the X-ray random intensity ratio of the {211} plane decreases as the finish rolling finish temperature (FT) in the hot rolling process increases as shown in FIG. Yes. The reason why the formula (6) of the lower limit of FT depends on Nb and B is that both are recrystallization-suppressing elements in austenite, and the non-recrystallization rolling temperature range depends on the contents of these elements. On the other hand, when the finish rolling finish temperature is higher than the upper limit of the FT equation (6), which is a function of the Nb content, the average grain size after transformation becomes coarser, exceeds 9 μm, and deteriorates toughness. The reason why the formula depends on Nb for the upper limit of FT is that Nb is an element that suppresses the grain growth of recrystallized grains in the austenite recrystallization temperature range. For the same reason, the time from the end of finish rolling to the start of cooling is preferably within 10 seconds.

  In the present invention, the rolling speed is not particularly limited. However, if the rolling speed on the final finishing stand side is less than 400 mpm, the γ grains grow and become coarse, and the area where ferrite can be precipitated for obtaining ductility is reduced. The ductility may deteriorate. Moreover, although there is no particular limitation on the upper limit, the effect of the present invention can be obtained, but 1800 mpm or less is realistic due to equipment restrictions. Therefore, it is desirable that the rolling speed in the finish rolling process be 400 mpm or more and 1800 mpm or less as necessary.

After the finish rolling process is completed, a cooling process is performed in which the obtained steel sheet is cooled at a cooling rate of 15 ° C./sec or higher from the finish rolling finish temperature to the winding start temperature in the winding process described later for the following reason. That is, during the cooling from the end of the finish rolling process to the winding process, there is competition between the formation of cementite and precipitation nuclei such as TiC and NbC, and the formation of cementite precipitation nuclei when the cooling rate is less than 15 ° C./sec. Is prioritized, and in the subsequent winding process, it grows to a cementite of more than 2 μm at the grain boundary, and the hole expandability deteriorates. Moreover, there is a concern that the growth of cementite suppresses fine precipitation of carbides such as TiC and NbC, and the strength decreases. Furthermore, as will be described later, even when the coiling temperature is 650 ° C. or lower or 550 ° C. or lower, if the cooling rate is less than 15 ° C./sec, the growth to cementite is promoted, and solid solution C and / or B The grain boundary segregation density is less than 4.5 / nm ² , and peeling may occur. For this reason, the cooling rate was set to 15 ° C./sec. Note that the upper limit of the cooling rate in the cooling step is not particularly limited, but the effect of the present invention can be obtained.

In the cooling process, it is desirable that the microstructure is a continuous cooling transformation structure (Zw) in order to obtain better stretch flange processing and burring workability. The cooling rate for obtaining this microstructure is 15 ° C. / Sec or more is sufficient.
In other words, the region where the production can be stably performed is about 15 ° C./sec or more and 50 ° C./sec or less, and the region of 30 ° C./sec or less is a region where the production can be performed more stably as shown in the examples.

In the case where the microstructure is a continuous cooling transformation structure (Zw), a polygonal ferrite having a volume ratio of 20% or less is included as necessary for the purpose of improving ductility without significantly degrading burring properties. You may do it. In this case, in the cooling process from the end of the finish rolling process to the start of the winding process, the temperature range from the Ar ₃ transformation point temperature to the Ar ₁ transformation point temperature (two-phase region of ferrite and austenite) is 1 to 20 It may be allowed to stay for 2 seconds. The residence here is carried out in order to promote ferrite transformation in the two-phase region, but if it is less than 1 second, ferrite transformation in the two-phase region is insufficient, so that sufficient ductility cannot be obtained. There is a possibility that the size of the precipitate containing Ti and / or Nb becomes coarse and does not contribute to the strength due to precipitation strengthening. Accordingly, it is desirable that the residence time for the purpose of including polygonal ferrite in the continuous cooling transformation structure in the cooling step is 1 second or more and 20 seconds or less as necessary. In addition, the temperature range in which the residence is performed for 1 to 20 seconds is preferably not less than the Ar ₁ transformation point temperature and not more than 860 ° C. in order to facilitate the ferrite transformation. Further, the residence time of 1 to 20 seconds is more preferably 1 to 10 seconds so as not to extremely reduce productivity. Moreover, in order to satisfy these conditions, it is necessary to quickly reach the temperature range at a cooling rate of 20 ° C./sec or more after the finish rolling is finished. Although the upper limit of the cooling rate is not particularly defined, an appropriate cooling rate is 300 ° C./sec or less because of the capacity of the cooling facility. Furthermore, if this cooling rate is too fast, the cooling end temperature cannot be controlled, and overshooting may result in overcooling to below the Ar ₁ transformation point temperature, and the effect of improving ductility is lost. The cooling rate is desirably 150 ° C./sec or less.
Note that the Ar ₃ transformation point temperature, simply indicated in relation to the steel ingredients, for example, by the following calculation formula. That is, the Si content (%) is [Si], the Cr content (%) is [Cr], the Cu content (%) is [Cu], the Mo content (%) is [Mo], When the Ni content is [Ni], it is described as the following mathematical formula (7).
Ar ₃ = 910-310 × [C] + 25 × [Si] −80 × [Mneq] (7)

However, when B is not added, [Mneq] is expressed by the following mathematical formula (8).
[Mneq] = [Mn] + [Cr] + [Cu] + [Mo] + [Ni] / 2 + 10 ([Nb] −0.02) (8)

Or, when B is added, [Mneq] is expressed by the following mathematical formula (9).
[Mneq] = [Mn] + [Cr] + [Cu] + [Mo] + [Ni] / 2 + 10 ([Nb] −0.02) +1 (9)

In the winding process, when the winding temperature is lower than the lower limit temperature represented by the formula (10) as a function of FT, the average particle size required to obtain a tensile strength of 780 MPa or more is 3 nm or less and the density is A precipitate containing TiC that is 1 × 10 ¹⁶ pieces / cm ³ or more cannot be obtained.

  8.12 × exp (4863 / (FT + 273)) ° C. ≦ CT (10)

  On the other hand, when the coiling temperature is lower than 480 ° C., the particle size of cementite precipitated at the grain boundaries exceeds 2 μm, and the hole expansion value deteriorates.

On the other hand, if the coiling temperature is higher than 560 ° C., the grain boundary segregation density of the solid solution C and the solid solution B becomes 4.5 pieces / nm ² or less, and a fracture surface crack occurs. Therefore, the winding temperature CT in the winding process satisfies the relationship of 8.12 × exp (4863 / (FT + 273)) ° C. ≦ CT ≦ 560 ° C. and is limited to 480 ° C. or more and 560 ° C. or less.

In the present invention, this winding temperature is set to a temperature range of 480 ° C. or more and 560 ° C. or less. As a result, TiC hardly precipitates, and a state in which solute C and solute B are likely to remain at the grain boundary is created, so that the grain boundary number density of solute C and solute B is 4.5 / nm ^2. It becomes possible to be super. Accordingly, the lower limit of the formula (1) is set to 0.004.

  For the purpose of improving ductility by correcting the shape of the steel sheet and introducing movable dislocations, it is desirable to perform skin pass rolling with a rolling reduction of 0.1% or more and 2% or less after the completion of all the steps. Moreover, after completion | finish of all the processes, you may pickle with respect to the hot-rolled steel plate obtained as needed for the purpose of the removal of the scale adhering to the surface of the obtained hot-rolled steel plate. Furthermore, after pickling, the obtained hot-rolled steel sheet may be subjected to skin pass or cold rolling with a reduction rate of 10% or less inline or offline.

  Furthermore, 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.

  In addition, when galvanizing the hot-rolled steel plate after pickling, the obtained steel plate may be immersed in a galvanizing bath and may be alloyed as necessary. By performing the alloying treatment, the hot-rolled steel sheet is improved in resistance to various types of welding such as spot welding in addition to the improvement in corrosion resistance.

  The present invention will be further described below based on examples.

  The slabs of a to x having the chemical components shown in Table 2 are melted in a converter and secondary refining process, directly sent or reheated after continuous casting, and finished rolling following rough rolling to 2.0 to 3 The steel sheet was rolled down to a thickness of 6 mm, wound after cooling with a run-out table, and a hot-rolled steel sheet was produced. More specifically, hot-rolled steel sheets were produced according to the manufacturing conditions shown in Table 3. In addition, all the displays about the chemical composition in a table | surface are the mass%.

Furthermore, 1 ^* in the table is [C] +12/11 [B] -12 / 48 × ([Ti] +48/93 [Nb] -48/14 [N] -48/32 [S]), Among them, 2 ^* represents [C] -12 / 48 × ([Ti] +48/93 [Nb] -48/14 [N] -48/32 [S]). Moreover, the remainder of the component in Table 2 refers to Fe and inevitable impurities, and the underline in Table 2, Table 3, and Table 4 refers to outside the scope of the present invention.

Here, “component” refers to steel having a component corresponding to each symbol shown in Table 2, “solution temperature” refers to the minimum slab reheating temperature calculated by Equation (1), and “FT lower limit”. "and the lower limit of the finish rolling end temperature calculated by the equation (6), the upper limit of the finish rolling end temperature calculated by the formula as" FT limit "(6)," Ar ₃ transformation point “Temperature” means the temperature calculated by Equation (7), and “CT lower limit” means the lower limit value of the coiling temperature calculated by Equation (10). Further, the “heating temperature” is the heating temperature in the heating step, the “holding time” is the holding time at the predetermined heating temperature in the heating step, and the “rough rolling end temperature” is the temperature at which the rough rolling ends, “Rough cumulative rolling reduction” refers to the cumulative rolling reduction of the final stage of rough rolling and the preceding stage, and “Rough / finishing pass time” refers to the time from the end of the rough rolling process to the start of the finishing rolling process. “Temperature” refers to the temperature at which the finish rolling process begins. Furthermore, the “final final pass reduction ratio” is the reduction ratio in the final pass in the finish rolling process, and the “finish rolling end temperature” is the temperature at which the finish rolling process ends, “time to start cooling”. Is the time from the end of the finish rolling process to the start of cooling in the cooling process. The “cooling rate” is the average cooling rate from the start of the cooling process in the run-out table to the winding process, excluding the residence time. “Winding temperature” indicates a temperature at which the coiler winds in the winding process.

The evaluation method of the obtained steel plate is the same as that described above. Here, “grain boundary segregation density” means the number of atoms per grain interface area in the amount of solid solution C and solid solution B in the vicinity of the grain boundary, and “cementite diameter” means cementite precipitated at the grain boundary. The “average crystal grain size” is the average crystal grain size measured by EBSP-OIM ^™ , and the “{211} plane X-ray random intensity ratio” is the X of the {211} plane parallel to the rolling plane The line random intensity ratio (α {211} plane strength), “TiC size” means the average precipitate size of TiC (which may contain Nb and some N) measured by 3D-AP, and “TiC density” Is the average number per unit volume of TiC measured by 3D-AP, “Ferrite fraction” is the fraction of KAM ≦ 1 ° measured by EBSP-OIM ^™ , and “inclusions” are stretched inclusions The sum M of the rolling direction length per unit area of the inclusion group is shown.

  Further, the “tensile test” result shows the result of the C direction JIS No. 5 test piece. In the table, “YP” indicates the yield point, “TS” indicates the tensile strength, and “EI” indicates the elongation. “Hole expansion” results are the results obtained by the hole expansion test method described in JFS T 1001-1996, “Fracture surface cracking” results are the results of visually confirming the presence or absence, and there is no peeling. “None” is indicated, and “existing” is indicated when there is peeling. “Toughness” indicates the transition temperature obtained in the sub-size V-notch Charpy test.

In accordance with the present invention are 11 steels with steel numbers 9, 10, 11, 13, 25, 26, 27, 28, 29, 30, 47. These steel sheets contain a predetermined amount of steel components, have a cementite grain size of 2 μm or less precipitated at grain boundaries, and have a grain boundary segregation density of solute C and solute B of 4.5 / nm. ^{2 over} 12 particles / nm ² or less, the average crystal grain size at the center of the plate thickness is 9 μm or less, the X-ray random intensity ratio of the {211} plane parallel to the rolling surface is 2 or less, and A high-strength steel sheet having a grade of 780 MPa or more is obtained, in which the average particle size of the precipitate containing TiC is 3 nm or less and the density thereof is 1 × 10 ¹⁶ pieces / cm ³ or more.

  Steels other than the above are outside the scope of the present invention for the following reasons. That is, since the heating temperature of Steel No. 3 is outside the range of the method for producing a hot-rolled steel sheet of the present invention, the density of TiC is outside the range of the present invention and sufficient tensile strength is not obtained.

In Steel No. 1, the Nb content and the value of the formula 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet of the present invention, so the TiC density is outside the scope of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 2, the Nb content and the value of Formula 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet according to the present invention. Therefore, the TiC density is outside the scope of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 3, since the Nb content and the value of the formula 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, the TiC density is outside the range of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 4, since the Nb content and the value of the formula 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, the TiC density is outside the range of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 5, since the Nb content and the value of the formula 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, the TiC density is outside the range of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 6, the Nb content and the value of Formula 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet according to the present invention. Therefore, the TiC density is outside the scope of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 7, the Nb content and the value of Formula 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet according to the present invention. Therefore, the TiC density is outside the scope of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

In Steel No. 8, since the Nb content and the value of the formula 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, the TiC density is outside the range of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

  Steel No. 12 has a heating temperature outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the TiC size and density are outside the range of the present invention, and sufficient tensile strength is not obtained.

  Steel No. 14 has a rough rolling temperature outside the range of the method for producing a hot-rolled steel sheet according to the present invention, so that the {211} X random strength ratio is outside the range of the present invention and the hole expansion value is low.

  In Steel No. 15, the cumulative rolling reduction of rough rolling is outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the average crystal grain size is outside the range of the present invention and the low temperature toughness is poor.

  In Steel No. 16, the time between the roughing / finishing passes is outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the size and density of TiC are outside the range of the present invention and sufficient tensile strength is not obtained. . Further, the grain boundary segregation density is outside the range of the present invention, and peeling occurs.

  Steel No. 17 has a finish rolling start temperature outside the range of the method for producing a hot-rolled steel sheet according to the present invention. Therefore, the size and density of TiC are outside the range of the present invention, and sufficient tensile strength is not obtained.

  Steel No. 18 has a final final pass reduction ratio outside the lower limit range of the method for producing a hot-rolled steel sheet of the present invention, so the average crystal grain size is outside the range of the present invention and the low temperature toughness is poor.

  Steel No. 19 has a finishing final pass reduction ratio outside the upper limit range of the method for producing a hot-rolled steel sheet of the present invention, so that the grain boundary segregation density and the size of TiC are both outside the scope of the present invention, and peeling occurs. The tensile strength is not obtained.

  Steel No. 20 has a finish rolling finish temperature outside the upper limit range of the method for producing a hot-rolled steel sheet of the present invention, so the average crystal grain size is outside the range of the present invention and the low temperature toughness is poor.

  Steel No. 21 has a finish rolling finish temperature outside the lower limit range of the method for producing a hot-rolled steel sheet of the present invention, so that the {211} X random strength ratio is outside the range of the present invention and the hole expansion value is low.

  Steel No. 22 has a cooling rate outside the range of the method for producing a hot-rolled steel sheet according to the present invention, so that the grain boundary segregation density, cementite particle size, and the size and density of TiC are outside the range of the present invention. Peeling occurs and the hole expansion value and tensile strength are low.

  Steel No. 23 has a coiling temperature outside the upper limit range of the method for producing a hot-rolled steel sheet of the present invention. Therefore, the grain boundary segregation density and the size and density of TiC are outside the range of the present invention. Low strength.

  Steel No. 24 has a coiling temperature outside the lower limit of the method for producing a hot-rolled steel sheet of the present invention, so the cementite particle size and the size and density of TiC are outside the scope of the present invention, and the hole expansion value and tensile strength Is low.

In Steel No. 31, the C content and the value of the formula 1 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention. Therefore, the grain boundary segregation density is outside the range of the present invention, and peeling occurs. It is outside the scope of the present invention and has low strength.

In Steel No. 32, the content of Nb and C and the value of Formula 1 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention. Therefore, the grain boundary segregation density is outside the range of the present invention, and peeling occurs. TiC density is outside the range of the present invention, strength is low, average grain size is outside the range of the present invention, and low temperature toughness is poor.

In Steel No. 33, the content of C and the value of 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the cementite particle size is outside the range of the present invention and the hole expansion value is low.

Steel No. 34 has an Nb, C content and a value of 2 ^* outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the cementite particle size is outside the range of the present invention, the hole expansion value is low, and the average The crystal grain size is out of the range of the present invention and the low temperature toughness is poor.

  In Steel No. 35, the Nb content is outside the upper limit range of the method for producing a hot-rolled steel sheet of the present invention, so the {211} X random strength ratio is out of the range of the present invention, the ferrite fraction is high, The spread value is low.

  In Steel No. 36, the Nb content is outside the lower limit range of the method for producing a hot-rolled steel sheet of the present invention, so the average crystal grain size is outside the range of the present invention and the low temperature toughness is poor.

  In Steel No. 37, the content of Nb and Ti is outside the range of the production method of the hot-rolled steel sheet of the present invention. Therefore, the grain boundary segregation density is outside the range of the present invention, and peeling is low. It is out of the scope of the invention and the low temperature toughness is poor.

In Steel No. 38, the Ti content and 1 ^* are outside the scope of the method for producing a hot-rolled steel sheet of the present invention. Therefore, the grain boundary segregation density is outside the scope of the present invention, and peeling occurs, and the ferrite fraction is high. turn into.

In Steel No. 39, the content of Nb and Ti and the value of 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet of the present invention. Therefore, the size and density of TiC are outside the scope of the present invention, and the strength is low. The cementite particle size is outside the range of the present invention and the hole expansion value is low, the average crystal particle size is outside the range of the present invention, and the low temperature toughness is poor.

In Steel No. 40, since the Ti content and the value of 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet according to the present invention, the size and density of TiC are outside the scope of the present invention, and the strength is low. The particle size is outside the range of the present invention and the hole expansion value is low.

In Steel No. 41, since the Nb and B contents and the value of 2 ^* are outside the range of the method for producing a hot-rolled steel sheet of the present invention, the grain boundary segregation density is outside the range of the present invention, and peeling occurs. The particle size is outside the range of the present invention and the low temperature toughness is poor, the cementite particle size is outside the range of the present invention, and the hole expansion value is low.

  In Steel No. 42, the content of B is outside the range of the method for producing a hot-rolled steel sheet of the present invention, so that the grain boundary segregation density is outside the range of the present invention, and peeling occurs.

  In Steel No. 43, the Nb content is outside the range of the method for producing a hot-rolled steel sheet of the present invention, so the average crystal grain size is outside the range of the present invention and the low temperature toughness is poor.

In Steel No. 44, the value of 1 ^* is outside the range of the method for producing a hot-rolled steel sheet of the present invention, and therefore, the grain boundary segregation density is outside the range of the present invention and peeling occurs.

In Steel No. 45, the Nb content and the value of Formula 2 ^* are outside the scope of the method for producing a hot-rolled steel sheet according to the present invention. Therefore, the TiC density is outside the scope of the present invention, and the cementite particle size is Out of the scope of the invention, the hole expansion value and the tensile strength are low.

Steel No. 46 has a value of 2 ^* that is outside the upper limit of the method for producing a hot-rolled steel sheet of the present invention, so the cementite particle size is outside the range of the present invention and the hole expansion value is low.

Steel plates manufactured according to the present invention are strictly required for high strength and hole expansibility, including automobile parts such as inner plate members, structural members, and suspension members, shipbuilding, construction, bridges, offshore structures, pressure It can be used for all uses such as containers, line pipes, and machine parts.

Claims (5)

% By mass
C: 0.02 to 0.06%,
Si: 0.01 to 2.0%,
Mn: 0.7 to 2.3%
P: 0.1% or less,
S: 0.03% or less,
N: 0.02% or less,
Al: 0.001 to 1%,
Nb: 0.005 to 0.05%,
Ti: 0.03-0.17%,
B: 0.0002 to 0.002% is contained,
When Nb content is [Nb], Ti content is [Ti], N content is [N], S content is [S], C content is [C], and B content is [B]. Satisfies the following formula,
0.004 ≦ [C] +12/11 [B] −12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]),
[C] -12 / 48 × ([Ti] +48/93 [Nb] −48/14 [N] −48/32 [S]) ≦ 0.012
The balance consists of Fe and inevitable impurities,
The total grain boundary number density of solute C and solute B is more than 4.5 / nm ^{2 and} 12 / nm ² or less, and the cementite grain size precipitated at the grain boundaries in the steel sheet is 2 μm or less. Yes, the average grain size at the center of the plate thickness is 9 μm or less, and the {211} random strength ratio at the center of the plate thickness is 2 or less,
A high-strength hot-rolled steel sheet having excellent burring characteristics, wherein the average grain size of precipitates containing TiC in crystal grains is 3 nm or less and the density is 1 × 10 ¹⁶ pieces / cm ³ or more. In addition,
Cu: 0.02 to 1.2%,
Ni: 0.01 to 0.6%,
Mo: 0.01 to 1%,
V: 0.01-0.2%
Cr: 0.01-1%,
The high-strength hot-rolled steel sheet having excellent burring properties according to claim 1, comprising one or more of the above. In addition,
Ca: 0.0005 to 0.005%,
REM: 0.0005 to 0.02%,
The high-strength hot-rolled steel sheet having excellent burring properties according to claim 1, wherein the high-strength hot-rolled steel sheet has excellent burring properties. A group of inclusions consisting of a collection of inclusions that are within 50 μm of each other and arranged within a straight line in the rolling direction on a cross section having the width direction as a normal line, and that have an equivalent circle diameter of 3 μm or more and have a length of 30 μm or more Per 1 mm ² cross-section of the length of the inclusions and other inclusions within 50 μm on the straight line in the rolling direction, and the length of inclusions with a circle equivalent diameter of 3 μm or more extended to 30 μm or more in the rolling direction The high-strength hot-rolled steel sheet having excellent burring properties according to any one of claims 1 to 3, wherein the total length of the steel sheets is 0.25 mm or less. A steel slab having the component according to any one of claims 1 to 3 is heated to a temperature SRTmin (° C) of 1260 ° C or less that satisfies the following formula:
SRTmin = 6670 / {2.26-log ([Nb] × [C])}-273
Furthermore, rough rolling is performed at a temperature of 1080 ° C. or higher and 1150 ° C. or lower, and the cumulative rolling reduction of the final stage of the rough rolling and the preceding stage is 40% or more and 65% or less,
After that, finish rolling is started at 1050 ° C. or more within 150 seconds, the rolling reduction of the final pass is over 15% and 25% or less, and finish rolling finish temperature FT is 848 + 2167 × [Nb] + 40353 × [B] ≦ FT ≦ 955 + 1389 × [Nb]
Finish rolling is finished in the temperature range to be cooled at a cooling rate of 15 ° C./sec or more, and the winding temperature CT is 8.12 × exp (4863 / (FT + 273)) with respect to the finish rolling finish temperature FT. A method for producing a high-strength hot-rolled steel sheet excellent in burring properties, characterized by satisfying the relationship of C ≦ CT ≦ 560 ° C. and winding at 480 ° C. or more and 560 ° C. or less.

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