JP2016204690A - High strength hot rolled steel sheet excellent in ductility, fatigue characteristic and corrosion resistance and manufacturing method therefor - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high strength hot rolled steel sheet excellent in ductility, fatigue characteristic and corrosion resistance and a manufacturing method therefor, especially achieving both of fatigue characteristics and press moldability of a material without notch.SOLUTION: There is provided a high strength hot rolled steel sheet excellent in ductility, fatigue characteristic and corrosion resistance and having a predetermined component composition, containing a crystal particle with orientation difference of a grain boundary of neighboring crystal particle of 15° or more and average of orientation difference in the crystal particle of 0 to 0.5° of 50% by area percentage, total of martensite, tempered martensite and retained austenite is 2% to 10% by area fraction, further Ti of 40 mass% or more of Tie represented by the (a) formula exists as Ti carbide and mass of the Ti carbide having circle equivalent particle diameter of 7 nm to 20 nm is 50% or more of mass of all Ti carbide. Tief=[Ti]-48/14×[N]-48/32×[S] (a)SELECTED DRAWING: None

Description

  The present invention relates to a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics, and corrosion resistance, and a method for producing the same.

  In recent years, various parts constituting an automobile have been reduced in weight for the purpose of improving the fuel efficiency of the automobile. The weight reduction means differs depending on the required performance of each part. For example, in the framework parts, the steel sheet is thinned by increasing the strength, and in the panel parts, the steel sheet is replaced with a light metal such as an Al alloy. However, in comparison with steel, light metals such as Al alloys are expensive and are currently applied mainly to luxury cars.

  On the other hand, automobile demand is shifting from developed countries to emerging countries, and it is expected that both weight reduction and price reduction will be required in the future. Regardless of the parts, it is necessary to achieve high weight and light weight by thinning the steel.

  From the viewpoint of design, aluminum wheels and forged products are advantageous for passenger car wheels. Recently, however, products with the same design characteristics as aluminum wheels have emerged for steel press products by devising materials and construction methods.

  In particular, steel wheels are required to have the same design and beauty as aluminum wheels, in addition to the excellent fatigue durability and corrosion resistance that have been required so far for wheel discs that can be seen by end users. Similarly, in steel plates for wheel disks, in addition to increasing strength to achieve thinning, fatigue resistance and corrosion resistance, and improving workability and aesthetics to improve design as a part. Improvements in surface properties to ensure security have been demanded.

  Up to now, emphasis workability, drawing workability and fatigue durability have been regarded as particularly important characteristics required for steel plates for wheel disks. This is because the hat part is severely processed in the wheel disk forming process, and the fatigue durability is controlled by the strictest standard in the wheel member characteristics. Fatigue durability is classified into two types: fatigue properties of smooth materials without notches, and notch fatigue properties under stress concentration, and it is desirable that both fatigue properties be good in a wheel disc.

  Currently, as a high-strength hot-rolled steel sheet for these wheel disks, a 590 MPa class ferritic-martensitic composite structure steel sheet (so-called dual phase steel) is used, which emphasizes fatigue durability in the member and is excellent in fatigue characteristics. . However, the strength level required for these steel sheets is moving from the 590 MPa class to the 780 MPa class to further increase the strength.

  In Non-Patent Document 1, uniform elongation is ensured even with the same strength by forming a microstructure of a steel sheet as a dual phase steel (hereinafter referred to as DP steel) composed of ferrite and martensite. A method is disclosed.

  On the other hand, it is known that DP steel has low local deformability represented by bending, hole expanding and burring. This is because, since the strength difference between ferrite and martensite is large, a large strain and stress concentration are generated in the ferrite in the vicinity of martensite, and cracks occur. Based on this knowledge, high-strength steel sheets with increased hole expansion rate by reducing the difference in strength between tissues have been developed.

  Patent Document 1 proposes a steel sheet that has bainite or bainitic ferrite as a main phase to ensure strength and greatly improve hole expansibility. By using single structure steel, the strain and stress concentration as described above do not occur, and a high hole expansion rate can be obtained.

  However, since the single-structure steel of bainite or bainitic ferrite is used, the elongation is greatly deteriorated, and it is not possible to achieve both the elongation and the hole expandability.

  Furthermore, in recent years, high strength steel sheets have been proposed in which ferrite having excellent elongation is used as the structure of a single structure steel, and the strength is increased using carbides such as Ti and Mo (for example, Patent Documents 2 to 2). 4).

The steel sheet proposed in Patent Document 2 contains a large amount of Mo.
The steel sheet proposed in Patent Document 3 contains a large amount of V.
The steel sheet proposed in Patent Document 4 needs to be cooled during rolling in order to refine crystal grains.
Therefore, these steel plates have a problem that alloy costs and manufacturing costs are increased. Also, in these steel plates, the elongation has deteriorated due to the fact that the ferrite itself is greatly increased in strength. Although these steel sheets exceeded the elongation of single-structure steels such as bainite and bainitic ferrite, the stretch-hole expanding balance was not always sufficient.

  Patent Document 5 proposes a composite-structure steel plate in which martensite in DP steel is bainite and the hole expandability is improved by reducing the difference in strength between the microstructures with ferrite. However, as a result of increasing the area ratio of the bainite structure in order to ensure the strength, the elongation deteriorated and the stretch-hole expanding property balance was not sufficient.

  Patent Documents 7 to 9 propose steel sheets in which the strength difference from the hard structure is reduced by precipitation strengthening ferrite of DP steel.

  However, in this technique, Mo is an essential element, and there is a problem that the manufacturing cost increases. Moreover, even if ferrite is strengthened by precipitation, the difference in strength from martensite, which is a hard structure, is large, and a high hole expandability improvement effect is not obtained.

  On the other hand, since these DP steels have a microstructure of a composite structure of ferrite and martensite, Si is often added for the purpose of promoting ferrite transformation. However, when Si is contained, a tiger stripe-like scale pattern called red scale (Si scale) is generated on the surface of the steel sheet, and therefore it is difficult to apply it to various steel sheets used for high-design wheel disks that require aesthetics.

  Patent Document 10 discloses a technique relating to a steel plate having an excellent balance between elongation and hole-expandability by controlling the martensite fraction of DP steel to 3 to 10% in a steel plate of 780 MPa or more. However, since Si is added in an amount of 0.5% or more and it is difficult to avoid the Si scale pattern, it is difficult to apply it to various steel plates used for high-design wheel disks that require aesthetics.

  Patent Documents 11 and 12 address this problem by suppressing the generation of red scale by suppressing the addition amount of Si to 0.3% or less, and further adding Mo to refine the precipitates. In addition, a technique of a high-tensile hot-rolled steel sheet having high strength and excellent stretch flangeability is disclosed. However, the steel sheet to which the techniques disclosed in Patent Documents 11 and 12 described above are applied has a Si addition amount of about 0.3% or less, but it is difficult to sufficiently suppress the occurrence of red scale, and is expensive. Since it is essential to add Mo which is an alloying element in an amount of 0.07% or more, there is a problem that the manufacturing cost is high.

  Patent Document 13 discloses a technique for improving the low cycle fatigue characteristics by adding Al. However, there is no technical disclosure about notch fatigue characteristics that are fatigue characteristics under stress concentration.

  Patent Document 14 discloses a technique for achieving both suppression of red scale and notch fatigue characteristics by defining the upper limit of the Si content and controlling the precipitate density and the hard phase fraction. However, there is no technical disclosure about the fatigue characteristics of a material without a notch (smooth material).

In Patent Document 15 and Patent Document 16, since it is effective to refine the structure in order to improve the fatigue characteristics of a material without a notch, the average particle diameter of ferrite is 2 μm or less. A hot-rolled steel sheet having a good balance and fatigue limit ratio (fatigue strength / TS) is described.
Patent Document 17 describes that since a fatigue crack is generated near the surface, the structure near the surface is refined.
Patent Document 18 describes an improvement in fatigue characteristics by refining the martensite structure. However, refinement improves the fatigue properties of materials without notches and does not impair ductility, but the strength increase is small compared to strengthening mechanisms such as solid solution strengthening and precipitation strengthening. There is a limit to guaranteeing fatigue characteristics only by making finer as the process progresses.

Solid solution strengthening and precipitation strengthening generally give a large increase in strength and are applied to high strength steel sheets.
In Non-Patent Document 1, the effects of solid solution strengthening / precipitation strengthening / fine grain strengthening on tensile strength and fatigue properties are investigated, and the amount of increase in fatigue strength relative to the amount of increase in tensile strength is determined by solid solution strengthening> precipitation strengthening> It is reported that the order of fine grain strengthening.
Non-Patent Document 2 reports that fatigue characteristics deteriorate when the state of Cu in steel is changed from solid solution to precipitation.

  Although precipitation strengthening has a higher tensile strength per added amount than solid solution strengthening and can easily increase the strength, there is a report that the amount of increase in fatigue characteristics is small compared to solid solution strengthening as described above. For this reason, solid solution strengthened steel sheets have been preferentially used for parts where fatigue characteristics are important.

For improvement of notch fatigue properties, it has been reported that reduction of crack propagation rate by compounding is effective.
In Patent Document 19, hard bainite or martensite is dispersed in a structure having fine ferrite as a main phase, thereby achieving both fatigue characteristics and notch fatigue characteristics of a material having no notches. However, no method for improving the press formability is described, and no special attention is paid to the hardness and shape of bainite and martensite. Therefore, it is considered that the press formability is not provided.

In Patent Document 20 and Patent Document 21, it is reported that the crack propagation speed can be reduced by increasing the aspect ratio of martensite in the composite structure. However, both are technologies applied to thick plates and do not have workability such as ductility and hole expansibility necessary for press forming. For this reason, it is difficult to use the steel sheets described in Patent Document 20 and Patent Document 21 as automobile steel sheets.
Thus, the fatigue properties of materials without cracks are a problem in precipitation-strengthened steel, and not only the precipitation-strengthened steel, but also the compatibility between the notch fatigue properties, fatigue properties of smooth materials without notches and press formability, No effective solution has been proposed.

JP 2003-193190 A JP 2003-089848 A JP 2007-063668 A JP 2004-143518 A JP 2004-204326 A JP 2007-302918 A JP 2003-321737 A JP 2003-321738 A JP 2003-321739 A JP 2011-184788 A JP 2002-322540 A JP 2002-322541 A JP 2010-150581 A International Publication No. 2014/051005 Japanese Patent Laid-Open No. 11-92859 Japanese Patent Laid-Open No. 11-152544 JP 2004-2111199 A JP 2010-70789 A JP 04-337026 A JP 2005-320619 A Japanese Patent Application Laid-Open No. 07-90478

Takashi Abe et al .: Iron and Steel, 70th (1984) No. 10, page 145 T.A. Yokoi et al .: Journal of Materials science, 36th (2001), p. 5757. Masaya Misui et al .: CAMP ISIJ, 5th Year (1992), p. 1867

  The present invention has been devised in view of the above-described problems, and an object thereof is to provide a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics, and corrosion resistance, and a method for producing the same. In particular, it is an object to satisfy both fatigue characteristics and press formability of a material having no notch.

In addition, the ductility defined here is a uniform elongation necessary for press formability. If the uniform elongation is small, a reduction in the plate thickness due to necking is likely to occur during press molding, causing press cracks. Further, the uniform elongation is a plastic elongation (%) at the maximum test force of JIS Z 2241: 2011, and may be described as (u-El) (uniform elongation).
The uniform elongation (u-El) needs to satisfy (TS) × (u-El) ≧ 8000 MPa% with tensile strength (TS) ≧ 540 MPa in order to ensure press formability. Also, the hole expansion rate (λ) by the test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996 is such that when the undercarriage part is pressed, the stretch flangeability is insufficient and cracks do not occur. It is necessary to satisfy (TS) × (λ) ≧ 36000 MPa%.
Here, (TS) represents the tensile strength measured based on JIS Z 2241: 2011.

  As a result of intensive studies, the present inventors have optimized the components and production conditions of a high-strength hot-rolled steel sheet and controlled the structure of the steel sheet, thereby improving the high-strength hot-rolled steel sheet having excellent ductility, fatigue properties, and corrosion resistance. Was successfully manufactured. The summary is as follows.

(1)
Chemical composition is mass%,
C: 0.030 to 0.200%,
Mn: 0.10 to 3.00%,
P: 0.100% or less,
S: 0.0300% or less,
Al: 0.100 to 2.000%
N: 0.0100% or less (excluding 0),
O: 0.0100% or less,
Ti: 0.010 to 0.380%,
The balance is Fe and inevitable impurities, and Tief represented by the formula (a) is 0.01 to 0.30%,
A crystal grain having an orientation difference between adjacent grain boundaries of 15 ° or more, wherein the crystal grains have an average orientation difference of 0 to 0.5 ° in the crystal grain and include an area ratio of 50% or more. Furthermore, the total of martensite, tempered martensite and retained austenite is 2% or more and 10% or less in area fraction,
Ti of 40% or more of Tief represented by the formula (a) is present as Ti carbide, and the mass of the Ti carbide having a circle equivalent particle diameter of 7 nm or more and 20 nm or less is 50% of the mass of all Ti carbides. % High-strength hot-rolled steel sheet with excellent ductility, fatigue characteristics and corrosion resistance.
Tief = [Ti] −48 / 14 × [N] −48 / 32 × [S] (a)
However, [Ti] [N] [S] represents mass% of Ti, N, and S, respectively.
(2)
In addition,
Si: 0.500% or less,
Nb: 0.010-0.100%
V: 0.010-0.300%,
Cu: 0.01 to 1.20%,
Ni: 0.01-0.60%,
Cr: 0.01-2.00%
Mo: 0.01 to 1.00%,
A high-strength hot-rolled steel sheet excellent in ductility, fatigue properties and corrosion resistance according to (1), characterized by containing one or more of the following.
(3)
In addition,
Mg: 0.0005 to 0.0100%,
Ca: 0.0005 to 0.0100%,
REM: 0.0005 to 0.1000%,
A high-strength hot-rolled steel sheet excellent in ductility, fatigue properties, and corrosion resistance according to any one of (1) and (2), characterized by containing one or more of the following.
(4)
In addition,
B: 0.0002 to 0.0020%,
A high-strength hot-rolled steel sheet excellent in ductility, fatigue properties, and corrosion resistance according to any one of (1) to (3).
(5)
(1) It is a manufacturing method of the high strength hot-rolled steel plate excellent in ductility, fatigue characteristics, and corrosion resistance as described in any one of (4), Comprising: It is described in any one of (1)-(4) When the slab having the component composition is heated to a temperature equal to or higher than T1 defined by the formula (b), and the heated slab is hot-rolled,
The time (t1) until the center temperature of the steel plate reaches the Ar3 temperature determined by the formula (d) from 1000 ° C. is within 9.0 seconds,
Of the hot rolling, in the finish rolling consisting of continuous rolling of a plurality of stages, the finish rolling stage at the rearmost stage among the finish rolling stages having a reduction rate of 10% or more is the temperature T2 defined by the formula (c). With respect to (T2-20) ° C. or more and (T2 + 100) ° C. or less, and the hot-rolled steel sheet obtained thereafter is
Cool at an average cooling rate of 20 ° C / second or more to a temperature of 730 ° C or more and 830 ° C or less, then air-cool in a temperature range of 730 ° C or more and 830 ° C or less for 3 seconds or more, and then cool at an average cooling rate of 40 ° C / second or more. Then, a method for producing a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics and corrosion resistance, characterized by winding at a temperature of 300 ° C. or lower.
T1 (° C.) = 7000 / {2.75-log ([Ti] × [C])}-273 (b)
T2 (° C.) = 870 + 10 × ([C] + [N]) × [Mn] + 350 × [Nb] + 250 × [Ti] + 40 × [B] + 10 × [Cr] + 100 × [Mo] + 100 × [V] ... (c)
Ar3 (° C.) = 868−396 × [C] −68.1 × [Mn] + 24.6 × [Si] −36.1 × [Ni] −24.8 × [Cr] −20.7 × [Cu ] + 250 × [Al] (d)
However, when the Ar3 temperature calculated by the formula (d) exceeds 900 ° C., Ar3 = 900 ° C., and [X] in the formula represents the mass% of the component element X contained in the steel.
In addition, square brackets ([]) in the formulas (a), (b), (c), and (d) indicate mass% of the elements in the brackets in the steel sheet (hereinafter the same in this specification).

  According to the present invention, a high-strength hot-rolled steel sheet having excellent ductility, fatigue characteristics, and corrosion resistance can be provided. If this steel plate is used, it will be possible to extend the press formability and fatigue life of materials applied to undercarriage parts for automobile materials, and it will be possible to promote weight reduction of automobile bodies, contributing to industrial contributions. It is remarkable.

It is a figure explaining a repeated stress-strain curve and repeated yield stress. It is a figure which shows the example of the low cycle fatigue test piece without a notch. It is a figure which shows the example of the fatigue test piece with a notch.

The contents of the present invention will be described in detail below.
[Chemical composition of steel sheet]
First, the reasons for limiting the chemical components of the hot-rolled steel sheet of the present invention will be described. In addition,% of content is the mass%.

(C: 0.030% to 0.200%)
C is one of the important elements in the present invention. In addition to generating martensite and stabilizing austenite, C forms a Ti carbide and contributes greatly to improving the strength of the hot-rolled steel sheet by strengthening the structure and precipitation. If it is less than 0.030%, a strength of 540 МPa cannot be secured. Further, since the notch fatigue characteristics tend to improve when the hard phase fraction increases, addition of 0.050% or more is desirable, and 0.060% is more desirable.
On the other hand, if added over 0.200%, the area ratio of the low-temperature transformation product, which is the hard second phase, increases, and the hole expandability decreases. Since the hole expansion property tends to improve as the addition amount decreases, it is preferably 0.180% or less, more preferably 0.150% or less.
Therefore, the C content is 0.030% to 0.200%.

(Si: 0-0.500%)
Si is a deoxidizing element and is also an element that is involved in the formation of ferrite and increases the temperature of the ferrite region to the higher temperature side as the content increases, thereby expanding the two-phase region of ferrite and austenite. Originally, it is desirable to contain Si in order to obtain the composite structure steel of the present invention. However, in the present invention, addition of Si is not essential because the ferrite region temperature is increased to the high temperature side by adding 0.100% or more of Al.
Further, Si significantly generates a Ti scale stripe-like Si scale pattern on the surface of the steel sheet, and remarkably deteriorates the surface properties. And the productivity of the scale removal process (pickling etc.) in a finishing line may be reduced extremely. If Si is contained in excess of 0.500%, the surface properties are remarkably deteriorated, and the productivity of the pickling process is extremely deteriorated. Moreover, whatever scale removal method is implemented, chemical conversion property deteriorates and post-coating corrosion resistance decreases. Therefore, the Si content is 0.500% or less.
On the other hand, when the Si content exceeds 0.070%, the Si scale pattern starts to be scattered on the steel plate surface. Therefore, the Si content is desirably 0.070% or less, and more desirably 0.050% or less.

(Mn: 0.10 to 3.00% or less)
In addition to solid solution strengthening, Mn is added to improve hardenability and to generate martensite or austenite in the steel sheet structure. When the Mn content is added to exceed 3.00%, a segregation band of Мn is generated at the central portion in the plate thickness direction of the steel sheet, and this segregation band becomes a starting point of cracking, so that the hole expansion rate is lowered. On the other hand, if the Mn content is less than 0.10%, it is difficult to exhibit the effect of suppressing pearlite that causes a decrease in the hole expansion rate during cooling. Therefore, the content of Mn is set to 0.10% to 3.00%.
Further, from the viewpoint of ensuring sufficient hardenability and ensuring the effect of suppressing pearlite, the Mn content is desirably 0.30% or more, and more desirably 0.50% or more. On the other hand, the content of Mn is desirably 2.50% or less, and more desirably 2.00% or less, from the viewpoint of suppressing a decrease in ductility due to an increase in center segregation.

(P: 0.100% or less)
P is an impurity contained in the hot metal and is an element that segregates at the grain boundary and lowers the low temperature toughness as the content increases. For this reason, the P content is preferably as low as possible, and if it exceeds 0.100%, the workability and weldability are adversely affected, so it is made 0.100% or less. In particular, considering the weldability, the P content is preferably 0.030% or less.

(S: 0.0300% or less)
S is an impurity contained in the hot metal. If the content is too large, S is an element that not only causes cracking during hot rolling, but also generates inclusions such as MnS that degrade the hole expansion property. For this reason, the S content should be reduced as much as possible. However, if it is 0.0300% or less, it is an acceptable range, so it is 0.0300% or less. However, the S content when a certain degree of hole expansibility is required is preferably 0.0100% or less, more preferably 0.0050% or less.

(Al: 0.100 to 2.000%)
Al is an element effective as a deoxidizer for molten steel. Al is a strong ferrite-forming element and has the effect of increasing the Ar3 temperature. Therefore, the crystal grains having an average orientation difference in the crystal grains of 0 to 0.5 ° are set to 50% or more in area ratio. It is an essential element. In order to obtain these effects, the lower limit of the Al content is 0.100%. A preferable lower limit of the Al content is 0.130%, and a more preferable lower limit of the Al content is 0.150%.
On the other hand, if the Al content exceeds 2.000%, cracks may occur during rolling. Therefore, the upper limit of the Al content is set to 2.000%. Further, when the Al content exceeds 1.000%, weldability, toughness, and the like begin to deteriorate. Therefore, the upper limit of the preferable Al content is 1.000%, and the more preferable upper limit of the Al content is 0.00. 500%.

(N: 0.0100% or less (excluding 0))
N may be added because it exists as TiN and contributes to improvement of low-temperature toughness through refinement of the crystal grain size during slab heating. However, nitrides in steel need to be 0.0100% or less in order to reduce the hole expansion rate. Desirably, it is 0.0050% or less.
On the other hand, 0.0005% or less is not economically desirable, so 0.0005% or more is desirable.

(O: 0.0100% or less)
O forms an oxide and degrades the moldability, so the amount added must be suppressed. In particular, when O exceeds 0.0100%, this tendency becomes remarkable, so it is necessary to make it 0.0100% or less.
On the other hand, since it is economically unpreferable to set it as less than 0.0010%, it is desirable to set it as 0.0010% or more.

(Ti: 0.010 to 0.380%)
Ti is added to achieve both excellent fatigue strength and high strength by precipitation strengthening. If Ti is less than 0.010%, the effect of precipitation strengthening cannot be obtained, so it is necessary to add 0.010% or more. If added over 0.380%, the above effect is saturated and the economic efficiency is lowered. Therefore, the Ti content is 0.010% to 0.380%.

(Tief: 0.010-0.300%)
Tief = [Ti] −48 / 14 × [N] −48 / 32 × [S] (a)
Ti nitride and Ti sulfide are generated at a higher temperature than Ti carbide. For this reason, when there are many N and S in steel, Ti carbide cannot fully be generated. Therefore, an index called Tief represented by the equation (a) was used as an index related to the generation of Ti carbide. If the Tief is less than 0.010%, the amount of Ti carbide precipitated is small, so that it is impossible to achieve both fatigue strength and high strength due to Ti carbide. Desirably, Tief is 0.025% or more. Moreover, even if Tief is added in excess of 0.300%, the above effect is saturated and the economy is lowered. Further, if Tief exceeds 0.150%, the tundish nozzle may be easily clogged at the time of casting, so 0.150% or less is desirable.

  The above are the basic chemical components of the hot-rolled steel sheet of the present invention, but can further contain the following components.

(Nb: 0 to 0.100%)
Nb may be added because this carbonitride or solute Nb delays grain growth during hot rolling, so that the grain size of the hot-rolled sheet can be reduced and the low-temperature toughness is improved. Even if the Nb content exceeds 0.100%, the above effect is saturated and the economy is lowered. Further, if the Nb content is less than 0.010%, the above effect cannot be obtained sufficiently. Therefore, when Nb is contained as necessary, the Nb content is preferably 0.010% to 0.100%.

(V: 0 to 0.300%)
V is an element having an effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the V content exceeds 0.300%, the above effect is saturated and the economy is lowered. Further, when the V content is less than 0.010%, the above effects cannot be sufficiently obtained. Therefore, when V is contained as necessary, the V content is preferably 0.010% to 0.300%.

(Cu: 0 to 2.00%)
Cu is an element having an effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Cu content exceeds 2.00%, the above effect is saturated and the economy is lowered. Further, if the Cu content is less than 0.01%, the above effect cannot be sufficiently obtained. When the Cu content exceeds 1.20%, scratches due to scale may occur on the surface of the steel sheet. Therefore, when Cu is contained as necessary, the Cu content is preferably 0.01% to 1.20%.

(Ni: 0.01% to 2.00%)
Ni is an element that has an effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Ni content exceeds 2.00%, the above effect is saturated and the economy is lowered. Further, if the Ni content is less than 0.01%, the above effects cannot be obtained sufficiently. When the Ni content exceeds 0.60%, the ductility starts to deteriorate. Therefore, when Ni is contained as necessary, the Ni content is preferably 0.01% to 0.60%.

(Cr: 0 to 2.00%)
Cr is an element having an effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Cr content exceeds 2.00%, the above effect is saturated and the economy is lowered. Further, when the Cr content is less than 0.01%, the above effect cannot be obtained sufficiently. Therefore, when Cr is contained as necessary, the Cr content is desirably 0.01% to 2.00%.

(Mo: 0 to 1.00%)
Mo is an element having an effect of improving the strength of the hot-rolled steel sheet by precipitation strengthening or solid solution strengthening, and may be added. Even if the Mo content exceeds 1.00%, the above effect is saturated and the economy is lowered. Further, if the Mo content is less than 0.01%, the above effect cannot be obtained sufficiently. Therefore, when Mo is contained as necessary, the Mo content is desirably 0.01% to 1.00%.

(Mg: 0 to 0.0100%)
Mg is an element that controls the form of non-metallic inclusions that become a starting point of fracture and cause deterioration of workability, and thus improves workability, so Mg may be added. Even if the Mg content exceeds 0.0100%, the above effect is saturated and the economy is lowered. In addition, the effect of Mg becomes remarkable when it is added in an amount of 0.0005% or more. Therefore, when Mg is contained as necessary, the Mg content is preferably 0.0005% to 0.0100%.

(Ca: 0 to 0.0100%)
Ca may be added because it is a starting point of destruction and is an element that controls the form of non-metallic inclusions that cause deterioration of workability and improves workability. Even if the Ca content exceeds 0.0100%, the above effect is saturated and the economy is lowered. In addition, when the Ca content is 0.0005% or more, the effect becomes remarkable. Therefore, when Ca is contained as necessary, the Ca content is preferably 0.0005% to 0.0100%.

(REM: 0 to 0.1000%)
REM (rare earth element) is an element that controls the form of non-metallic inclusions, which is a starting point of destruction and causes deterioration of workability, and may be added. Even if the content of REM exceeds 0.1000%, the above effect is saturated and the economic efficiency is lowered. Moreover, the effect becomes remarkable when the content of REM is 0.0005% or more. Therefore, when REM is contained as necessary, the REM content is preferably 0.0005% to 0.1000%.

(B: 0 to 0.0100%)
B segregates at the grain boundaries and improves the low temperature toughness by increasing the grain boundary strength. Therefore, it may be added. When the addition amount of B is more than 0.0100%, the effect is saturated and the economy is inferior. Moreover, this effect becomes remarkable when the amount of B added to the steel sheet is 0.0002% or more. B is a strong quenching element, and when adding more than 0.0020%, the area ratio of crystal grains in which the average of the orientation difference in the crystal grains is 0 to 0.5 ° is reduced. There is a fear. Therefore, when B is contained as necessary, the B content is preferably 0.0002% to 0.0020%.

  In addition, about other elements, even if it contains Sn, Zr, Co, Zn, and W 1% or less in total, it has confirmed that the effect of this invention is not impaired. Among these elements, Sn is preferably 0.05% or less because wrinkles may occur during hot rolling.

[Microstructure of steel sheet]
The microstructure of the steel sheet will be described.
The microstructure of the steel sheet is subjected to EBSD analysis at a measurement interval of 1 μm or less using an EBSD method (electron beam backscatter diffraction pattern analysis method) often used for crystal orientation analysis. With respect to the orientation of measurement points with a measurement interval of 1 μm or less obtained by EBSD analysis, the case where the orientation difference between adjacent measurement points is 15 ° or more is defined as a grain boundary, and the region surrounded by this grain boundary is defined as a crystal grain To do.
And about the orientation of all the measurement points inside this crystal grain, the orientation difference of adjacent measurement points is calculated | required, and let the average value of these orientation differences be the average of the orientation difference in a crystal grain.
The steel sheet of the present invention is characterized in that the crystal grains having an average orientation difference within the crystal grains of 0 to 0.5 ° are contained in an area ratio of 50% or more with respect to all the crystal grains. Such crystal grains have high ductility and are further strengthened by precipitation with Ti carbide. Therefore, by securing such a crystal grain at a certain ratio or more, ductility can be improved while maintaining the tensile strength (TS) at 540 MPa or more.

The proportion of crystal grains having an average misorientation within the crystal grains of 0 to 0.5 ° can be measured, for example, by the following method.
A sample is taken so that a cross section (width direction cross section) of the width direction of the steel sheet viewed from the rolling direction becomes an observation surface at either 1/4 W (width) or 3/4 W (width) of the plate width, EBSD analysis is performed on a rectangular region having a width direction of 200 μm × thickness direction of 100 μm at a measurement depth of 0.2 μm at a position that is ¼ depth from the steel plate surface.
Here, the EBSD analysis is performed using a device configured with a thermal field emission scanning electron microscope (for example, JSMOL JSM-7001F) and an EBSD detector (for example, TSL HIKARI detector) at 200 to 300 points / second. Perform at analysis speed. With respect to the obtained crystal orientation information, crystal grains having an orientation difference within a crystal grain of 15 ° or more and an equivalent circle diameter (diameter) of 0.3 μm or more are extracted, and an average orientation difference within the crystal grain is calculated. And the area ratio of the crystal grain whose average orientation difference in a crystal grain is 0-0.5 degree is calculated | required. The definition of the crystal grains and the calculation of the average orientation difference in the crystal grains can be obtained using, for example, software “OIM Analysis ^™ ” attached to the EBSD analyzer.

  When crystal grains having an average orientation difference within a crystal grain of 0 to 0.5 ° are less than 50% in area ratio, ductility deteriorates and (TS) × (u−El) ≧ 8000 MPa% is satisfied. Disappear. Since the ductility improves as the area ratio of the crystal grains having an average orientation difference within the crystal grains of 0 to 0.5 ° increases, the area ratio is preferably 60% or more, more preferably 80% or more. Good. However, if the area ratio exceeds 98%, the hard phase fraction of martensite, tempered martensite, and retained austenite structure decreases, and the notch fatigue characteristics decrease. Therefore, the upper limit is desirably 98%.

  In the present embodiment, the crystal grains whose average orientation difference in the crystal grains is 0 to 0.5 ° and the ferrite defined from the observation result of the optical microscope are not directly related. In other words, for example, even if there is a hot-rolled steel sheet having a ferrite area ratio of 50% or more, the proportion of crystal grains having an average misorientation within the crystal grains of 0 to 0.5 ° is 50% or more. Is not limited. Therefore, the characteristics corresponding to the hot-rolled steel sheet according to this embodiment cannot be obtained only by controlling the ferrite area ratio.

[Martensite + Tempered martensite + Residual austenite: 2% to 10%]
The steel sheet of the present invention is characterized in that the total structure of martensite, tempered martensite and retained austenite is 2% or more and 10% or less in terms of area fraction. The hard phase of martensite or tempered martensite or retained austenite becomes an obstacle to fatigue crack propagation in the soft phase and has the effect of reducing the fatigue crack propagation rate, thus contributing to the improvement of notch fatigue properties. From this, the total of the hard phases of martensite, tempered martensite and retained austenite is made 2% or more in area fraction. If the hard phase fraction is less than 2%, it will not satisfy (c-FL) / (TS) ≧ 0.25, which is an indication of a high-strength steel sheet having excellent notch fatigue properties. % Or more is desirable. More desirably, it is 5% or more. However, (c-FL) indicates the fatigue limit of the notch fatigue test, and (TS) indicates the tensile strength. In addition, the method of calculating | requiring the fatigue limit (c-FL) of a notch fatigue test is mentioned later.

  On the other hand, when the area fraction of these structures exceeds 10%, the hole expansion ratio (λ), which is one of press formability, decreases, so the total structure of martensite, tempered martensite, and retained austenite is The area fraction must be 10% or less.

  The area fraction of martensite and tempered martensite constituting the steel sheet structure of the present invention is measured using an FE-SEM (field emission scanning electron microscope). The area fraction of retained austenite is measured using the EBSD method (electron beam backscatter diffraction pattern analysis method). Specifically, the area fraction of the martensite and tempered martensite constituting the steel sheet structure of the present invention is the width direction at either 1/4 W or 3/4 W as the sheet width W from one end of the steel sheet. A sample is collected so that the cross section becomes the observation surface (hereinafter, the collected material is referred to as “tissue measurement sample”), the observation surface is polished, nital etched, 1/4 of the plate thickness, 3 The range of / 8 thickness and 1/2 thickness was obtained by observing with FE-SEM (field emission scanning electron microscope). When observed by FE-SEM, martensite was a lath-like (thin and long plate-like) structure and no carbide was precipitated. Tempered martensite is a lath-like structure in which carbides are precipitated in multiple variants (cementite is precipitated in various directions).

  In addition, it was judged that the carbide | carbonized_material precipitated with the single variant (Cementite precipitates in one direction.) Is bainite. Using a FE-SEM, a field of 120 μm × 100 μm was measured at a magnification of 1000 times and 10 fields were measured for each of the 1/4 thickness, 3/8 thickness, and 1/2 thickness ranges. For each field of view, the area fractions of martensite and tempered martensite were determined, and the average value thereof was taken as the representative area fraction of martensite and tempered martensite.

  The area fraction of residual austenite was measured using the EBSD method (electron beam backscatter diffraction pattern analysis method) after removing the processed layer from the structure measurement sample by electropolishing. The crystal grain in which the back scattering characteristic of FCC metal was obtained was defined as retained austenite. Using the EBSD method, regions of 100 μm × 100 μm or more were observed in each of the thickness ranges of ¼ thickness, 3/8 thickness, and ½ thickness. For each range, the area fraction of retained austenite was obtained, and the average value thereof was taken as the representative area fraction of retained austenite.

[Ti carbide]
Next, the state and amount of Ti carbide in the structure will be described. Conventionally, it has been reported that precipitation strengthening by Ti carbide is inferior in fatigue properties to solid solution strengthening by Si or the like. However, as a result of intensive studies by the inventors, Ti% having a mass% of 40% (0.4 times) or more of Tief calculated from Ti, N, and S in steel (determined by the equation (a)). However, among the precipitated Ti carbides, those whose Ti carbide equivalent particle diameter (in this specification, simply referred to as the particle diameter) is 7 nm to 20 nm. It has been found that when the mass fraction is 50% or more, fatigue characteristics equal to or better than solid solution strengthening can be obtained.

  As described in Non-Patent Document 3, a composite steel having a high repeated yield stress (c-YP) has good low cycle fatigue characteristics and high cycle fatigue characteristics. Repeated yield stress (c-YP) is the resistance to deformation of a material after repeated deformation, and is a solid solution strengthened steel that does not reduce the strengthening function of the steel sheet even when repeatedly deformed. In the strengthening, the one using a precipitate having a large particle size is increased. A material having a large ratio of the repeated yield stress (c-YP) to the yield stress (YP) measured by a tensile test is known to have both low cycle fatigue characteristics and high cycle fatigue characteristics. The steel sheet satisfying (c-YP) / (YP) ≧ 0.90 has good fatigue characteristics despite its low yield stress (YP), and has an excellent balance between productivity and fatigue characteristics during press forming.

  When the Ti carbide particle size is small and the area fraction of Ti carbide particles having a particle size of 7 nm to 20 nm is less than 50%, (YP) increases greatly due to precipitation strengthening, but the increase in (c-YP) is It is small, and (c−YP) / (YP) <0.90. This is because, as a result of dislocation motion due to repetitive deformation, Ti carbide having a small particle size shears and the effect of suppressing dislocation motion due to Ti carbide decreases. A method for obtaining the repeated yield stress (c-YP) will be described later.

On the other hand, when Ti carbide having a particle diameter of 7 nm to 20 nm is sufficiently present, the dislocation is considered to move around the Ti carbide. Around the bypassed Ti carbide, annular dislocations called the Orowan loop remain, so the Orowan loop proliferates due to the movement of the dislocation, the dislocation density increases, the dislocation strengthens, and the yield stress increases from before the repeated deformation. As a result, (c−YP) / (YP) ≧ 0.90.
Further, when the Ti carbide particle size is too large and the weight fraction of Ti carbide particles having a particle size of 7 nm to 20 nm is less than 50% of the total Ti carbide weight, the effect of suppressing dislocation motion by the Ti carbide is small. Thus, the effect of improving these fatigue characteristics is reduced.

  Similarly, even when Ti corresponding to less than 40% (0.4 times) mass% of Tief calculated from Ti, N, and S in steel precipitates as Ti carbide, dislocation motion due to Ti carbide is suppressed. As a result, the effect of improving fatigue characteristics is reduced. Therefore, the above (c−YP) / (YP) ≧ 0.90 cannot be satisfied. Ti corresponding to 40% (0.4 times) or more of Tief calculated from Ti, N, and S in steel needs to be precipitated as Ti carbide. From the viewpoint of ensuring the effect of improving fatigue characteristics, Ti deposited as Ti carbide is desirably 45% (0.45 times) or more of Tief.

  The amount of Ti carbide deposited can be measured by electrolyzing the steel sheet and identifying the weight of Ti in the residual residue by chemical analysis or the like. Specifically, the total weight of precipitated Ti is obtained from the weight of Ti in the residue. Further, the weight of Ti precipitated as Ti nitride (TiN) can be obtained from the weight of nitrogen contained in the steel. By subtracting the weight of Ti deposited as TiN from the total weight of Ti deposited, the weight of Ti in the Ti carbide can be determined. Thereby, the mass% of Ti contained in Ti carbide can be obtained.

  The means for identifying the Ti carbide and measuring the particle size is not particularly specified. For example, by using 3D-AP (three-dimensional atom probe), the Ti carbide is measured by measuring the positions of Ti and C in the steel sheet. The particle diameter can be measured with high accuracy even for Ti carbide having a small particle diameter. Specifically, at least 20 fields of view having a width of 10 μm × 10 μm of the sample are observed, and the magnification of the Ti carbide particles in the range of 1 nm to 100 nm is increased while the magnification is increased according to the particle size of the Ti carbides. What is necessary is just to obtain | require distribution and to obtain | require the ratio of the particle size of 7 nm-20 nm which occupies in it by a weight ratio. In addition, if it is Ti carbide with a particle size of 5 nm or more, an electron diffraction pattern is acquired using Fe-TEM, and Ti carbide is identified and measured by measuring the orientation relation and lattice spacing with Fe as a parent phase. Is possible.

  The repeated yield stress (c-YP) can be measured by executing a low cycle fatigue test under different strain amplitude conditions. In this study, low cycle fatigue tests were performed at strain amplitude levels of 0.2%, 0.3%, 0.5%, 0.8%, and 1.0%. The maximum stress of each strain amplitude at the time when half the number of fatigue tests was performed was connected to create a repeated stress-strain curve. As shown in FIG. 1, a straight line having a Young's modulus slope is inserted into the repeated stress-strain curve, and the intersection with the repeated stress-strain curve is repeated as the yield stress (c-YP). Defined.

  The high-strength hot-rolled steel sheet of the present invention having the structure and composition as described above is provided with a hot-dip galvanized layer by hot dip galvanizing treatment on the surface, and further, an alloyed galvanized layer by alloying after plating. Corrosion resistance can be improved. The plating layer is not limited to pure zinc, and elements such as Si, Mg, Al, Fe, Mn, Ca, and Zr may be added to further improve the corrosion resistance. By providing such a plating layer, the excellent punching fatigue characteristics and workability of the present invention are not impaired. Moreover, the effect of the present invention can be obtained regardless of the surface treatment layer formed by organic film formation, film lamination, organic salt / inorganic salt treatment, non-chromic treatment, or the like.

[Steel plate manufacturing method]
The production method preceding hot rolling is not particularly limited. That is, various secondary smelting may be performed subsequent to smelting in a blast furnace, electric furnace, or the like to adjust to the above-described component composition, and then cast by a method such as normal continuous casting or thin slab casting. At that time, as long as it can be controlled within the component range of the present invention, scrap may be used as a raw material.

  The casting slab is heated to a predetermined temperature when starting hot rolling. In the case of continuous casting, it may be cooled once to a low temperature and then heated again and then hot rolled, or it may be heated and subsequently hot rolled following continuous casting without being cooled.

The slab heating time of hot rolling needs to be T1 ° C. or higher represented by the formula (b). When normal slab casting is performed, the slab temperature is lowered to an Ar3 temperature or lower, so that Ti carbide precipitates in the structure. In order to control the particle size of the Ti carbide, it is necessary to first solutionize the Ti carbide precipitated in the slab. When the slab heating temperature is less than T1 ° C., Ti carbide precipitated in the slab does not sufficiently dissolve, and the particle size of the Ti carbide cannot be controlled. Therefore, the temperature of the heating furnace is set to T1 ° C. or higher.
Moreover, the upper limit of the slab heating temperature is not particularly defined. However, an excessively high heating temperature is not economically preferable because the surface of the slab is oxidized and becomes scale. Therefore, the upper limit of the slab heating temperature is desirably 1300 ° C.
T1 (° C.) = 7000 / {2.75-log ([Ti] × [C])}-273 (b)

After the slab heating, a hot rolled steel sheet is obtained by a hot rolling rough rolling process and a subsequent finish rolling process on the slab extracted from the heating furnace. During this hot rolling, the time (t1) from when the center temperature of the steel sheet drops to 1000 ° C. to the Ar3 temperature (transformation point temperature) represented by the formula (d) needs to be within 9.0 seconds. There is.
Ar3 (° C.) = 868−396 × [C] −68.1 × [Mn] + 24.6 × [Si] −36.1 × [Ni] −24.8 × [Cr] −20.7 × [Cu ] + 250 × [Al] (d)
However, when the Ar3 temperature calculated by the formula (d) exceeds 900 ° C., Ar3 = 900 ° C., and the time from when the steel sheet temperature falls to 1000 ° C. until it reaches 900 ° C. is t1. The center temperature of the steel sheet described here can be obtained by thermal analysis from the temperature of the steel sheet surface and the cooling history.

  In order to obtain the effect of the present invention, it is necessary that 50% or more of the precipitated Ti carbides have a mass equivalent particle diameter of 7 nm to 20 nm. Since Ti carbide precipitated in the austenite region is 20 nm or more and does not contribute to improvement of fatigue characteristics, it is necessary to suppress precipitation of Ti carbide in the austenite region. Precipitation of Ti carbide in the austenite region is significant when the center temperature of the steel sheet is 1000 ° C. or less. Therefore, it is effective to shorten the time during which the center temperature of the steel sheet is maintained at Ar 3 temperature or more and 1000 ° C. or less. According to the study by the present inventors, when this time is longer than 9.0 seconds, coarse Ti carbide is precipitated, and the area fraction of Ti carbide having a particle diameter of 7 nm to 20 nm is reduced, so that a desired area fraction is obtained. 50% cannot be obtained. There is a possibility that the center temperature of the steel sheet will change from the Ar3 temperature to 1000 ° C in the latter part of the rough rolling, the finish rolling, and the front part of the cooling device, and the temperature history of this process is controlled and maintained in this temperature range. It must be within 9.0 seconds.

Finish rolling is usually performed by multi-stage (for example, 6-stage or 7-stage) continuous rolling. And the finish rolling performed by this multi-stage continuous rolling is rolled with a reduction ratio higher on the upstream side (upstream side) than on the downstream side (downstream side), and on the downstream side (downstream side) with a lower reduction ratio. Sometimes. In the present invention, in the finish rolling performed by this multi-stage continuous rolling, the finish rolling stage on the most downstream side (downstream side) among the finish rolling stages having a reduction ratio of 10% or more is expressed by the formula (c). It is good to carry out at the rolling temperature of (T2-20) degreeC or more and (T2 + 100) degreeC or less using temperature T2 prescribed | regulated. Desirably, it is good to set it as (T2-20) degreeC or more and (T2 + 30) degreeC or less.
T2 (° C.) = 870 + 10 × ([C] + [N]) × [Mn] + 350 × [Nb] + 250 × [Ti] + 40 × [B] + 10 × [Cr] + 100 × [Mo] + 100 × [V] ... (c)
That is, it does not necessarily manage the rolling temperature of the finishing rolling stage on the most downstream side (downstream side), but the finishing rolling on the most downstream side (downstream side) among finishing rolling stages having a reduction ratio of 10% or more. The rolling temperature of the stage is managed.
In addition, a rolling reduction is calculated | required with the following formula | equation for every stage.
Reduction ratio = (sheet thickness on the entry side of the finishing mill−sheet thickness on the exit side of the finishing mill) / (sheet thickness on the entry side of the finishing mill) × 100%

  When the rolling temperature of the finishing rolling stage at the rearmost stage among the finishing rolling stages with a reduction ratio of 10% or more is less than (T2-20) ° C., rolling is performed with austenite recrystallization suppressed. As a result, the aspect ratio of austenite increases. Since the shape is inherited by hard phases such as martensite, tempered martensite, and retained austenite, the aspect ratio of the hard phase particularly in the central portion of the plate thickness increases, and the hole expansion rate decreases.

When the rolling temperature of the finish rolling stage at the rearmost stage among the finish rolling stages with a rolling reduction of 10% or more is (T2 + 100) ° C., the average orientation difference in the crystal grains is 0 to 0.5. The area ratio of the crystal grains that are ° becomes less than 50%, and the ductility decreases. This is presumably because austenite coarsened after recrystallization and the transformation temperature decreased. Further, among the finish rolling stages having a rolling reduction of 10% or more, the most subsequent finish rolling stage is (T2-20) ° C. or more (T2 + 100) with respect to the temperature T2 defined by the equation (c). The reason why the rolling is performed at a temperature of ℃ or less is that when the rolling reduction is less than 10%, there is no crystal grain refining effect, and the rolling reduction needs to be 10% or more.
Rolling with a rolling reduction of 40% or more places a heavy burden on the rolling mill, so the rolling reduction is preferably less than 40%.

The steel plate obtained is cooled to a temperature of 730 ° C. or higher and 830 ° C. or lower (primary cooling) after the finish rolling step of the rearmost stage among the finish rolling stages having a rolling reduction of 10% or more. In this primary cooling, it is necessary to cool at an average cooling rate of 20 ° C./s or more. When the average cooling rate of primary cooling is less than 20 ° C./s, the area ratio of crystal grains having an average orientation difference within the crystal grains of 0 to 0.5 ° is less than 50%, and ductility is lowered.
This is presumably because austenite coarsens due to grain growth after rolling, so that the austenite grain interfacial area, which is the nucleation site of ferrite transformation, decreases and the transformation temperature decreases.

  If the average cooling rate of the primary cooling is large, the transformation temperature rises and the ductility is improved, so the upper limit is not specified. However, since it is difficult to control the cooling stop temperature when the average cooling rate exceeds 200 ° C./s, it is desirable that the average cooling rate be 200 ° C./s or less.

In the intermediate air cooling following the primary cooling, air cooling is performed for 3 seconds or more in a temperature range of 730 ° C. or higher and 830 ° C. or lower. This is indispensable for setting the area ratio of crystal grains having an average orientation difference in the crystal grains of 0 to 0.5 ° to 50% or more, preferably 5 seconds in a temperature range of 750 ° C. or more and 830 ° C. or less. As described above, it is more preferable to set the temperature in the temperature range from 780 ° C. to 830 ° C. for 5 seconds or more.
By increasing the intermediate air cooling temperature, the area ratio of crystal grains having an average misorientation within a crystal grain of 0 to 0.5 ° is increased. There is a possibility that the crystal grains of 5 ° are those transformed at a higher temperature among the ferrites defined by observation with a conventional optical microscope. The upper limit of the intermediate air cooling time is not specified, but if it is 15 seconds or more, the effect of increasing the area ratio of the crystal grains whose average orientation difference in the crystal grains is 0 to 0.5 ° is saturated and the productivity is lowered. Therefore, the intermediate air cooling time is preferably less than 15 seconds.

  When the intermediate air cooling temperature exceeds 830 ° C., Ti having a mass corresponding to less than 40% of Tief represented by the formula (a) is precipitated as Ti carbide in the structure, and the fatigue characteristics are deteriorated. On the other hand, when the intermediate air cooling temperature is less than 730 ° C., the area ratio of the crystal grains having an average orientation difference in the crystal grains of 0 to 0.5 ° is less than 50%, and the ductility is lowered.

In the secondary cooling subsequent to the intermediate air cooling, it is necessary to cool at an average cooling rate of 40 ° C./s or higher and to take up the steel plate at a temperature of 300 ° C. or lower. When the average cooling rate of the secondary cooling is less than 40 ° C./s, the total structure of martensite, tempered martensite, and retained austenite has an area fraction of less than 2%, and the notch fatigue characteristics deteriorate.
Hard martensite, tempered martensite, and retained austenite are obtained as the average cooling rate of the secondary cooling increases, and the notch fatigue characteristics are superior. For this reason, although the upper limit of the average cooling rate is not specified, a large-scale capital investment is required to perform cooling at 500 ° C./second or more, and therefore, it is preferably less than 500 ° C./second.

  In the winding process, the winding temperature needs to be 300 ° C. or lower. When the winding temperature is higher than 300 ° C., the fraction of bainite in the structure is increased, it is difficult to secure the fraction of the hard phase, and notch fatigue characteristics are deteriorated. When emphasizing notch fatigue characteristics, the desired winding temperature is 150 ° C. or lower. This is because the notch fatigue characteristics improve as the hardness of the hard phase increases.

  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.

  In addition to the hot rolling step described above, excellent ductility and fatigue characteristics, which are the effects of the present invention, can be ensured even if production is carried out by removing some of the accompanying continuous casting, pickling and the like. Further, once the hot-rolled steel sheet is manufactured, even if heat treatment is performed online or offline in the temperature range of 100 to 600 ° C. for the purpose of improving ductility, the excellent rolling direction that is the effect of the present invention is achieved. Fatigue properties and workability can be ensured.

Test results for confirming the effects of the present invention will be described.
Table 1 shows the components of the steel subjected to the test.
Table 2-1 and Table 2-2 (in this specification, these are collectively referred to as Table 2) show the steel types of the test pieces subjected to the test and the production conditions thereof.
Table 3-1 and Table 3-2 (in this specification, these are collectively referred to as Table 3) show the evaluation results of each test piece.

Among the mechanical properties, the tensile strength characteristics (yield stress, tensile strength, uniform elongation) are determined in the width direction at either 1/4 W or 3/4 W from the plate edge, where W is the plate width. It evaluated based on JISZ2241: 2011 using the JISZ2241: 2011 No. 5 test piece extract | collected as the longitudinal direction.
The yield stress may be 0.2% proof stress, and the uniform elongation is the plastic elongation (%) at the maximum test force.
The hole expansion rate was evaluated in accordance with the test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996 by collecting test pieces from the same position as the tensile test piece collection position in the hole expansion test.
In addition, it was confirmed that the steel plate in the present invention was (TS) ≧ 540 MPa, (TS) × (u−El) ≧ 8000 MPa%, and (TS) × (λ) ≧ 36000 MPa%. However, (TS) is the tensile strength, (u-El) is the uniform elongation, and (λ) is the hole expansion rate.

  Test specimens for evaluating fatigue properties were subjected to a fatigue test by collecting fatigue test specimens having the shapes shown in FIGS. 2 and 3 so that the width direction is the long side from the same position as the tensile test specimen collection position. . The fatigue test piece shown in FIG. 2 is a test piece for obtaining repeated yield stress, which is an index of the fatigue characteristics of a material without a notch, and the fatigue test piece shown in FIG. 3 is used for obtaining the fatigue strength of a notch material. It is a notch test piece produced in (1). The fatigue test piece was ground to a depth of about 0.05 mm from the outermost layer.

  Low cycle fatigue using the test piece shown in Fig. 2 with strain rate control of 0.4% / s, strain amplitude of 0.2%, 0.3%, 0.5% and 1.0%. The test was run. The maximum stress of each strain amplitude at the time when the number of fatigue tests that were half the number of fatigue tests at the time of fatigue rupture was connected to create a repeated stress-strain curve. As shown in FIG. 1, a straight line having a Young's modulus slope is inserted into the repeated stress-strain curve, and the intersection with the stress-strain curve is defined as repeated yield stress (c-YP). did.

  Using the test piece shown in FIG. 3, a stress control axis fatigue test was conducted at a stress ratio R = 0.1 and a frequency of 5 Hz to evaluate notch fatigue characteristics. The stress that did not break after 10 million cycles was defined as the notch fatigue limit (c-FL).

  The surface characteristics were evaluated by “surface defects” and “roughness” before pickling. If this score is below the standard, the surface quality may be inferior and evaluated by the customer even after pickling due to the scale defects and surface irregularities. Here, “surface defect” indicates the result of visually confirming the presence or absence of a scale defect such as Si scale, scale, spindle, etc., “x” indicates that there is a scale defect, and “○” indicates that there is no scale defect. It showed. In addition, those having these defects partially or in area ratio of 5% or less are indicated by “Δ” as “minor”. “Roughness” is evaluated by Rz and indicates a value obtained by the measurement method described in JIS B 0601: 2001. If Rz is 20 μm or less, the surface quality is at a level with no problem.

Corrosion resistance was evaluated by "chemical conversion treatment" and "corrosion resistance after painting". First, after the pickled steel sheet was pickled, a phosphorylation treatment was applied to attach a 2.5 g / m ² zinc phosphate coating. At this stage, as “chemical conversion treatment”, the presence or absence of skeins (parts to which the chemical conversion film does not adhere) and the P ratio (X-ray diffraction intensity P of the phosphophyllite (100) surface measured using an X-ray diffractometer), The ratio of the X-ray diffraction intensity H of the face (020) plane: P ratio = P / (P + H)) was measured.

The phosphorylation treatment is a treatment using a chemical solution containing phosphoric acid and Zn ions as main components, and phosphophyllite: FeZn ₂ (PO ₄ ) ₂ .4H ₂ O between the Fe ions eluted from the steel sheet. This is a chemical reaction that produces crystals. The technical point of phosphorylation treatment is
(1) leaching Fe ions to promote the reaction;
(2) The formation of phosphophyllite crystals densely on the steel sheet surface.
Especially for (1), if oxides resulting from the formation of Si scale remain on the steel sheet surface, the elution of Fe is hindered, and no skein appears or Fe does not elute,
Hopeite: An abnormal chemical conversion coating that is not originally formed may be formed on the iron surface called Zn ₃ (PO ₄ ) ₂ .4H ₂ O, which may deteriorate the performance after coating. Therefore, it becomes important to normalize the surface so that Fe on the steel sheet surface is eluted by phosphoric acid and sufficient Fe ions are supplied.

  About this scale, it can be confirmed by observation with a scanning electron microscope. About 20 fields of view are observed at a magnification of 1000 times, and the case where the entire surface is uniformly adhered and the scale cannot be confirmed is marked as “o”. Also, if the field of view where the squeezing could be confirmed was 5% or less, it was considered as “△” as minor. More than 5% was evaluated as “x” because there was a scale.

  On the other hand, the P ratio represents the ratio of phosphite and phosphophyllite in the film obtained by chemical conversion treatment. The higher the P ratio, the more phosphophyllite is contained, and the phosphophyllite crystal It means that it is densely formed on the surface. In general, the P ratio ≧ 0.80 is required in order to satisfy the corrosion resistance performance and the coating performance. In a severe corrosive environment such as a snowmelt salt application area, the P ratio ≧ 0.85. It is required to be. Therefore, if this P ratio <0.80, the chemical conversion processability is inferior.

  Next, regarding “corrosion resistance after painting”, after the chemical conversion treatment, a 25 μm-thick electrodeposition coating was applied, and after baking at 170 ° C. for 20 minutes, the electrodeposition coating film was applied to the iron bar with a sharp knife. A 130 mm long cut was made until it reached, and 5% salt water spraying at a temperature of 35 ° C. was continued for 700 hours under the salt water spray conditions shown in JIS Z 2371. (Nichiban 405A-24 JIS Z 1522) was affixed 130 mm in length parallel to the cut portion, and the maximum coating film peeling width was measured when this was peeled off. When the maximum coating film peeling width is more than 4.0 mm, the post-coating corrosion resistance is inferior.

  As shown in Table 3, it was confirmed that the hot-rolled steel sheet according to the present invention has excellent ductility and fatigue characteristics.

  On the other hand, in Steel Nos. 2 and 20, since the heating temperature was T1 ° C. or less defined by the formula (b), the Ti carbide was not formed into a mass, and the mass of Ti carbide having a particle diameter of 7 to 20 nm with respect to the mass of all Ti carbide. Was less than 50%, and (c−YP) / (YP) ≧ 0.90 could not be satisfied.

  Steel Nos. 3 and 21 have a rolling temperature of 10% or more of the finish rolling stage where the rolling temperature of the finishing rolling stage at the rearmost stage is (T2-20) with respect to the temperature T2 defined by the formula (c). ) Since it was lower than ° C., (TS) × (λ) ≧ 36000 MPa% could not be satisfied. This is because the aspect ratio of austenite is increased by rolling in a state where austenite recrystallization is suppressed, and the shape is inherited by hard phases such as martensite, tempered martensite and residual austenite. This is probably because the aspect ratio of the hard phase at the center of the thickness has increased.

  Steel Nos. 6 and 24 are (T2 + 100) ° C. with respect to the temperature T2 where the rolling temperature of the finishing rolling stage at the rearmost stage among the finishing rolling stages with a rolling reduction of 10% or more is defined by the formula (c). As described above, the area ratio of crystal grains having an average orientation difference within the crystal grains of 0 to 0.5 ° is less than 50%, and (TS) × (u−El) ≧ 8000 MPa% can be satisfied. There wasn't.

  In Steel Nos. 7 and 25, the time until the temperature reached 1000 ° C. to Ar 3 temperature was 9.0 seconds or more, so the mass of Ti carbide having a particle diameter of 7 to 20 nm with respect to the mass of all Ti carbides was less than 50%, −YP) / (YP) ≧ 0.90 could not be satisfied.

  Steel Nos. 8 and 26 were primary cooling, and the average cooling rate was less than 20 ° C / s, so the area ratio of crystal grains with an average misorientation within the crystal grains of 0 to 0.5 ° was less than 50%. , (TS) × (u−El) ≧ 8000 MPa% could not be satisfied.

  Since steel Nos. 10 and 28 had an intermediate air cooling temperature of less than 730 ° C, the area ratio of crystal grains having an average orientation difference within the crystal grains of 0 to 0.5 ° was less than 50%, and (TS) x ( u−El) ≧ 8000 MPa% could not be satisfied.

  Since steel Nos. 13 and 31 had an intermediate air cooling start temperature of 830 ° C. or higher, Ti having a weight of less than 40% of Tief represented by the formula (a) was precipitated as Ti carbide in the structure, (c− YP) / (YP) ≧ 0.90 could not be satisfied.

  Steel Nos. 14 and 32 had an intermediate air cooling time of less than 3 seconds, so the area ratio of crystal grains having an average misorientation within the crystal grains of 0 to 0.5 ° was less than 50%, and (TS) × ( u−El) ≧ 8000 MPa% could not be satisfied.

  Steel Nos. 16 and 34 had an average secondary cooling rate of less than 40 ° C / s, so the total structure (hard phase) of martensite, tempered martensite, and retained austenite was less than 2% in area fraction. , (C-FL) / (TS) ≧ 0.25 could not be satisfied.

  Steel Nos. 17 and 35 had a coiling temperature of over 300 ° C., so the total structure (hard phase) of martensite, tempered martensite, and retained austenite was less than 2% in area fraction, (c-FL) /(TS)≧0.25 could not be satisfied.

  In Steel No. 37, the amount of C added was less than 0.030%, so the total structure (hard phase) of martensite, tempered martensite, and retained austenite was less than 2% in area fraction, and a strength of 540 МPa was secured. could not.

  In Steel No. 40, the amount of C added was over 0.200%, so the total structure (hard phase) of martensite, tempered martensite, and retained austenite was 10% or more in terms of area fraction, and (TS) × ( λ) ≧ 36000 MPa% could not be satisfied.

  In Steel No. 41, since the amount of Si added was over 0.500%, chemical conversion property and post-coating corrosion resistance decreased, and P ratio ≧ 0.80 and maximum peel width could not satisfy 4.0 mm or less. It was.

  Steel No. 46 could not satisfy (TS) × (λ) ≧ 36000 MPa% because the amount of Mn added exceeded 3.00%.

  In Steel No. 47, since the amount of Mn added was less than 0.10%, pearlite was generated during cooling, and (TS) × (λ) ≧ 36000 MPa% could not be satisfied.

  In steel No. 48, the amount of addition of P was more than 0.100%, so the workability was lowered, and (TS) × (u−El) ≧ 8000 MPa% and (TS) × (λ) ≧ 36000 MPa% were satisfied. could not.

  In Steel No. 49, since the addition amount of S was over 0.0300%, the hole expansion rate was lowered and (TS) × (λ) ≧ 36000 MPa% could not be satisfied.

  In Steel No. 50, the amount of Al added exceeded 2.000%, so cracking occurred during rolling.

  In Steel No. 55, since the amount of Al added was less than 0.100%, the area ratio of crystal grains having an average misorientation within the crystal grains of 0 to 0.5 ° was less than 50%, and (TS) × (U-El) ≧ 8000 MPa% could not be satisfied.

  In Steel No. 56, since the amount of N added was over 0.0100%, the hole expansion rate was lowered, and (TS) × (λ) ≧ 36000 MPa% could not be satisfied.

  Since Steel No. 57 had a Tief of over 0.300%, the tundish nozzle was clogged during casting.

  In Steel No. 59, since Tief was less than 0.010%, Ti having a weight of less than 40% of Tief represented by the formula (a) was precipitated as Ti carbide in the structure, and (c-YP) / (YP) ≧ 0.90 could not be satisfied.

  The hot-rolled steel sheet according to the present invention is excellent in ductility, fatigue characteristics, and corrosion resistance, and particularly excellent in fatigue characteristics and press formability in a material having no notch. Therefore, the hot-rolled steel sheet according to the present invention can be used for mechanical products that require designability and the like. In particular, it can be applied to automobile wheels.

Claims (5)

Chemical composition is mass%,
C: 0.030 to 0.200%,
Mn: 0.10 to 3.00%,
P: 0.100% or less,
S: 0.0300% or less,
Al: 0.100 to 2.000%
N: 0.0100% or less (excluding 0),
O: 0.0100% or less,
Ti: 0.010 to 0.380%,
The balance is Fe and inevitable impurities, and Tief represented by the formula (a) is 0.01 to 0.30%,
A crystal grain having an orientation difference between adjacent grain boundaries of 15 ° or more, wherein the crystal grains have an average orientation difference of 0 to 0.5 ° in the crystal grain and include an area ratio of 50% or more. Furthermore, the total of martensite, tempered martensite and retained austenite is 2% or more and 10% or less in terms of area fraction, and 40% or more by mass of Tief represented by the formula (a) is Ti as Ti carbide. High strength with excellent ductility, fatigue properties, and corrosion resistance, characterized in that the mass of the Ti carbide having a circle equivalent particle diameter of 7 nm or more and 20 nm or less is 50% or more of the mass of all Ti carbides Hot rolled steel sheet.

Tief = [Ti] −48 / 14 × [N] −48 / 32 × [S] (a)
However, [Ti] [N] [S] represents mass% of Ti, N, and S, respectively.
In addition,
Si: 0.500% or less,
Nb: 0.010-0.100%
V: 0.010-0.300%,
Cu: 0.01 to 1.20%,
Ni: 0.01-0.60%,
Cr: 0.01-2.00%
Mo: 0.01 to 1.00%,
The high-strength hot-rolled steel sheet having excellent ductility, fatigue properties, and corrosion resistance according to claim 1, characterized by containing at least one of the following.
In addition,
Mg: 0.0005 to 0.0100%,
Ca: 0.0005 to 0.0100%,
REM: 0.0005 to 0.1000%,
The high-strength hot-rolled steel sheet having excellent ductility, fatigue properties, and corrosion resistance according to claim 1 or 2, characterized by containing at least one of the following.
In addition,
B: 0.0002 to 0.0020%,
The high-strength hot-rolled steel sheet having excellent ductility, fatigue characteristics, and corrosion resistance according to any one of claims 1 to 3.
It is a manufacturing method of the high intensity | strength hot-rolled steel plate excellent in the ductility, fatigue characteristics, and corrosion resistance of any one of Claims 1-4, Comprising: The component composition of any one of Claims 1-4 is included. In heating the slab having a temperature equal to or higher than T1 defined by the formula (b) and hot rolling the heated slab,
The time t1 until the center temperature of the steel sheet reaches the Ar3 temperature determined by the formula (d) from 1000 ° C. is within 9.0 seconds,
Of the hot rolling, in the finish rolling consisting of continuous rolling of a plurality of stages, the finish rolling stage at the rearmost stage among the finish rolling stages having a reduction rate of 10% or more is the temperature T2 defined by the formula (c). With respect to (T2-20) ° C. or more and (T2 + 100) ° C. or less, and the hot-rolled steel sheet obtained thereafter is
Cool to a temperature of 730 ° C. or more and 830 ° C. or less at an average cooling rate of 20 ° C./second or more,
Then, it is air-cooled for 3 seconds or more in a temperature range of 730 ° C. or more and 830 ° C. or less,
Then, cool at an average cooling rate of 40 ° C / second or more,
A method for producing a high-strength hot-rolled steel sheet excellent in ductility, fatigue characteristics, and corrosion resistance, which is then wound at a temperature of 300 ° C. or lower.

T1 (° C.) = 7000 / {2.75-log ([Ti] × [C])}-273 (b)
T2 (° C.) = 870 + 10 × ([C] + [N]) × [Mn] + 350 × [Nb] + 250 × [Ti] + 40 × [B] + 10 × [Cr] + 100 × [Mo] + 100 × [V] ... (c)
Ar3 (° C.) = 868−396 × [C] −68.1 × [Mn] + 24.6 × [Si] −36.1 × [Ni] −24.8 × [Cr] −20.7 × [Cu ] + 250 × [Al] (d)

However, when the Ar3 temperature calculated by the formula (d) exceeds 900 ° C., Ar 3 = 900 ° C.
[X] in the formula represents mass% of the component element X contained in the steel.

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