JP5609712B2 - High-strength hot-rolled steel sheet having good ductility, stretch flangeability, and material uniformity and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength hot-rolled steel sheet useful for the use of a skeleton member of a large vehicle automobile such as a truck frame and a manufacturing method thereof.

In recent years, there has been an urgent need to improve the fuel efficiency of automobiles in order to regulate CO ₂ emissions from the viewpoint of global environmental protection, and there has been a demand for weight reduction by reducing the thickness of members used. Since safety improvement is also required in order to ensure the safety of the occupant at the time of a collision, it is essential to increase the strength of the member used for thinning.

  Many automobile parts made of steel plates are manufactured by press molding. In general, steel sheets have problems such as formability and shape stability because ductility, stretch flangeability, and the like decrease due to increased strength and springback increases. In recent years, the amount of springback can be accurately predicted by CAE (Computer Assisted Engineering). When the material variation is large, the accuracy of prediction by CAE is lowered. Therefore, there is a demand for a high-strength steel sheet excellent in material uniformity with small strength variation in addition to formability.

  At present, development that achieves both high strength and good moldability is being actively promoted. For example, Patent Document 1 describes C: 0.06 to 0.15% by mass%, Si: 1.2% or less, Mn : 0.5 to 1.6%, P: 0.04% or less, S: 0.005% or less, Al: 0.05% or less, Ti: 0.03 to 0.20%, the balance being Fe and inevitable Having a component composition consisting of mechanical impurities, 50 to 90% in volume occupancy is the ferrite phase, and the balance is substantially the bainite phase, and the total volume occupancy of the ferrite phase and bainite phase is 95% or more In the ferrite phase, a precipitate containing Ti is precipitated, the precipitate has an average diameter of 20 nm or less, and 80% or more of the amount of Ti in the steel is precipitated. , TS excellent in stretch flange characteristics and tensile fatigue characteristics is 780 MPa or more Strength hot rolled steel sheet is disclosed.

  In addition, Patent Document 2 is C: 0.015-0.06%, Si: less than 0.5%, Mn: 0.1-2.5%, P: 0.10% or less, and S: 0 in mass%. 0.01% or less, Al: 0.005 to 0.3%, N: 0.01% or less, Ti: 0.01 to 0.30%, B: 2 to 50 ppm, the balance being Fe and inevitable 0.75 <(C% / 12) / (Ti% / 48) -N% / 14-S% / 32) <1.25 and 1.0 <(Mn% + Bppm) / 10-Si%), the total area ratio of the ferrite phase and bainitic ferrite phase is 90% or more, the area ratio of cementite is 5% or less, TS is 690 to 850 MPa, and λ is 40 %, A high-strength hot-rolled steel sheet excellent in stretch flange formability is disclosed.

  Patent Document 3 includes, in mass%, C: 0.1% or less, Mo: 0.05 to 0.6%, Ti: 0.02 to 0.10%, and atoms in a metal structure having a ferrite structure. A high-formability, high-tensile hot-rolled steel sheet with a TS of 610 to 830 MPa, which is excellent in material uniformity, in which carbides containing Ti and Mo are dispersed and precipitated in a range satisfying Ti / Mo ≧ 0.1, is disclosed. Yes.

  Furthermore, in Patent Document 4, C: 0.05 to 0.12% by mass, Si: 0.5% or less, Mn: 0.8 to 1.8%, P: 0.030% or less, S: 0.01% or less, Al: 0.005 to 0.1%, N: 0.01% or less, Ti: 0.030 to 0.080%, the balance is composed of Fe and inevitable impurities The amount of Ti present in precipitates having a polygonal ferrite phase area ratio of 70% or more and less than 20 nm in size is calculated by the formula [Ti * = [Ti] − (48/14) × [N]]. A high-strength hot-rolled steel sheet that is 50% or more of the value of Ti * and has a TS of 540 to 780 MPa and excellent strength uniformity with small strength variation is disclosed.

  However, in the high-strength hot-rolled steel sheet described in Patent Document 1, it is necessary to produce hard ferrite and bainite phase with a prescribed volume fraction, but the transformation behavior is not constant with respect to the chemical composition of the steel, There is a problem that it is difficult to control during intermediate air cooling that promotes ferrite transformation. The high-strength hot-rolled steel sheets described in Patent Documents 2 and 3 have a low elongation El, and there is a problem that a steel sheet having good stretch flangeability and material stability cannot always be obtained. The high-strength hot-rolled steel sheet described in Patent Document 4 has obtained TS590 MPa or more by solid solution strengthening with Mn, but the solid solution strengthening is inferior in cost because the strengthening ratio with respect to the amount of added elements is smaller than the precipitation strengthening with Ti. Moreover, since there is much C addition amount with respect to Ti, hard cementite production | generation is inevitable. Therefore, there is also a problem inferior in stretch flangeability.

JP 2007-9322 A JP 2007-302992 A JP 2002-322541 A JP 2009-185361 A

  In order to solve the above problems, the present invention provides a hot-rolled steel sheet having high strength, excellent ductility and stretch flangeability, and good material uniformity with small variations in strength within the coil, and a method for producing the same. The purpose is to provide.

As a result of repeated studies on the above-described high-strength hot-rolled steel sheet, the present inventors have found the following knowledge.
1) After optimizing the composition of components for the purpose of controlling the TiC precipitation efficiency and the amount of cementite produced, the tensile strength is increased by a steel structure having a ferrite phase area ratio of 95% and a ferrite grain size of 10 μm or less. (TS) is 590 to 780 MPa, and total elongation (El) is 2.
A hot-rolled steel sheet having a hole expansion ratio (λ) of 8% or more and 100% or more is obtained.
2) In order to improve the material uniformity, it is important to suppress the coarsening of TiC while keeping the ferrite fraction in the steel sheet constant. For this reason, the ferrite transformation can be completed in a short time and the manufacturing cost can be reduced by setting the content of Mn content of the austenite former to 0.4 to 0.8%. it can. TS590M
In order to achieve Pa or higher, Ti needs to be contained in an amount of 0.08 to 0.16%. However, when the content of Ti as a constituent element of the precipitate is large, there is a problem that the precipitate is likely to be coarsened. The For this problem, it is important to wind up at a low temperature after obtaining precipitates during ferrite transformation. Specifically, the coiling temperature needs to be 560 ° C. or less.

The present invention is based on such knowledge and employs the following means in order to solve the above problems.
[1] By mass%
C: 0.020 to 0.065%
Si: 0.1% or less Mn: 0.40 to less than 0.80% P: 0.030% or less S: 0.005% or less Ti: 0.08 to 0.20%
Al: 0.005 to 0.1%
N: steel containing 0.005% or less, the balance being Fe and inevitable impurities, and Ti ^* defined by the following formula (1) satisfying the following formula (2) and formula (3) The ferrite structure has a steel structure with an area ratio of 95% or more and the balance is at least one of a pearlite phase, a bainite phase, and a martensite phase, and the average ferrite particle size of the ferrite is 10 μm or less The average particle size of Ti carbide precipitated in steel is 10 nm or less, and the amount of Ti precipitated as Ti carbide is 80% or more of Ti ^* . High strength hot rolled steel sheet with excellent material uniformity.
Ti ^* = Ti− (48/14) × N (1)
Ti ^* ≧ 0.08 (2)
0.300 ≦ C / Ti ^* ≦ 0.375 (3)
Here, Ti, N, and C in the formula indicate the content (% by mass) of each element.
[2] After heating the steel slab having the steel component described in [1] in a range of 1200 to 1300 ° C, the steel slab is hot-rolled at a finishing temperature of 900 ° C or higher, and 30 ° C within 2 seconds after the hot rolling. Cooling is started at a cooling rate of not less than / s and stopped at a temperature of 650 to 750 ° C., followed by a cooling process of 5 to 20 seconds, and then cooled at a cooling rate of not less than 30 ° C./s. The steel structure is coiled at 560 ° C. or lower, and the steel structure is a ferrite phase with an area ratio of 95% or more, and the balance is one or more of a pearlite phase, a bainite phase and a martensite phase. The average ferrite particle diameter of ferrite is 10 μm or less, the average particle diameter of Ti carbide precipitated in steel is 10 nm or less, and the amount of Ti precipitated as Ti carbide is 80% or more of Ti ^*. Some ductility, elongation hula Process for producing a high strength hot rolled steel sheet having excellent di properties and material homogeneity.

  According to the present invention, as a hot-rolled steel sheet having high strength, excellent ductility and stretch flangeability, and good material uniformity with small variation in strength within the steel sheet, the tensile strength (TS) is 590 to 780 MPa or more. A high-strength hot-rolled steel sheet having a total elongation (El) of 28% or more, a hole expansion value (λ) of 100% or more, and a TS variation ΔTS of 15 MPa or less can be produced. The high-strength hot-rolled steel sheet of the present invention is suitable for structural members such as pillars and members of passenger cars and truck frames.

  Details of the present invention will be described below. The unit representing the element content is mass%, and is simply indicated by “%” hereinafter.

1) Steel component The reason which limited the steel component (chemical component) in this invention is demonstrated.
C: 0.020-0.065%
C is an element that contributes to high strength by forming fine Ti carbides in the ferrite phase. In order to obtain a hot rolled steel sheet having a TS of 590 MPa or more, the C content needs to be 0.020% or more. On the other hand, if the amount of C exceeds 0.065%, not only El and λ are lowered, but also the rate of progress of ferrite transformation is slowed, causing a reduction in material uniformity. Therefore, the C content is 0.020 to 0.065%. Preferably, the C content is 0.020% or more and 0.055% or less, more preferably 0.050% or less.

-Si: 0.1% or less If the Si amount exceeds 0.1%, the Ar ₃ point rises too much, making it difficult to obtain a fine and sized structure of the ferrite phase. Furthermore, since an increase in the amount of Si leads to deterioration of toughness and fatigue resistance, the amount of Si is set to 0.1% or less, preferably 0.05% or less.

Mn: 0.40 to less than 0.80% Mn is effective for increasing strength and refining ferrite grains. In order to obtain a hot-rolled steel sheet having a TS of 590 MPa or more and a ferrite particle size of 10 μm or less, the Mn amount needs to be 0.40% or more. On the other hand, if the amount of Mn is 0.80% or more, the progress of ferrite transformation becomes slow and the uniformity of the material is lowered. Therefore, the amount of Mn is made 0.40 to less than 0.80%.

-P: 0.030% or less When the amount of P exceeds 0.03%, segregation to the grain boundary becomes remarkable, leading to deterioration of toughness and weldability. Therefore, the P content is 0.03% or less, but it is desirable to reduce it as much as possible.
-S: 0.005% or less S forms sulfide with Mn and Ti, and reduces stretch flangeability. Therefore, the S amount is 0.005% or less, but it is preferable to reduce it as much as possible.

-Al: 0.005-0.1%
Al is used as a deoxidizing element and is an effective element for improving the steel cleanliness. In order to obtain such an effect, the Al amount needs to be 0.005% or more. On the other hand, if the Al content exceeds 0.1%, surface defects are likely to occur, and cost increases. Therefore, the Al content is 0.005 to 0.1%.

N: 0.005% or less N is an element having a strong affinity with Ti, and forms Ti nitride that does not contribute to strengthening. Therefore, if the N amount exceeds 0.005%, a large amount of Ti is required to secure the amount of Ti carbide that contributes to strengthening, resulting in an increase in cost. Therefore, although it is 0.005% or less, it is desirable to reduce it as much as possible.

Ti: 0.08 to 0.20%
Ti is an important element in the present invention, and carbide such as fine TiC or Ti ₄ C ₂ S ₂ having a particle size of less than 10 nm in the ferrite phase during cooling (air cooling) following primary cooling after hot rolling. As a result, it contributes to high strength. In order to achieve TS of 590 MPa or more, at least the Ti content needs to be 0.08% or more. On the other hand, if the Ti amount exceeds 0.20%, it becomes difficult to dissolve coarse Ti carbide during slab heating prior to hot rolling, and fine Ti carbide that contributes to strengthening after hot rolling cannot be obtained. Moreover, the coarse Ti carbide | carbonized_material melt | dissolves at the time of a slab heating is caused, and uniformization of TS in a steel plate is inhibited. Therefore, the Ti content is 0.08 to 0.20%, desirably 0.08 to 0.16%, and more desirably 0.08 to 0.13%.
The balance is Fe and inevitable impurities.

-About formula (1)-(3) Control of the amount of cementite to precipitate is needed to obtain a hot-rolled steel sheet with λ of 100% or more, as will be described later. Therefore, the present invention utilizes the fact that Ti combines with C to produce Ti carbides such as TiC and Ti ₄ C ₂ S ₂ .
Therefore, it is necessary to secure an amount of Ti capable of forming Ti carbide, and Ti ^* defined by the following formula (1) needs to satisfy the following formula (2).
Ti ^* = Ti− (48/14) × N (1)
Ti ^* ≧ 0.08 (2)
Ti ^* represents the amount of Ti that can form Ti carbide.
In order to obtain good stretch flangeability, it is necessary to control the amount of cementite. In the steel of the present invention, the amount of surplus C that does not form Ti carbide is the amount of cementite produced. As the amount of cementite generated increases, the stretch flangeability tends to decrease. To obtain λ100% or more, the value of (C / Ti ^* ) needs to be 0.375 or less. On the other hand, if this value is less than 0.300, the amount of fine Ti carbide produced is insufficient, and a predetermined strength (TS590 MPa or more) cannot be obtained.
That is, (C / Ti ^* ) must satisfy the following formula (3).
0.300 ≦ (C / Ti ^* ) ≦ 0.375 (3)
In the formulas (1) to (3), Ti, N, and C indicate the content (% by mass) of each element.

2) Steel structure Next, the steel structure of the present invention will be described.
In order to achieve TS of 590 to 780 MPa, El of 28% or more, and λ of 100% or more, it is important to have a steel structure mainly composed of a hard ferrite phase. This is because a steel sheet having high strength and high ductility can be obtained by precipitating Ti carbide in the ferrite phase rich in ductility during the progress of ferrite transformation. In order to suppress the precipitation of cementite which adversely affects stretch flangeability, the contained C needs to be fixed as fine Ti carbide. Since cementite is very hard, it becomes the base point for voids during stamping and stretch flange molding. The generated voids grow and connect, and lead to destruction, but in the steel sheet having a steel structure with a ferrite phase area ratio of 95% or more, the particle spacing between cementite is sufficiently wide, so even if cementite is included, The progress of void connection can be slowed down, and the stretch flangeability is good compared to the case where the ferrite area ratio is less than 95%. Furthermore, if the area ratio of the ferrite phase is 95% or more, El can be 28% or more.

  In addition, if the area ratio of a ferrite phase is 95% or more, even if a 2nd phase contains any one or more types of a martensite phase, a bainite phase, and a pearlite phase, the effect of this invention will be impaired. Absent.

In order to obtain a steel sheet with high strength and uniform material, it is necessary to satisfy the condition that the area ratio of the ferrite phase is 95% or more and to make the ferrite grain size and Ti carbide fine and uniform in size. is there. Furthermore, it is necessary to obtain as much Ti carbide as possible. Specifically, if the average ferrite particle diameter is 10 μm or less, the average particle diameter of Ti carbide is 10 nm or less, and the Ti amount precipitated as Ti carbide is 80% or more of Ti ^* (Ti amount capable of forming Ti carbide). For example, it is possible to achieve TS of 590 MPa or more and ΔTS of 15 MPa or less.

3) Manufacturing conditions The manufacturing conditions of this invention are demonstrated.
-Slab heating temperature: 1200-1300 ° C
In order to precipitate fine Ti carbide in the ferrite phase after hot rolling, it is necessary to dissolve coarse Ti carbide precipitated in the slab before hot rolling. For that purpose, it is necessary to heat a slab at 1200 degreeC or more. On the other hand, heating above 1300 ° C. increases the production of scales and leads to a decrease in yield. Therefore, the heating temperature of a slab shall be 1200-1300 degreeC.

-Hot rolling finishing temperature: 900 ° C. or higher Since the Mn content of the austenite former is small, the Ar ₃ point is relatively high. Specifically, if the finishing temperature is lower than 900 ° C., it causes coarsening of ferrite grains and abnormal structure, resulting in a decrease in strength and material uniformity. Therefore, the finishing temperature is set to 900 ° C. or higher.

・ Cooling start time after hot rolling: within 2 seconds ・ Average cooling rate during primary cooling after hot rolling: 30 ° C./s or more Coarse if the time from hot rolling to the start of primary cooling exceeds 2 seconds Since strong ferrite grains and coarse Ti carbide are generated, strength and material uniformity are lowered. Therefore, the cooling start time after rolling is set to be within 2 seconds. For the same reason, the average cooling rate during primary cooling after hot rolling is 30 ° C./s or more.

-Cooling stop temperature for primary cooling: 650-750 ° C
It is necessary to stop the primary cooling in a temperature range of 650 to 750 ° C. to promote ferrite transformation and formation of fine Ti carbide during subsequent cooling (air cooling). When the cooling stop temperature is less than 650 ° C., ferrite is not sufficiently generated, and an area ratio of 95% or more cannot be secured, and the amount of Ti precipitated as Ti carbide cannot be made 80% or more of Ti ^*. . On the other hand, if the cooling stop temperature exceeds 750 ° C., ferrite grains and Ti carbides are coarsened, and it becomes difficult to achieve a ferrite particle size of 10 μm or less and an average particle diameter of Ti carbides of 10 nm or less. Therefore, the primary cooling stop temperature is set to 650 to 750 ° C.

Air cooling time after primary cooling: 5 to 20 seconds When the air cooling time is less than 5 seconds, the ferrite phase is not sufficiently formed, the area ratio of the ferrite phase is 95% or more, and the amount of Ti precipitated as Ti carbide is 80 of Ti ^* . % Or more becomes difficult. If the air cooling time exceeds 20 seconds, ferrite grains and Ti carbides are coarsened, and it becomes difficult to achieve a ferrite particle size of 10 μm or less and an average particle diameter of Ti carbides of 10 nm or less. Therefore, the air cooling time after the primary cooling is 5 to 20 seconds.

Secondary cooling conditions: average cooling rate of 30 ° C./s or more Air cooling to maintain a ferrite particle size of 10 μm or less and an average particle size of Ti carbide of 10 nm or less obtained by a combination of primary cooling after hot rolling and an air cooling process Until the subsequent winding, it is necessary to perform secondary cooling at an average cooling rate of 30 ° C./s or more.

-Winding temperature: 560 degrees C or less In the manufacturing method of this invention, the state of a steel plate structure and Ti carbide is determined before winding, and winding processing will be performed after that. However, when the coiling temperature exceeds 560 ° C., the Ti carbide becomes coarse and the strength decreases. Therefore, the coiling temperature is 560 ° C. or less. In addition, from the viewpoint of securing a good steel plate shape, the winding temperature is preferably 350 ° C. or higher.

  Normal conditions can be applied to other manufacturing conditions. For example, steel having a desired component composition is manufactured by melting in a converter or electric furnace and then performing secondary refining in a vacuum degassing furnace. The subsequent casting is preferably performed by a continuous casting method from the viewpoint of productivity and quality. After casting, hot rolling is performed according to the method of the present invention. After hot rolling, the characteristics of the steel sheet are not impaired even if the scale is attached to the surface or the scale is removed by pickling. Further, after hot rolling, temper rolling, hot dip galvanizing, electrogalvanizing, and chemical conversion treatment can be performed. Here, zinc-based plating is plating mainly composed of zinc and zinc (containing 90% or more of zinc), and is alloyed after plating containing alloy elements such as Al and Cr in addition to zinc or zinc-based plating. It is the plating which processed.

Steels A to H having chemical components (compositions) shown in Table 1 were melted in a converter and made into slabs by a continuous casting method. These steel slabs were heated to 1250 ° C. and coiled hot-rolled steel No. 2 having a thickness of 3.2 mm under the hot rolling conditions shown in Table 2. 1-18 were produced.
In addition, the underline in Table 1 and Table 2 shows that it remove | deviates from the conditions of this invention.

  Then, after the pickling, after trimming the innermost winding and the outermost winding of the coil and both ends 10 mm in the coil width direction, the coil is divided into 20 equal parts in the longitudinal direction of the coil and 8 equal parts in the width direction. JIS No. 5 tensile test specimens were collected from the position of 189 points including the coil end in parallel to the rolling direction, and subjected to a tensile test at a crosshead speed of 10 mm / min in accordance with JIS Z 2241 to obtain an average tensile strength. As a measure of (TS), total elongation (El), and material uniformity, the variation of TS in the coil after trimming, that is, the standard deviation ΔTS of TS was obtained.

  In addition, a test piece for hole expansion test was collected from the position of 189 points, and a hole expansion test was performed in accordance with the iron standard JFST1001, and an average hole expansion ratio λ was obtained. The area ratio of the ferrite phase and the second phase occupying the entire structure is obtained by taking a specimen for a scanning electron microscope (SEM) from the position of 189 points, corroding the plate thickness section parallel to the rolling direction, and then corroding nital. Observe 10 views of SEM photograph at 1000x magnification near the center of plate thickness, identify phases other than ferrite phase such as ferrite phase and martensite phase by image processing, and ferrite phase such as ferrite phase and martensite phase by image analysis processing The area of the other phases was measured and determined as a ratio (percentage) to the area of the observation field, and the area ratio of the ferrite phase was the minimum value of 189 points.

The average ferrite particle diameter was determined by the cutting method from the above 10 SEM photographs. That is, three vertical and horizontal lines are drawn on each SEM photograph to obtain the ferrite grain section length, and the obtained grain section length multiplied by 1.13 (corresponding to the ASTM nominal grain diameter) And 10 fields of view were averaged to obtain an average ferrite particle size.
In addition, the maximum value of the average ferrite particle diameter calculated | required in said 189 point position by said method was shown in Table 3 mentioned later. The average particle diameter of Ti carbide was measured by dividing the thin film by the twin jet method from the center of the plate thickness at 21 points in the coil width direction, including the coil end, and the transmission electron microscope. Observation was performed using (TEM), and the particle size of 3000 or more Ti carbides was measured by image analysis, and the average value was obtained. The amount of Ti deposited as Ti carbide was about 0.2 g in 21% sampling position where TEM observation was performed in 10% AA electrolyte (10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol). Constant current electrolysis was performed at a density of 20 mA / cm ² to extract Ti carbide, and the extraction amount was analyzed.

The results are shown in Table 3. An underline in the table indicates that the condition of the present invention is not satisfied.
In Table 3, steel plate No. 1-3, 11 and 13 are examples of invention, and steel plate No.1. 4-10, 12, and 14-18 are comparative examples.
In addition, although the ferrite area ratio is described in Table 3, the phase other than ferrite was a pearlite or bainite phase.

  Invention Example No. 1-3, 11 and 13 all have TS of 590 to 780 MPa, El of 28% or more and λ of 100% or more, high strength, excellent ductility and stretch flangeability, and TS variation ΔTS of 15 MPa. It is as follows, strength variation is small in the coil, and material uniformity is excellent.

  On the other hand, no. In No. 4, the steel type is A and the composition is within the scope of the present invention, but the time until the start of the primary cooling time after rolling is 3.0 seconds, and the production conditions are outside the present invention exceeding 2 seconds. For this reason, the ferrite grain size is coarsened to 11 μm, the strength is low at 586 MPa, ΔTS is 28 MPa, and the material uniformity is poor.

  Comparative Example No. 5 is steel grade A and the composition is within the scope of the present invention, but the average cooling rate during the primary cooling after rolling is 20 ° C./s, less than 30 ° C./s, and the production conditions are outside the scope of the present invention. is there. For this reason, no. 4, the ferrite grain size is as coarse as 12 μm, TS is 565 MPa and the strength is low, ΔTS is 31 MPa, and the material uniformity is inferior.

Comparative Example No. No. 6 is steel grade A and the composition is within the scope of the present invention, but the cooling stop temperature of the primary cooling after rolling is 600 ° C., which is less than 650 ° C., and the production conditions are outside the present invention. Therefore, the ferrite phase is not sufficiently formed, the ferrite area ratio is as low as 76%, the amount of Ti precipitated as Ti carbide is 76% of Ti ^* , does not reach 80%, and El 26% and λ are as low as 78%, especially ΔTS is 35 MPa, which is inferior in material uniformity.

Moreover, No. of the comparative example. In No. 7, the steel type is A and the composition is within the scope of the present invention, but the cooling stop temperature of the primary cooling after rolling is 800 ° C., which exceeds 750 ° C., and the production conditions are outside the present invention. For this reason, the average particle diameter of Ti carbide is 12 nm, exceeds 10 nm, and the amount of Ti precipitated as Ti carbide is 64% of Ti ^* , which is less than 80%. Further, the ferrite area ratio is 61%, which is lower than 85%. For this reason, TS is as low as 532 MPa, ΔTS reaches 47 MPa, and is inferior in strength and material uniformity. Further, El is 27%, λ is 64%, and ductility and stretch flangeability are also inferior.

  Comparative Example No. In No. 8, the steel type is A and the composition is within the range of the present invention, but the air cooling time after the primary cooling is 25 seconds, which exceeds 20 seconds, and the production conditions are outside the present invention. For this reason, Ti carbide has an average particle diameter of 11 nm and is coarse. For this reason, TS is 578 MPa and ΔTS is 21 MPa, which is slightly inferior in strength and material uniformity.

  Comparative Example No. No. 9 is steel grade A and is within the scope of the present invention, but the average cooling rate of secondary cooling is 20 ° C. and less than 25 ° C., which is outside the production conditions of the present invention. For this reason, the ferrite grain size is 13 μm, which is coarse. For this reason, TS is 574 MPa and ΔTS is 27 MPa, which is slightly inferior in strength and material uniformity.

  Comparative Example No. In No. 10, the steel type is A and within the scope of the present invention, but the coiling temperature is 600 ° C., which exceeds 560 ° C., which is outside the production conditions of the present invention. The average particle diameter and ferrite particle diameter of Ti carbide are larger than 10 nm and 10 μm, respectively, and are coarsened. For this reason, TS is 564 MPa and ΔTS is 22 MPa, which is slightly inferior in strength and material uniformity.

  Invention Example No. 11 and Comparative Example 12 both have a steel grade of B and the composition is within the scope of the present invention. No. 11 is 910 ° C., which satisfies the production conditions of the present invention. 12 is 880 degreeC and is outside the manufacturing conditions of this invention. For this reason, Comparative Example 12 has a coarsened ferrite particle diameter of 11 μm, resulting in poor strength and material uniformity.

Comparative Example No. In No. 14, the steel type is D, the C content is 0.019%, the (C / Ti ^* ) value is 0.187, and the composition deviates from the conditions of the present invention. For this reason, TS is 549 MPa and the strength is low.
Comparative Example No. In No. 15, the steel type is E, the C content is 0.077%, the (C / Ti ^* ) value is 0.806, and the composition deviates from the conditions of the present invention. For this reason, λ is 67%, which is inferior in moldability.

  Comparative Example No. In No. 16, the steel type is F, the Si amount is 0.56%, and the composition deviates from the condition of the present invention (0.1% or less). For this reason, the ferrite grain size is 11 μm, exceeding 10 μm, ΔTS is 25, and the material uniformity is inferior.

Comparative Example No. In No. 17, the steel type is G, the amount of Mn is 1.25%, and the composition deviates from the condition of the present invention (less than 0.80%). Moreover, the ratio of Ti amount precipitated as Ti carbide to Ti ^* amount is also as low as 71%, which is lower than the conditions of the present invention. Therefore, the ferrite area ratio is low, ΔTS is 18 MPa, the material uniformity is inferior, El is 26%, λ is 86%, and the ductility and stretch flangeability are also inferior.

Comparative Example No. In No. 18, the steel type is H, the Ti amount is 0.075%, and the composition deviates from the conditions of the present invention (0.08 to 0.16%). Further, Ti ^* is 0.060, less than 0.08, and (C / Ti ^* ) also exceeds 0.603 and 0.375, both of which are outside the conditions of the present invention. For this reason, TS is 574 MPa and the strength is inferior.

  As described above, in the present invention, it is possible to obtain a hot-rolled steel sheet having a TS of 590 to 780 MPa, an El of 28% or more, a λ of 100% or more, and a ΔTS of 15 MPa or less, ductility (elongation characteristics) and elongation. It can be seen that it has excellent flangeability and excellent material uniformity.

Claims (2)

% By mass
C: 0.020 to 0.065%
Si: 0.1% or less Mn: 0.40 to less than 0.80% P: 0.030% or less S: 0.005% or less Ti: 0.08 to 0.20%
Al: 0.005 to 0.1%
N: 0.005% or less
And the balance is composed of Fe and inevitable impurities, and Ti ^* defined by the following formula (1) has a steel component satisfying the following formula (2) and formula (3), and has a steel structure. Is a ferrite phase with an area ratio of 95% or more and the balance is at least one of a pearlite phase, a bainite phase, and a martensite phase, and the ferrite has an average ferrite particle size of 10 μm or less and precipitates in the steel. The average particle size of the Ti carbide is 10 nm or less, and the amount of Ti precipitated as Ti carbide is 80% or more of Ti ^* , which is excellent in ductility, stretch flangeability and material uniformity. Strength hot-rolled steel sheet.
Ti ^* = Ti− (48/14) × N (1)
Ti ^* ≧ 0.08 (2)
0.300 ≦ C / Ti ^* ≦ 0.375 (3)
Here, Ti, N, and C in the formula indicate the content (% by mass) of each element. The steel slab having the steel component according to claim 1 is heated in a range of 1200 to 1300 ° C, then hot-rolled at a finishing temperature of 900 ° C or higher, and 30 ° C / s or higher within 2 seconds after the hot rolling. The cooling is started at a cooling rate of 650 to 750 ° C., followed by a cooling process of 5 to 20 seconds, followed by cooling at a cooling rate of 30 ° C./s or more and 560 ° C. The steel structure is coiled in the following, and the steel structure is a ferrite phase with an area ratio of 95% or more and the balance is at least one of a pearlite phase, a bainite phase and a martensite phase, and the ferrite The average ferrite particle diameter is 10 μm or less, the average particle diameter of Ti carbide precipitated in steel is 10 nm or less, and the amount of Ti precipitated as Ti carbide is 80% or more of Ti ^* , Stretch flange And the method of producing a high strength hot-rolled steel sheet excellent in material homogeneity.