JP4905147B2 - Thin high tensile hot-rolled steel sheet and manufacturing method thereof - Google Patents

Description

The present invention is a thin high-tensile hot-rolled steel sheet having a tensile strength of 980 MPa or more and a plate thickness of 2.5 mm or less, which is suitable as a material for automobile members, and is particularly thin and high-tensile hot-rolled with excellent stretch flangeability after processing. It is related with a steel plate and its manufacturing method.

  In recent years, from the viewpoint of improving fuel efficiency that leads to environmental conservation, there has been a strong demand for thin steel sheets for automobiles with high strength and thinness, and there is an increasing demand for those exceeding 980 MPa. There is also a strong demand for thin objects with a thickness of 2.5 mm or less from the viewpoint of thinning the steel sheet to make it lighter. Further, many automotive members have complicated shapes obtained by press molding, and materials that are high in strength and excellent in workability are required.

  In recent years, in order to obtain parts with complex shapes and high strength, steel plates with good hardenability are used, and the steel plates are pressed in a state of being heated to a high temperature, and the steel plates are cooled with a mold simultaneously with the press processing. Therefore, a technique for transforming martensite is also used (die quench method). However, this processing method not only has low productivity, but also the problem of scale formation on the surface of the steel sheet when it is heated, and once it is processed into parts, the structure is martensite, so the subsequent processing is completely There are various problems such as inability to do so. Therefore, there is an urgent need to develop a steel sheet having high strength and excellent workability (for example, stretch flangeability) after processing.

  Conventionally, various steel sheets of this type have been proposed. For example, Patent Document 1 proposes a steel sheet excellent in stretch flangeability in which a bainitic ferrite structure having a high dislocation density is generated. However, this steel sheet has a drawback that it has poor elongation because it contains a bainitic ferrite structure having a high dislocation density. In addition, strong cooling on the run-out table is inevitable for the production of bainitic ferrite, and there is a problem with strip runnability on the run-out table when producing thin products, so thin products with a thickness of 2.5 mm or less are produced. It is unsuitable to do.

  Patent Document 2 proposes a steel sheet excellent in stretch flangeability in which most of the structure is polygonal ferrite and precipitation strengthening and solid solution strengthening are centered on TiC. However, it is difficult to increase the tension to 980 MPa or more with generally well-known precipitates used in this steel sheet. That is, when a large amount of Ti is added to increase the strength of 980 MPa or more, precipitates having a large size are likely to be generated, and it is difficult to obtain the intended strength and the characteristics are likely to become unstable. . Moreover, the slab heating temperature necessary for dissolving TiC increases with an increase in the amount of Ti added, and manufacturing is likely to be difficult with normal equipment.

  Patent Document 3 proposes a steel sheet having an acicular ferrite structure in which fine TiC and / or NbC are precipitated and having excellent stretch flangeability. However, like this steel sheet proposed in Patent Document 1 described above, this steel sheet is a structure having a high dislocation density called acicular ferrite, so that sufficient elongation is not obtained.

Patent Document 4 proposes a hot-rolled steel sheet having a ferrite having an average particle diameter of 1 to 5 μm as a main phase and precipitation strengthened with V carbonitride having an average particle diameter of 50 nm or less. Here, in order to finely precipitate the V precipitate, winding at a low temperature is necessary. As a result, it is difficult to increase the amount of the precipitate, and there is a limit to strengthening. For this reason, in the technique of Patent Document 4, it is necessary to combine with the ferrite fine graining as described above for high strength. Here, in the technique of the cited document 4, in order to make ferrite finer, at the time of finish rolling, rolling is performed at ₃ or more points of Ar in a rolling stand one stage before the end of the tandem rolling mill row, and then 50 ° C./second or more. After cooling to a temperature of “Ar ₃ points−50 ° C.” or less at an average cooling rate of 20% or less, it is necessary to apply a reduction of 20% or less in the final stand, which is difficult to manufacture in a normal production line. Moreover, in this steel plate, since generation | occurrence | production of pearlite etc. is accept | permitted, there exists a possibility that elongation and stretch flangeability may fall.

  Moreover, the technique disclosed by patent document 5 and patent document 6 is developed as a technique of obtaining an ultra high strength steel plate. This technique is a technique in which fine carbides composed of C, Ti, and Mo are dispersed in a ferrite single phase to obtain an ultra-high-strength steel sheet having excellent elongation and stretch flangeability. However, like the technique of the above-mentioned patent document 2, in this technique, when a large amount of C or Ti is added to obtain a tensile strength of 980 MPa or more, the slab is not heated at a normal slab heating temperature (about 1150 to 1250 ° C.). In some cases, the deposited TiC or the like cannot be completely dissolved. That is, in order to completely dissolve TiC or the like in order to obtain high strength, a higher temperature may be required, which may make manufacture difficult, and even if possible, a heavy load is placed on the equipment.

In any of the examples, there is no mention of high strength and further workability after press working.
JP-A-6-172924 Japanese Patent Laid-Open No. 6-200351 JP-A-7-11382 JP 2004-143518 A JP 2003-89848 A JP 2002-322539 A

The present invention has been made in view of such circumstances, and is excellent in stretch flangeability after processing, suitable for applications in which the cross-sectional shape in press molding is complicated, such as an automobile member, and is easier to manufacture than conventional ones. such have a strength above 980 MPa, and an object thereof is to provide the following thin high strength hot-rolled steel sheet thickness 2.5mm and method of manufacturing such a thin high strength hot rolled steel sheet. Moreover, an object of this invention is to provide the manufacturing method which can reduce such an equipment burden and can obtain such a thin high-tensile-strength hot-rolled steel plate.

As a result of intensive studies to achieve the above object, the present inventors have obtained the following knowledge.
(A) When the structure has a low dislocation density and is strengthened with fine precipitates, the strength-elongation balance is improved.
(B) When the ferrite single phase structure is substantially formed and strengthened with fine precipitates, the strength-elongation balance is improved.
(C) When C, Ti, Mo, and V are added and the addition balance is appropriately controlled, carbides in which these are combined are finely precipitated.
(D) When the proportion of V in the composite precipitate is lowered, the precipitate is coarsened, so that the elongation is lowered.
(E) Steel added with V dissolves carbides at a lower temperature than steel added only with Ti and Mo, and fine precipitates effective for strengthening can be obtained efficiently.
(F) In addition to fine dispersion of carbide, addition of a small amount of Ta is effective for improving stretch flangeability after processing.

The present invention has been completed based on these findings, and provides the following (1) to ( 4 ).
(1) By mass%, C: more than 0.06 to 0.24%, Si ≦ 0.3%, Mn: 0.5 to 2.0%, P ≦ 0.06%, S ≦ 0.005% , Al ≦ 0.06%, N ≦ 0.006%, Mo: 0.05 to 0.5%, Ti: 0.03 to 0.2%, V: more than 0.15 to 1.2%, Ta : Containing 0.0010 to 0.0050%, the balance being Fe and inevitable impurities, the C, Ti, Mo, V content has a component composition that satisfies the following formula (I), and in area ratio 95% or more is ferrite, and carbides containing Ti, Mo and V having an average particle size of less than 10 nm are dispersed and precipitated, and the carbides containing Ti, Mo and V are Ti, Mo, V expressed in atomic%. but the plate with the features and tensile strength than 980MPa to have an average composition satisfying the V / (Ti + Mo + V ) ≧ 0.3 The following thin high-strength hot-rolled steel sheet 2.5mm.
0.8 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} ≦ 1.5 (I)
(However, C, Ti, Mo, and V represent mass% of each component)
(2) In the thin high-strength hot-rolled steel sheet of the above (1), thin high-strength hot rolled steel sheet characterized by having a hot-dip galvanized coating on the surface.
( 3 ) By mass%, C: more than 0.06 to 0.24%, Si ≦ 0.3%, Mn: 0.5 to 2.0%, P ≦ 0.06%, S ≦ 0.005% , Al ≦ 0.06%, N ≦ 0.006%, Mo: 0.05 to 0.5%, Ti: 0.03 to 0.2%, V: more than 0.15 to 1.2%, Ta : A steel slab containing 0.0010 to 0.0050%, with the balance being Fe and inevitable impurities, and having a component composition in which the C, Ti, Mo, and V contents satisfy the following formula (I): Hot rolling is performed at a finish rolling finish temperature of 880 ° C. or higher and a coiling temperature of 570 ° C. or higher. The area ratio is 95% or more of ferrite, and carbides containing Ti, Mo, and V having an average particle size of less than 10 nm are dispersed. The carbide containing Ti, Mo, and V precipitates and Ti, Mo, V expressed in atomic% is V / (T + Mo + V) ≧ 0.3 has an average composition satisfying the tensile strength thin high-strength hot rolled steel sheet manufacturing method, characterized in that to obtain the following thin high strength hot-rolled steel sheet thickness 2.5mm above 980 MPa.
0.8 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} ≦ 1.5 (I)
(However, C, Ti, Mo, and V represent mass% of each component)
(4) above (3) of the thin high-strength hot-rolled steel sheet Oite the manufacturing method, a thin high-strength hot rolled steel sheet characterized by having a hot-dip galvanized coating on the surface of the steel sheet after the hot rolling Production method.

In the steel sheet of the present invention having a tensile strength of 980 MPa or more, the carbide containing Ti, Mo, and V having an average particle size of less than 10 nm is about 5 × 10 ⁵ or more per 1 μm ³ , and a higher strength is required. It is considered that about 1 × 10 ⁶ or more per 1 μm ³ is dispersed and precipitated.

  According to the present invention, V is added in an appropriate balance in addition to Ti and Mo, and fine carbides containing Ti, Mo and V are substantially dispersed and precipitated in the ferrite single phase structure, and a trace amount of Ta is further added. By adding, a high-tensile steel sheet excellent in stretch flangeability after processing can be obtained.

  In addition, according to the present invention, the steel pieces containing C, Ti, Mo, V are hot-rolled under appropriate conditions, and fine carbides containing Ti, Mo, V are dispersed and precipitated. It is easy to secure the amount (number), and the steel can be strengthened without requiring special equipment for increasing the slab heating temperature, for example.

  Hereinafter, the present invention will be specifically described by dividing it into a metal structure, a chemical component, and a production method.

[Metal structure]
The high-tensile steel sheet according to the present invention has a substantially ferrite single-phase structure, and carbides including Ti, Mo, and V are precipitated.

・ Substantially ferrite single phase structure The reason why the matrix is made substantially ferrite single phase structure is that ferrite with low dislocation density is effective for improving elongation and single phase structure for improving stretch flangeability. This is because the effect is remarkable particularly in a ferrite single-phase structure rich in ductility. However, the matrix does not necessarily have a complete ferrite single phase structure, and may be substantially a ferrite single phase structure. That is, trace amounts of other phases or precipitates are acceptable. Preferably, the ferrite may have an area ratio of 95% or more, and in the present invention, the ferrite single phase structure substantially means 95% or more of ferrite in the area ratio. Note that ferrite having a high dislocation density such as bainitic ferrite and acicular ferrite is not included in the ferrite phase in the present invention, and is handled as another phase.

-Carbide containing Ti, Mo, and V Carbides containing Ti, Mo, and V are effective in strengthening steel because they become finer and the amount of precipitation is easily secured. Conventionally, it has been mainstream to use TiC containing no Mo or V as a precipitate. However, since Ti has a strong tendency to form precipitates, it is likely to be coarsened and the effect on strengthening is reduced. Therefore, precipitates are required until the workability is deteriorated in order to obtain a necessary strengthening amount.

  On the other hand, as disclosed in the above-mentioned Patent Document 5, the precipitates are refined by adding Mo to Ti, and a certain degree of strength can be achieved. However, in order to obtain a tensile strength of 980 MPa or more with only carbides containing Ti and Mo, if Ti corresponding to this level is added, a high temperature exceeding the heating temperature before general hot rolling is required as described above. In order to increase the temperature, for example, special equipment is required, which increases the cost. On the other hand, when only V is added to Ti, sufficient precipitate refinement cannot be obtained.

  On the other hand, composite carbide containing Ti, Mo, and V precipitates finely, and since it is easy to ensure the amount (number) of precipitates, the steel can be strengthened without degrading workability. The inventors have discovered.

  This is because Mo and V, particularly Mo, has a tendency of precipitate formation (carbide formation tendency) to be weaker than that of Ti, so that composite carbide containing Ti, Mo, and V does not become coarse precipitates that do not contribute to strengthening. It is presumed that it can exist stably and finely, and can be effectively strengthened with a relatively small addition amount that does not deteriorate the workability.

  On the other hand, the combination of V and C has a very low melting temperature, and can be easily dissolved at a normal heating temperature even when added in a relatively large amount in order to obtain a high strength of 980 MPa or more. However, when V is added alone, the deposition rate of V becomes low. For this reason, it is considered effective to add both Mo and V in addition to Ti in order to deposit a precipitate having a size and amount sufficient to obtain a high tensile strength of 980 MPa or more.

  Conventionally, when a large amount of V is added to steel containing Ti, Mo or the like, the elongation tends to decrease, and the addition of V has been suppressed to a relatively low range. However, as a result of detailed studies by the present inventors, as the amount of V added is increased, the precipitation rate of V is increased, that is, the added V is sufficiently precipitated as carbides, and the carbides are stably fined. It was found that high strength can be achieved while ensuring sufficient elongation.

  In order for the carbide to exist stably and finely, the composition of the carbide influences. Specifically, when Ti, Mo, and V, in which the average composition of carbides is expressed in atomic%, satisfy V / (Ti + Mo + V) ≧ 0.3, the effect of suppressing the coarsening of precipitates increases, Desired fine precipitates can be obtained. Therefore, in the present invention, it is a requirement that carbides containing Ti, Mo, and V are dispersed and precipitated in a range where Ti, Mo, and V expressed in atomic% satisfy V / (Ti + Mo + V) ≧ 0.3. To do. The upper limit of V / (Ti + Mo + V) is preferably limited to about 0.7.

  As the inventors have found, the optimum carbide composition for miniaturization is approximately 1: 1: 2 in terms of atomic ratio of Ti: Mo: V. Therefore, in the average composition of carbides, the atomic ratio of Ti: Mo: V is Ti = 0.6 to 1.4, Mo = 0.6 to 1.4, V = 1.4 to 2.8, It is further desirable to satisfy Ti + Mo + V = 4.

  Further, by setting the average particle size of the composite carbide to less than 10 nm, the distortion around the precipitate becomes more effective as the resistance of dislocation movement, and the steel can be strengthened efficiently. For this reason, in this invention, it is required that the carbide | carbonized_material containing Ti, Mo, and V with an average particle diameter of less than 10 nm has precipitated. More preferably, the average particle size is 5 nm or less.

  In some cases, carbides containing Ti, Mo, and V are precipitated on coarse precipitates that hardly affect the strength. Since it is inappropriate to use such a precipitate as an object of particle size evaluation, the average particle size is measured by excluding precipitates having a particle size exceeding 100 nm.

In the steel sheet of the present invention having a tensile strength (TS) of 980 MPa, a larger number of the composite carbides having an average particle size of less than 10 nm are observed than a conventional TS780 MPa class steel sheet. It is considered that about 5 × 10 ⁵ or more of this composite carbide in the steel of the present invention is dispersed and precipitated per 1 μm ³ based on the estimation based on the data of Patent Document 6. Since the data in the region exceeding TS800 MPa is not disclosed in the above-mentioned Patent Document 6, it is extrapolated to TS980 MPa (logarithm) assuming that a linear relationship is simply established between the logarithm of TS and the logarithm of fine carbide density. .

[Chemical composition]
Hereinafter, each of these components will be described. Hereinafter, “%” represents “% by mass” unless otherwise specified.
・ C: more than 0.06 to 0.24%
C forms carbides and is effective for strengthening steel. However, if it is 0.06% or less, the steel is not sufficiently strengthened, and if it exceeds 0.24%, spot welding becomes difficult, so the C content is preferably more than 0.06 to 0.24%. . More preferably, it is 0.07% or more, and in particular, in order to obtain a tensile strength of 1100 MPa or more, 0.1% or more is desirable. A more preferable C content range is 0.11 to 0.2%.

-Si: 0.3% or less Si has been actively used as an element effective for solid solution strengthening in the past, and is often added to high-strength steel by about 0.4% or more. The Si content is 0.3% or less. This is because if added over 0.3%, C precipitation from the ferrite is promoted, and coarse iron carbide tends to precipitate at the grain boundaries, and the stretch flangeability is lowered. Further, in the present invention, by reducing Si, the rolling load of austenite is reduced, and the manufacture of thin objects is facilitated. That is, if added over 0.3%, rolling of a material of 2.5 mm or less becomes unstable and the plate shape is also deteriorated. For these reasons, the Si content is preferably 0.3% or less. More preferably, it is 0.15% or less, and desirably 0.05% or less. In addition, since extremely reducing Si worsens the manufacturing cost, the practical lower limit without significant cost increase is about 0.001%.

・ Mn: 0.5-2.0%
Mn is preferably 0.5% or more from the viewpoint of strengthening the steel by solid solution strengthening, but if added over 2.0%, segregation occurs, a hard phase is formed, and stretch flangeability deteriorates. For this reason, the Mn content is preferably 0.5 to 2.0%. More preferably, it is 1.0 to 2.0%.

P: 0.06% or less P is effective for solid solution strengthening, but if it exceeds 0.06%, it segregates and deteriorates stretch flangeability, so it is preferable to make it 0.06% or less. Note that extremely reducing P worsens the manufacturing cost. Therefore, a practical lower limit value that does not significantly increase the manufacturing cost is about 0.001%.

S: 0.005% or less S is preferably as small as possible, and if it exceeds 0.005%, the stretch flangeability deteriorates, so 0.005% or less is preferable. In addition, extremely reducing S worsens manufacturing cost. Therefore, a practical lower limit that does not significantly increase the manufacturing cost is about 0.0005%.

-Al: 0.06% or less Al may be added as a deoxidizer. However, if the amount of Al in the steel exceeds 0.06%, elongation and stretch flangeability deteriorate, so 0.06% or less is preferable. There is no particular lower limit, but in order to sufficiently obtain the effect as a deoxidizer, the Al content is preferably 0.01% or more.

N: 0.006% or less N is preferably as small as possible, and if it exceeds 0.006%, coarse nitrides increase and stretch flangeability deteriorates, so 0.006% or less is preferable. In addition, extremely reducing N worsens manufacturing cost. Therefore, a practical lower limit that does not greatly increase the manufacturing cost is about 0.0005%.

Mo: 0.05-0.5%
Mo is an important element in the present invention, and has an effect of suppressing pearlite transformation by adding 0.05% or more. Furthermore, Ti, V and fine precipitates (composite carbides) can be formed, and the steel can be strengthened while ensuring excellent elongation and stretch flangeability. However, if added over 0.5%, a hard phase is formed and stretch flangeability is lowered, so the Mo content is preferably 0.05 to 0.5%. A more preferred lower limit is 0.15%, and a more preferred upper limit is 0.4%.

Ti: 0.03-0.2%
Ti is an important element in the present invention. By forming composite carbide with Mo and V, steel can be strengthened while ensuring excellent elongation and stretch flangeability. However, if it is less than 0.03%, the effect of strengthening the steel is insufficient. If it exceeds 0.2%, the stretch flangeability deteriorates, and the slab heating temperature before hot rolling must be 1300 ° C or higher. Since the carbide does not dissolve, it cannot be effectively precipitated as a fine precipitate even if it is added more than this. Therefore, the Ti content is preferably 0.03 to 0.2%. More preferably, it is 0.08 to 0.2%.

・ V: Over 0.15 to 1.2%
V is an important element in the present invention. As described above, the composition of the carbide has an influence on the fact that the carbide can exist stably and finely. Specifically, when Ti, Mo, and V, in which the average composition of carbides is expressed in atomic%, satisfy V / (Ti + Mo + V) ≧ 0.3, the effect of suppressing the coarsening of precipitates is enhanced. A desired fine precipitate can be obtained. In this regard, as a result of detailed studies by the present inventors, addition of a large amount of C exceeding 0.06% and addition of a large amount of V increase the precipitation efficiency of V, and V / (Ti + Mo + V) ≧ It was found that a precipitate satisfying 0.3 could be obtained. Further, in order to obtain a desired precipitate, as described above, in addition to satisfying V / (Ti + Mo + V) ≧ 0.3 with Ti: Mo: V in which the average composition of the carbide is expressed in atomic%. In the average composition of carbides, the atomic ratio of Ti: Mo: V is Ti = 0.6 to 1.4, Mo = 0.6 to 1.4, V = 1.4 to 2.8, where Ti + Mo + V = 4 It is preferable to satisfy. Also in this case, as described above, C is added in a large amount exceeding 0.06%, and by adding a large amount of V, the atomic ratio of Ti: Mo: V becomes Ti = 0 in the average composition of carbides. 6 to 1.4, Mo = 0.6 to 1.4, V = 1.4 to 2.8, but it was found that a precipitate satisfying Ti + Mo + V = 4 can be obtained.

FIG. 1 shows the relationship between the amount of V added to steel and the precipitation rate of V. Here, the precipitation rate of V means the ratio of V that actually formed precipitates to the added V, and indicates the deposition efficiency of V. The results are as follows: C: 0.11 to 0.15%, Si: 0.02%, Mn: 1.34%, P: 0.01%, S: 0.001%, Al: 0.04 %, N: 0.002%, Ta: 0.0021%, Mo: 0.33%, Ti: 0.15%, V: 0.1% to 0.3% steel, It is obtained using a hot-rolled steel sheet obtained by hot rolling at a finish rolling finish temperature of 920 ° C and a winding temperature of 620 ° C. Here, the amount of C and the amount of V are (C amount, V amount) = (0.11%, so that the atomic ratio of C and (Ti + Mo + V) is substantially constant (about 1.0 to 1.1). 0.1%), (0.13%, 0.2%), and (0.15%, 0.3%). Moreover, the amount of precipitation V of the hot rolled steel sheet is measured by quantitative analysis of the extraction residue,
V precipitation rate (%) = (precipitation V amount (mass%) / V addition amount (mass%)) × 100
As sought.

  As shown in FIG. 1, as the V addition amount increases, the precipitation rate of V, that is, the ratio of V that actually forms precipitates in the added V increases, and V> 0.15% The deposition rate is> 50%, resulting in very good deposition efficiency. In addition, it confirmed that the structure | tissue of these steel plates was a ferrite single phase structure. Moreover, an example of a precipitate when such a good precipitation efficiency is obtained is shown in FIG. The left side of FIG. 2 is a transmission electron microscope (TEM) photograph showing the precipitate. Moreover, the right side of FIG. 2 is a figure which shows the measurement result by the energy dispersive X-ray-spectrometer (EDX) of Ti, Mo, and V in a precipitate. It was confirmed from the position of the X-ray diffraction peak that these precipitates were mainly composed of carbides. The results are as follows: C: 0.15%, Si: 0.02%, Mn: 1.34%, P: 0.01%, S: 0.001%, Al: 0.04%, N: 0.00. 002%, Ta: 0.0021%, Mo: 0.33%, Ti: 0.15%, V: 0.3% steel, and finish rolling finish temperature 920 ° C, coiling temperature 620 ° C It is obtained using a hot-rolled steel sheet produced by hot rolling. In addition, the precipitates are observed by pickling the obtained hot-rolled steel sheet, producing a thin film from the steel sheet, and observing by TEM. The composition of Ti, Mo, and V in the precipitate is equipped in the TEM. Determined from analysis by EDX. In the analysis result of FIG. 2, Ti: Mo: V was 1.1: 0.9: 2.0 by atomic ratio, and thus V / (Ti + Mo + V) was 0.50.

  As a result of further investigation by the inventors based on such experimental results, it was found that carbide was included in the steel in an amount exceeding 0.15% to obtain a very good precipitation efficiency. Ti, Mo, V, whose average composition is expressed in atomic%, satisfies V / (Ti + Mo + V) ≧ 0.3, forms fine composite carbide with Ti, Mo, and has excellent elongation and stretch flangeability It was found that the steel can be strengthened while securing the above. Further, by containing 0.2% or more of V in the steel, the atomic ratio of Ti: Mo: V is Ti = 0.6 to 1.4 and Mo = 0.6 to 1.4 in the average composition of carbides. V = 1.4 to 2.8 However, it has also been found that the condition of Ti + Mo + V = 4 can be stably satisfied, and the tension can be increased more efficiently. Therefore, the content of V is preferably more than 0.15, and more preferably 0.20% or more. However, if the content of V exceeds 1.2%, center segregation appears strongly, leading to a decrease in elongation and toughness, so 1.2% or less is preferable. More preferably, it is 0.8% or less. Accordingly, the V content is preferably more than 0.15 to 1.2%, more preferably 0.2 to 0.8%. A more preferred lower limit is 0.3%. Even when 1.2% of V is contained, if the slab heating temperature is a normal temperature of about 1200 ° C., the carbide is completely dissolved.

  In addition, although the range of the suitable addition amount of Ti, Mo, V is as above-mentioned, Ti = 0.6-1.4 which is the atomic ratio of preferable Ti: Mo: V in the average composition of the target carbide | carbonized_material. , Mo = 0.6 to 1.4, V = 1.4 to 2.8, but it is more preferable to add at an addition amount ratio corresponding to Ti + Mo + V = 4.

  In order to convert the weight% into the atomic ratio, Ti, Mo, and V may be divided by the atomic weight (48, 96, 51), respectively, to obtain the ratio. However, even if the steel composition does not satisfy the above ratio, the atomic ratio in the fine carbide does not immediately deviate from the preferred range.

0.8 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} ≦ 1.5
(However, C, Ti, Mo, and V in the formula represent mass% of each component)
In the present invention, the balance of addition of C, Ti, Mo and V is very important. Theoretically, the atomic ratio between C and (Ti, Mo, V) in the steel is 1, that is, (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51). } = 1, it is expected that carbon is precipitated as a composite carbide without excess or deficiency. However, according to the investigation by the present inventors, the content of C, Ti, Mo, V in the above predetermined range is set. Thus, by setting (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} to 0.8 to 1.5, Ti, Mo, and V become V / (Ti + Mo + V ) It was found that a large amount of carbide having a composition satisfying ≧ 0.3 can be finely dispersed and precipitated in ferrite finely, that is, with an average particle size of less than 10 nm. If (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} is less than 0.8, the precipitate becomes coarse and a strength of 980 MPa or more cannot be obtained stably. , (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} is more than 1.5, C is excessive and pearlite is generated, so that the moldability is lowered. A more preferable range is 0.8 to 1.3. Even when the C content is excessive, the carbide tends to become coarse.

・ Ta: 0.0010 to 0.0050%
In the present invention, Ta is an important element for ensuring further workability after processing. Here, FIG. 3 shows the influence of the amount of Ta on the hole expansion ratio λ after 10% cold working with respect to the hole expansion ratio λ, which is an index of stretch flangeability, as an evaluation of workability. The reason why the cold working rate is set to 10% is based on the assumption that further processing is performed after a strain corresponding to about 10% cold working is applied in the manufacturing process of automotive parts. It is. From FIG. 3, it can be seen that the hole expansion ratio λ increases rapidly when the Ta content is 0.0010% or more, and is saturated at about 0.0050%. Therefore, in order to obtain the above effect, the lower limit of the Ta content is 0.0010%, and adding more than 0.0050% only saturates the effect and increases the cost. The upper limit of was made 0.0050%. Preferably, it is 0.0015 to 0.0040%.

-Other components In a high-tensile steel plate, other carbide forming elements, such as Nb and W, may be added. However, in the case of the present invention, there is a possibility that the optimum balance of Ti, Mo, and V in the carbide may be lost. Therefore, addition of these is avoided, and the content thereof is preferably within a range acceptable as an impurity. In particular, Nb increases the hot rolling load and makes it difficult to produce thin materials, and in the steel composition of the present invention, it may promote the coarsening of carbides and reduce the strength. Therefore, Nb is preferably 0.02% or less, and more preferably 0.003% or less. W is also preferably 0.02% or less, and more preferably 0.005% or less.

  The balance in the chemical composition of the steel sheet of the present invention is iron and inevitable impurities. In addition to the above, unavoidable impurities include Cr, Cu, Sn, Ni, Ca, Zn, Co, B, As, Sb, Pb, Se, and the like. Cr is allowed to be contained in an amount of 1% or less, preferably 0.6% or less, more preferably 0.1% or less. The other elements are allowed to contain 0.1% or less, but more preferably 0.03% or less.

[Production method]
In the present invention, a steel having the above composition is melted to form a steel slab (including ingot, slab, and thin slab), which is hot-rolled at a finish rolling end temperature of 880 ° C. or higher and a winding temperature of 570 ° C. or higher. I do.

  The plate thickness of the steel plate of the present invention, that is, the plate thickness after hot rolling is preferably about 1.4 to 5.0 mm, but particularly for the production of thin materials with a plate thickness of 2.5 mm or less, which has been difficult in the past. The steel plate of the present invention can be applied without problems. Moreover, in manufacturing a thin hot-rolled steel sheet having a tensile strength of 980 MPa or more and a thickness of 2.5 mm or less, the present application deposits a precipitate bearing the strength after rolling. For this reason, steel is soft during rolling, and can be manufactured without particularly increasing the equipment burden related to rolling.

Steel slab heating conditions Steel slabs such as steel slabs may be cooled once and then reheated to a predetermined temperature (so-called slab heating temperature) before hot rolling. Hot rolling may be performed immediately before the temperature becomes lower. Furthermore, before the steel slab is completely cooled, it may be heated to the predetermined temperature for a short time and hot rolled.

  The slab heating temperature is preferably about 1150 to 1280 ° C. in order to re-dissolve the carbide (or not to precipitate it). In the case of the steel composition of the present invention, re-dissolution can be achieved at a slab heating temperature lower than that of conventional steels of similar components (Ti carbide type, Ti-Mo carbide type).

Finish finish rolling temperature: 880 ° C. or higher Finish finish rolling temperature is important for securing elongation and stretch flangeability and reducing rolling load. If it is less than 880 degreeC, a surface layer will become a coarse grain and elongation and stretch flangeability will be impaired. Further, the amount of strain accumulated due to the progress of rolling due to non-recrystallization increases, and the rolling load significantly increases, making it difficult to hot-roll thin materials. For this reason, finishing rolling finish temperature shall be 880 degreeC or more.

  In addition, in the case of the steel composition of this invention, intensity | strength can be ensured with the finish rolling completion temperature lower than the conventional steel (Ti carbide type | system | group, Ti-Mo carbide type | system | group) of a similar component. In addition, for this reason, it is easy to produce a thin object that is difficult with these conventional steels. The upper limit of the finish rolling finish temperature does not need to be specifically determined. However, since the crystal grains become coarse when finished at a high temperature, the strength of the crystal structure is lowered, and extra strengthening with fine carbides or the like is required. From this viewpoint, the rolling end temperature is preferably set to 1000 ° C. or less.

-Winding temperature of 570 ° C or higher In order to obtain a ferrite structure and to ensure sufficient carbide precipitation, the amount of water injection on the run-out table installed between the finishing mill and the winding device is further suppressed. The winding temperature is set to 570 ° C. or higher in order to allow the thin material to pass through stably. In order to ensure the running stability of the steel sheet on the run-out table, 600 ° C. or higher is preferable. In addition, in order to suppress the production | generation of pearlite, it is desirable for the coiling temperature to be 700 degrees C or less.

  By satisfying the above hot rolling conditions for a steel having a predetermined composition, in the average composition of the precipitated carbide, the atomic ratio of Ti: Mo: V is V / (Ti + Mo + V) ≧ 0.3 or the average composition of the carbide. Ti = 0.6 to 1.4, Mo = 0.6 to 1.4, V = 1.4 to 2.8, where Ti + Mo + V = 4 is satisfied and an average particle size of less than 10 nm is achieved.

  The high-tensile steel sheet according to the present invention includes those having a surface subjected to surface treatment or surface coating treatment. In particular, the steel sheet of the present invention can be suitably applied to a hot-dip galvanized steel sheet that forms a hot-dip galvanized film. That is, since the high-tensile steel sheet of the present invention has good workability, good workability can be maintained even when a hot dip galvanized film is formed. Here, the hot dip galvanizing is hot dip plating mainly composed of zinc and zinc (that is, containing about 80% by mass or more), and includes those containing alloy elements such as Al and Cr in addition to zinc. Moreover, even if hot dip galvanizing is performed, alloying treatment may be performed after plating.

Example 1
A steel slab having the chemical composition shown in Table 1 was heated to 1250 ° C., and finished to a plate thickness of 3.5 mm at a finish rolling end temperature of 880 to 930 ° C. by a normal hot rolling process. Thereafter, steel sheets having various structures were manufactured by changing the cooling rate and the coiling temperature at a coiling temperature exceeding 600 ° C. In Table 1, the A value represents the value of (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} in the above formula (I).

  The obtained steel sheet is pickled, and the thin film prepared by sampling from 1/8, 1/4, 3/8, and 1/2 positions of the thickness of the steel sheet is observed with a transmission electron microscope (TEM). At the same time, the size of the precipitate was measured. Note that the structure observation was mainly performed at a magnification of 5000 to 10,000 times, and the precipitate observation was mainly performed at 260000 to 340000 times. The composition of Ti, Mo, and V in the precipitate was determined from the analysis by an energy dispersive X-ray spectrometer (EDX) equipped in the TEM in the observation at 340000 times, and the V ratio (atomic ratio) of the precipitate = V / (Ti + Mo + V) (wherein Ti, Mo, V are atomic%) and the atomic ratio of Ti: Mo: V were determined. Moreover, about the plate | board thickness direction cross-section of the obtained steel plate, structure | tissue observation was performed 1000-5000 times about the plate | board thickness position which performed TEM observation also with the scanning electron microscope (SEM), and the structure | tissue observation result in TEM was confirmed.

  Here, 30 precipitates having a particle size of 100 nm or less were selected at random, and the particle size and the contents of Ti, Mo, and V were measured for each. The particle diameter was determined by image processing using circle approximation based on the observation result of 340000 times, and the arithmetic average of the 30 was used as the average particle diameter. About the value of V ratio and Ti: Mo: V, content of Ti, Mo, and V was calculated | required by said 30 arithmetic average, and it was set as the average composition, and it computed based on this. The average particle size and average composition obtained for the precipitate having a particle size of 100 nm or less were used as the average particle size and average composition of carbides containing Ti, Mo and V.

JIS No. 5 tensile test specimens were collected from the vertical direction of rolling of the obtained steel sheet, and the pickled steel sheet was cold-rolled at a reduction ratio of 10%, and a hole expansion test was performed on the steel sheet after cold working. Pieces were collected. In the hole expansion test, a test piece having a hole punched with a clearance (one side) of 12.5% of the plate thickness is prepared in the center of a 130 mm square steel plate with a punch of 10 mmφ. The hole diameter d (mm) at the time when the crack penetrated the steel sheet was measured, and the hole expansion ratio λ was calculated from the following equation.
λ (%) = {(d−10) / 10} × 100
Table 2 shows the structure, the average particle size of precipitates, the composition of precipitates (V ratio), the atomic ratio of Ti: Mo: V, the tensile strength (TS), the elongation (El), and the hole expansion ratio after processing (λ). Indicates.

  As shown in Table 2, No. of the steel of the present invention. 1 to 5 are all composed of a ferrite structure, the average particle size of the precipitate is less than 10 nm, the V ratio (atomic ratio) of the precipitate is 0.3 or more, and the tensile strength (TS) is 980 MPa or more. It was confirmed to have excellent elongation. At the same time, the stretch flangeability after 10% cold working was also good.

  On the other hand, No. which is a comparative example. In No. 6, since the amount of C and V is small, the amount of precipitates necessary for strengthening the steel is small, and the tensile strength (TS) is less than 980 MPa. No. In No. 7, since the amount of C is too large and the amount of Mo is small, pearlite exceeding 5% in area ratio is generated to form a ferrite-pearlite structure, the precipitate is coarsened, and the elongation is low. No. No. 8 has a large amount of V, precipitates are coarse, and martensite with an area ratio exceeding 5% is generated to form a ferrite-martensite structure, so that the elongation is low. No. No. 9 has a small amount of Ti and V, so that precipitates necessary for strengthening the steel are insufficient and the tensile strength (TS) is less than 980 MPa. No. It can be seen that No. 10 has a low amount of Ta and therefore has a low stretch flangeability after processing.

(Example 2)
Chemical component in mass%, C: 0.15%, Si: 0.02%, Mn: 1.34%, P: 0.01%, S: 0.001%, Al: 0.04%, N : 0.002%, Ta: 0.0022%, Mo: 0.33%, Ti: 0.15%, V: 0.3% steel (A value: (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} = 0.96) was made into a steel slab. Next, hot rolling was performed after heating to the austenite region. After rolling, the product was cooled to the coiling temperature and wound at the coiling temperature. Table 3 shows finish rolling finishing temperature, winding temperature, and plate thickness. In addition, steel No. shown in Table 3 11 to 17 are all the same chemical components. A sample was taken from the widthwise center of the obtained coil, and a JIS No. 5 tensile test piece was taken so that the tensile direction was perpendicular to the rolling direction, and a tensile test was performed. Moreover, the deposit was investigated from the sample extract | collected from the same position by the method similar to Example 1, and the steel structure was also observed. Furthermore, the pickling plate of the obtained steel plate was cold-rolled at a reduction rate of 10%, and a hole-expansion specimen was taken from the steel plate after cold working, and a hole-expansion test was performed. Moreover, the plate shape after rolling was determined visually. The results are also shown in Table 3. In addition, as for the evaluation criteria of the plate shape after rolling, the case where it was visually flat was marked as ◯, and the case where the wave was remarkable was marked as x. Furthermore, the composition of precipitates (V ratio) and the atomic ratio of Ti: Mo: V are shown in Table 3.

That is, Table 3 shows an example in which the thickness, finish rolling finish temperature, and winding temperature are changed in a 1180 MPa class steel plate having the same chemical composition. Steel No. having a finish rolling finish temperature of 880 ° C. or higher and a winding temperature of 570 ° C. or higher is secured. In Nos. 11 to 15, precipitates having an average particle size of less than 10 nm were formed regardless of the plate thickness, and the target tensile strength (TS) and elongation were achieved. Moreover, the stretch flangeability after processing was also good, and the plate shape was also good. These steel sheets were confirmed to have a ferrite single phase structure as a result of structural observation. On the other hand, no. In No. 16, the finish rolling finish temperature was low, so the crystal grains were coarsened in the surface layer portion, and the precipitates were also coarsened, so the target strength was not satisfied and the elongation was low. The plate shape was also wavy. No. Since No. 17 had a low coiling temperature, precipitates necessary for strengthening the steel were insufficient, the tensile strength (TS) did not reach the target, and waving was remarkable. Steel No. In Nos. 11 to 15, the number of precipitates is about 1 × 10 ⁶ per 1 μm ³ . 16 and 17 are estimated to be about 2.5 to 4 × 10 ⁵ .

Example 3
Steels having chemical components shown in Table 4 were hot-rolled at a finish rolling end temperature of 920 ° C. or higher and a winding temperature of 620 ° C. to produce a hot-rolled steel plate having a thickness of 1.6 mm. The heating temperature of the steel slab was 1250 ° C. These hot-rolled steel sheets were pickled, hot-dip galvanized using zinc as a plating bath, and then alloyed to obtain alloyed hot-dip galvanized steel sheets. As in Example 1, the thin film prepared from the obtained steel sheet was observed with a transmission electron microscope (TEM), the size of the precipitate was measured, and Ti, Mo, and V of the precipitate were further measured. The composition was determined from analysis by an energy dispersive X-ray spectrometer (EDX) equipped with a TEM. Moreover, a JIS No. 5 tensile test piece and a hole expansion test piece were collected from these plated steel sheets and subjected to a tensile test. Further, the obtained alloyed hot-dip galvanized steel sheet was subjected to cold working with a cold reduction of 10% by cold rolling, and a hole-expanding test piece was collected from the cold-worked steel sheet and subjected to a hole-expanding test. . In addition, the A value in Table 4 shows the value of (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} of the formula (I) as in Table 1. Table 5 shows the structure, the average particle size of the precipitate, the composition of the precipitate (V ratio), the atomic ratio of Ti: Mo: V, the tensile strength (TS), the elongation (El), and the hole expansion ratio after processing (λ). Indicates.

  As shown in Table 5, No. 1 as an example of the present invention. No. 18 stretches even when hot dip galvanizing is performed, and the stretch flangeability after processing shows a good value. In No. 19, the precipitate was coarsened, and V was not included in the precipitate, so that the elongation was low. Comparative steel No. Since No. 20 has a small Ta content, stretch flangeability after processing was inferior.

  According to the present invention, by adding V in an appropriate balance in addition to Ti and Mo, fine carbides containing Ti, Mo and V are dispersed and precipitated, and a small amount of Ta is added, thereby reducing the rolling reduction rate. A steel plate having a tensile strength of 980 MPa or more and excellent in stretch flangeability after processing corresponding to cold processing of about 10% is obtained. The steel sheet of the present invention is optimal for automotive parts and the like that require high tensile strength and good workability, and by using the steel sheet of the present invention, it contributes to improvement in productivity and quality improvement in the production of automobile parts. it can.

The graph which shows the relationship between the addition amount of V, and precipitation efficiency. An example of the fine carbide | carbonized_material containing Ti, Mo, and V obtained by this invention (Observation result by a transmission electron microscope). The graph which shows the influence of the amount of Ta exerted on the stretch flangeability after performing a 10% cold work on the high-tensile steel plate having a tensile strength of 980 MPa or more.

Claims (4)

In mass%, C: more than 0.06 to 0.24%, Si ≦ 0.3%, Mn: 0.5 to 2.0%, P ≦ 0.06%, S ≦ 0.005%, Al ≦ 0.06%, N ≦ 0.006%, Mo: 0.05 to 0.5%, Ti: 0.03 to 0.2%, V: more than 0.15 to 1.2%, Ta: 0.00. 0010-0.0050%, the balance is Fe and inevitable impurities, the C, Ti, Mo, V content has a component composition that satisfies the following formula (I), 95% or more by area ratio Is a ferrite, and carbides containing Ti, Mo, and V having an average particle size of less than 10 nm are dispersed and precipitated, and the carbides containing Ti, Mo, and V are Ti, Mo, and V expressed in atomic%. thickness 2 in tensile strength than 980MPa characterized by having an average composition satisfying the /(TitasuMotasuV)≧0.3. mm or less of thin high-strength hot-rolled steel sheet.
0.8 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} ≦ 1.5 (I)
(However, C, Ti, Mo, and V represent mass% of each component) The thin high-tensile hot-rolled steel sheet according to claim 1 , which has a hot-dip galvanized film on the surface. In mass%, C: more than 0.06 to 0.24%, Si ≦ 0.3%, Mn: 0.5 to 2.0%, P ≦ 0.06%, S ≦ 0.005%, Al ≦ 0.06%, N ≦ 0.006%, Mo: 0.05 to 0.5%, Ti: 0.03 to 0.2%, V: more than 0.15 to 1.2%, Ta: 0.00. Finishing rolling into a steel slab containing 0010 to 0.0050%, with the balance being Fe and inevitable impurities, and having a component composition in which the C, Ti, Mo, and V contents satisfy the following formula (I) When hot rolling is performed at a temperature of 880 ° C. or higher and a coiling temperature of 570 ° C. or higher , 95% or more of ferrite is an area ratio, and carbides containing Ti, Mo, and V having an average particle size of less than 10 nm are dispersed and precipitated. The carbide containing Ti, Mo, and V has Ti / Mo / V expressed in atomic% as V / (Ti + Mo Has an average composition satisfying V) ≧ 0.3, a tensile strength thin high-strength hot rolled steel sheet manufacturing method, characterized in that to obtain the following thin high strength hot-rolled steel sheet thickness 2.5mm above 980 MPa.
0.8 ≦ (C / 12) / {(Ti / 48) + (Mo / 96) + (V / 51)} ≦ 1.5 (I)
(However, C, Ti, Mo, and V represent mass% of each component) The method for producing a thin, high-tensile hot-rolled steel sheet according to claim 3 , comprising a hot-dip galvanized coating on the surface of the steel sheet after the hot rolling.

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