JP5510025B2 - High strength thin steel sheet with excellent elongation and local ductility and method for producing the same - Google Patents

Description

  The present invention relates to a high-strength thin steel sheet excellent in elongation, hole expansibility, and local ductility suitable for undercarriage parts and structural materials such as automobiles that are mainly used by pressing and a method for producing the same.

High press workability and strength are required for steel plates used in automobile body structures. As a high-strength thin steel sheet having both press workability and high strength, one having a ferrite / martensite structure, a ferrite / bainite structure, or one containing residual austenite in the structure is known.
Especially, the mixed structure steel plate containing retained austenite can achieve both strength and elongation at a high level. In this retained austenitic steel, in order to further increase the elongation, for example, in Patent Document 1, uniform elongation is achieved by controlling two types of ferrites (bainitic ferrite and polygonal ferrite) while ensuring a high fraction of retained austenite. A technique for ensuring the above is disclosed. Patent Document 2 discloses a technique for defining the shape of an austenite phase by an aspect ratio for the purpose of securing elongation and shape freezing property. Furthermore, Patent Document 3 states that higher elongation can be secured by optimizing the distribution of the austenite phase.
On the other hand, the retained austenitic steel causes work-induced martensitic transformation during work hardening, and a large stress concentration occurs at the interface between the ferrite and martensite, so that the local ductility is low. As a technique for suppressing this stress concentration, as disclosed in Patent Documents 4 to 6, there is a technique for precipitation strengthening a ferrite phase in a mixed structure of ferrite and martensite.

However, in order to meet the demands for further weight reduction and complex parts of today's automobiles, a mixed structure steel plate having higher elongation and better local ductility than before is required.
In Patent Document 7, in a structure composed of a ferrite phase and a hard second phase (martensite, retained austenite), a steel in which the ferrite phase is reinforced by precipitating an alloy carbide in the ferrite phase during cooling after hot rolling. The technology related to is disclosed. However, local ductility and punching workability due to the dispersion of precipitates are a concern.

JP 2006-274418 A JP 2007-154283 A JP 2008-56993 A JP 2002-180188 A JP 2002-180189 A JP 2002-180190 A JP 2009-84648 A

R. W. K. Honeycombe: Metall. Trans. A, 7A, (1976), 915.

  The present invention has been made in order to solve the conventional problems, and as shown in Non-Patent Document 1, it occurs mainly by grain boundary diffusion at the phase interface during the transformation from austenite to ferrite. Provided is a high-strength steel sheet excellent in elongation and local ductility, and a method for producing the same, using a mixed structure of precipitation strengthened ferrite and retained austenite whose precipitation distribution is controlled by a precipitation phenomenon (hereinafter referred to as interphase interface precipitation). This is the issue.

Residual austenitic steel with retained austenite in the steel structure is known to have a very high elongation using the TRIP effect despite its high strength. On the other hand, local ductility is a characteristic that depends on the uniformity of the tissue. A mixed structure of a ferrite phase and a martensite phase having a large hardness difference in the structure generally deteriorates the local ductility. For this reason, the residual austenite which produces a process induction martensite during a process also deteriorates a local ductility similarly. Therefore, to achieve high local ductility and high strength at high elongation, it is important to increase the hardness of the ferrite phase.
Until now, strengthening of the ferrite phase using precipitation strengthening has been studied, but in order for the retained austenite phase to remain in the steel, the coiling temperature must be controlled within a very narrow range, and precipitation control is required. Must be performed during the hot rolling cooling process before winding. Until now, techniques related to the size control of precipitates have been studied, but with this alone, the dispersion state of the precipitates cannot be properly controlled, and an effective improvement effect of local ductility cannot be obtained. It was.

  Therefore, as a result of extensive studies, the present inventors have been able to effectively improve the local ductility by controlling the precipitate distribution (array) in addition to the size of the precipitate. We have found that control does not impair the TRIP effect, which is essential to achieve high growth. Further, as a result of intensive investigations on the method for controlling the distribution of the precipitate, it was found that it can be achieved by a combination of slow cooling and rapid cooling in an extremely limited temperature range, and the present invention has been completed.

That is, the high-strength thin steel sheet excellent in elongation and local ductility of the present invention is
(1) In mass%,
C: 0.08% or more, 0.30% or less,
Si: 0.005% or more, 2.0% or less,
Al: 0.010% or more, 2.0% or less,
Mn: 0.3% or more, 3.0% or less,
P: 0.08% or less,
S: 0.010% or less,
N: 0.010% or less,
Furthermore, Nb: 0.01% or more and 0.10% or less, Ti: 0.01% or more and 0.20% or less of one or two kinds of steel composition, the balance iron and unavoidable impurities and , 30% or more of ferrite phase, a high strength thin steel sheet having a steel structure containing 3% or more retained austenite phases,
The carbonitride the ferrite phase are precipitated by phase interfacial precipitation, spacing of the deposition surface of the interphase surface precipitation at 40% or more regions of the ferrite phase Ri der than 60nm or less 20 nm, the interphase interfaces precipitation high strength thin steel sheet average carbonitrides size in the column is excellent in elongation and local elongation, wherein less der Rukoto 6 nm.
( 2 ) The elongation according to (1 ) above, wherein an in-plane precipitate density of the interfacial interface precipitation is 1 × 10 ⁸ pieces / mm ² or more and 5 × 10 ⁹ pieces / mm ² or less. High-strength thin steel sheet with excellent local ductility.

( 3 ) Further, in the steel composition, V: 0.005% or more, 0.10% or less, Mo: 0.02% or more, 0.5% or less, Cr: 0.1% or more, in the steel composition. 5.0% or less, W: 0.01% or more, 5.0% or less of 1 type or 2 types or more, containing the elongation and local ductility according to the above ( 1 ) or (2) Excellent high-strength thin steel sheet.
( 4 ) The steel composition further containing one or more of Ca, Mg, Zr, and REM in mass% in a range of 0.0005% to 0.05% (%) A high-strength thin steel sheet having excellent elongation and local ductility according to any one of 1) to (3 ).
( 5 ) Further, in the steel composition, Cu: 0.04% or more, 2.0% or less, Ni: 0.02% or more, 1.0% or less, B: 0.0003% or more in the steel composition. The high-strength thin steel sheet having excellent elongation and local ductility according to any one of ( 1 ) to ( 4 ) above, containing one or more of 0.007% or less.

( 6 ) A manufacturing method for manufacturing the high-strength thin steel sheet according to any one of (1) to ( 5 ) above, and during cooling after hot rolling, between 780 ° C or lower and 620 ° C or higher. In the temperature range, air cooling is performed for 1.5 seconds or more, and further, cooling is performed so that the average cooling rate in regions other than air cooling in the temperature range of 800 ° C. or less and 600 ° C. or more is 15 ° C./second or more. A method for producing a high-strength thin steel sheet having excellent elongation and local ductility, wherein the steel sheet is wound at a temperature of from ℃ ° C to less than 450 ℃ to form a hot-rolled steel plate.
( 7 ) In the production method for producing the high-strength thin steel sheet described in ( 6 ) above, the slab temperature before hot rolling is set to 1100 ° C. or higher as it is or after re-heating, after continuous casting. The rolling is finished at 1050 ° C. or higher, hot rolling is performed at a hot rolling finishing temperature of Ar ₃ or higher and 970 ° C. or lower , and when cooling after hot rolling, a temperature range of 800 ° C. or higher is further reduced to 10 ° C./second. A method for producing a high-strength thin steel sheet excellent in elongation and local ductility, characterized by cooling at the above average cooling rate.

The high-strength thin steel sheet of the present invention can ensure sufficient strength by precipitation strengthening due to precipitates in the ferrite phase. In addition, by leaving residual austenite, it is possible to obtain high elongation due to the TRIP effect. In addition, by controlling the distribution of precipitates using interfacial interface precipitation, which is the basis of the present invention, it becomes possible to disperse the precipitates uniformly in the ferrite phase in a short time, resulting in a significant increase in local ductility. The resulting effective hole expansion rate can be improved in high-strength steel.
This effect can be continued as long as heat treatment is not performed in a subsequent process that significantly breaks the structure of the hot-rolled sheet. That is, the same effect can be obtained not only in a hot-rolled steel sheet but also in a hot-rolled plated steel sheet using this as a base plate.

It is a figure which compares and shows the hole expansibility of the steel plate of this invention, and the steel plate of a comparative example. It is a figure which compares and shows the local ductility of the steel plate of this invention, and the steel plate of a comparative example.

  The high-strength thin steel sheet of the present invention focuses on the dispersion state of precipitates in the ferrite phase in a mixed structure including a ferrite phase and a retained austenite phase that enhance hole expansibility. It has been found that strength, elongation and local ductility can be achieved at a high level in a specific dispersed state as described.

In the present invention, the structure is a mixed structure containing 30% or more of the ferrite phase and 3% or more of the retained austenite phase. Thus, high strength and high elongation can be achieved by using the TRIP effect of the retained austenite phase including the ferrite phase.
At this time, if the ferrite phase is less than 30%, the strength is ensured but the elongation deteriorates. Further, if the retained austenite phase is less than 3%, the TRIP effect is small and the effect of improving elongation is small. If the structure is a bainite single phase or a ferrite single phase, the elongation deteriorates.

The dispersion of carbonitrides precipitated in the ferrite phase is one of the most important in the present invention. The carbonitride is deposited at least by interfacial precipitation in order to facilitate dispersion control. At this time, the interval between the precipitation surfaces where precipitation occurs periodically is 20 nm or more and 60 nm or less. When it is deposited at a thickness of less than 20 nm, the precipitate is deposited at a high density, so that the elongation is deteriorated or the precipitate density in the precipitation surface is remarkably reduced. On the other hand, if it exceeds 60 nm, the ferrite phase is not sufficiently strengthened and the desired strength cannot be ensured.
Further, in order to sufficiently harden the ferrite phase and improve the local ductility, it is necessary that the region having a plane spacing of 20 nm or more and 60 nm or less is present at 40% or more of the ferrite phase. If it is less than this value, the ferrite phase is not sufficiently cured and the effect of improving the local ductility cannot be sufficiently obtained.

The average size of carbonitride deposited by interfacial interface precipitation is preferably 6 nm or less. If the average size exceeds 6 nm, the density of the precipitates cannot be secured sufficiently, and the strength of the steel decreases.
Further, the density of carbonitride precipitates in the precipitation surface is also important, and is desirably 1 × 10 ⁸ pieces / mm ² or more in order to ensure the strength of the steel sheet. When the precipitate density is less than 1 × 10 ⁸ pieces / mm ² , the precipitation surface is small and the weakest surface is formed, and the local ductility is reduced by causing stress concentration. On the other hand, 5 in × 10 ⁹ pieces / mm ^2, more than to cause a decrease in local ductility large amount of precipitates cause embrittlement of the deposition surface, and 5 × 10 ⁹ pieces / mm ² or less.

  In order to form carbonitride, either Nb or Ti, or both are included as alloy elements. Nb and Ti are elements having a high driving force for precipitation of carbonitrides in the ferrite phase, and it is easy to control the spacing between the precipitation surfaces and the in-plane precipitate density of the precipitation surfaces depending on the amount of addition. Moreover, since the carbonitride by Ti and Nb has high precipitation strengthening ability, it is set as the carbonitride containing these. However, as long as Ti and Nb are contained, even if other carbide generating elements are contained, the driving force is not greatly reduced, so the effect of the present invention is not impaired.

The chemical components of steel used in the high strength thin steel sheet of the present invention will be described below. % Of content is the mass%.
Nb and Ti are effective elements for strengthening steel by precipitating fine carbonitride by interfacial interface precipitation. When Ti and Nb are both less than 0.01% due to interfacial precipitation, the precipitation ability is low, and the targeted precipitation distribution cannot be obtained. On the other hand, even if Nb exceeds 0.10% and Ti exceeds 0.20%, dissolution of precipitates is insufficient in heating before hot rolling, and not only does not give an advantage to the interfacial interface precipitation amount, The undissolved coarse carbonitride deteriorates local ductility.

  Since C is an element that deteriorates hole expansibility, it is desirable that the addition amount be low. However, it is an effective element for producing carbonitride and increasing strength. If the C content is less than 0.08%, the strength cannot be sufficiently increased. On the other hand, if it exceeds 0.30%, the decrease in elongation becomes large. Therefore, the range of C is set to 0.08% or more and 0.30% or less. When the local ductility requirement is high, the upper limit of C is preferably 0.20%.

  Si is an element effective in suppressing the generation of harmful carbides and generating a ferrite phase. However, addition of more than 2.0% lowers the ductility and also reduces the chemical conversion property, so the amount of Si added is made 2.0% or less. In addition, when the chemical conversion property requirement is high, Si is desirably 1.3% or less. Si is added for deoxidation, but if it is less than 0.01%, the deoxidation effect is not sufficient, so the lower limit of Si is preferably 0.005% or more.

  Al is added as a deoxidizer. For this purpose, Al needs to be added in an amount of 0.010% or more. On the other hand, when Al is added, there is an effect of generating a ferrite phase like Si, and it may be added from this point. However, since addition exceeding 2.0% causes embrittlement, the upper limit was made 2.0%. In addition, when the request | requirement of chemical conversion property is high, it is desirable to set it as 1.5% or less.

  Mn is an effective element for enhancing the hardenability and strengthening the steel. If Mn is less than 0.3%, the strength cannot be sufficiently increased. However, if Mn exceeds 3.0%, the hardenability is increased more than necessary, and the ferrite phase cannot be sufficiently secured and interfacial interface precipitation cannot be obtained. Therefore, the amount of Mn added is 3.0% or less. .

  When P is contained in a large amount, it segregates to the grain boundary, so that the local ductility is degraded and the weldability is degraded. Therefore, the upper limit is made 0.08%. In addition, since it will lead to the cost increase at the time of refining to reduce P unnecessarily, it is desirable that the lower limit is 0.001%.

S is an element that forms MnS and significantly deteriorates local ductility and weldability. Therefore, the upper limit is made 0.010%. Moreover, it is desirable that the lower limit is 0.0005% due to the problem of refining costs.
N is effective for precipitating AlN and the like to refine crystal grains, but if N exceeds 0.010%, solute nitrogen remains and ductility decreases, so the upper limit Is 0.010%. In addition, it is desirable that the lower limit is 0.0010% because of cost problems during refining.

  The steel can further contain one or more of V, Mo, Cr, and W. These are elements that generate carbonitrides in combination with Ti or Nb. For this purpose, one or more of V: 0.005% or more, Mo: 0.02% or more, Cr: 0.1% or more, W: 0.01% or more are contained. Good. However, adding V: more than 0.10%, Mo: more than 0.5%, Cr: more than 5.0%, W: more than 5.0% not only saturates the effect of increasing the strength, This will cause a reduction in local ductility. Therefore, V: 0.10% or less, Mo: 0.5% or less, Cr: 5.0% or less, and W: 5.0% or less are set as upper limits.

  The steel can further contain one or more of Ca, Mg, Zr, and REM (rare earth elements) alone or in total from 0.0005% to 0.05%. Ca, Mg, Zr, and REM improve the local ductility and hole expansibility by controlling the shapes of sulfides and oxides. For this purpose, it is preferable to add one or more of these elements alone or in total of 0.0005% or more. However, excessive addition deteriorates workability, so the upper limit was made 0.05%.

  Further, the steel is Cu: 0.04% or more, 2.0% or less, Ni: 0.02% or more, 1.0% or less, B: 0.0003% or more, 0.007% or less, or 1 or 2 More than seeds can be contained. These elements can improve the hardenability and increase the strength of the steel, but Cu: less than 0.04%, Ni: less than 0.02%, B: less than 0.0003%, the hardenability is weak and high temperature. In order to promote ferrite formation, the necessary interfacial precipitation cannot be obtained. On the other hand, addition exceeding this range makes the hardenability too strong and delays the ferrite transformation necessary for interfacial interface precipitation.

  In addition to the above elements, steel contains elements inevitably mixed such as Sn and As, and is made of the remaining iron.

  The high strength thin steel sheet excellent in elongation and local ductility of the present invention is composed of a structure containing 3% or more of retained austenite phase in a ferrite / bainite mixed structure mainly composed of ferrite. Thus, high strength and high elongation can be achieved by using the TRIP effect of the retained austenite phase including the ferrite phase.

If the amount of the ferrite phase is small, the decrease in ductility increases, so it is desirable that the ferrite phase fraction be 30% or more. When there are many ferrite phases, it becomes difficult to concentrate C in the retained austenite phase, and when a large amount of retained austenite remains, the C concentration in the retained austenite phase becomes insufficient, resulting in a decrease in elongation. For this purpose, it is desirable that the ferrite phase is 80% or less and the residual austenite phase is 30% or less.
In addition, if it is 5% or less, even if a martensite exists, the effect of this invention will not be impaired. In addition, pearlite is inevitably included, but if the pearlite is 5% or less, the material is not significantly deteriorated.

Below, the manufacturing method of the high intensity | strength thin steel plate based on this invention is demonstrated.
As a result of intensive studies, the present inventors have found that, when producing the high-strength thin steel sheet of the present invention, it is possible to control the distribution of interfacial interface precipitation by cooling control in a predetermined temperature range after hot rolling. I found it. Specifically, the transformation rate changes greatly in the region of 800 ° C. or lower and 600 ° C. or higher, and this utilizes the fact that the phase interface precipitation distribution changes greatly.

In order to realize the necessary interfacial interface precipitation, it is necessary to perform air cooling in a temperature range of 780 ° C. or lower and 620 ° C. or higher. When the temperature is higher than this, the distance between the precipitation surfaces becomes too large, causing a decrease in strength. On the other hand, when carried out at a temperature lower than this, the transformation rate becomes too high and the interplanar spacing becomes too small or interfacial interface precipitation does not occur, so that local ductility and elongation deteriorate.
When the air cooling time is less than 1.5 seconds, the region where interphase interface precipitation occurs is not sufficient, and the effect of improving the local ductility cannot be obtained. Since pearlite is produced when the air cooling time is long, the air cooling time is preferably 15 seconds or less.

  Furthermore, since the deposition state of interfacial interface precipitation changes greatly in the temperature range of 800 ° C. or lower and 600 ° C. or higher, the cooling rate during this period is very important. The average cooling rate during this period, excluding the air cooling region, needs to be 15 ° C./second or more. If it is less than this speed, the interplanar spacing, precipitation size, and in-plane precipitate density vary greatly between structures, and the hole expandability and local ductility deteriorate due to the nonuniformity of the structure.

In addition, the effect of the interfacial interface precipitation is greatly influenced by the form of the second phase. Specifically, the coiling temperature is set to 200 ° C. to less than 450 ° C., the second phase is bainite, and the residual austenite phase is left, thereby achieving the optimization of the structure and the reduction of the structure hardness difference due to interfacial precipitation. Thus, an excellent high-strength thin steel sheet having excellent elongation and local ductility is obtained.
When the coiling temperature is less than 200 ° C., martensitic transformation occurs and local ductility deteriorates. On the other hand, when the temperature is 450 ° C. or higher, cementite precipitates vigorously, the residual austenite phase cannot remain, and the strength and local elongation decrease due to the generation of the pearlite phase.

In the present invention, as described above, it can be controlled by hot rolling conditions and is not affected by casting conditions. For example, the influence of the difference in casting method and slab thickness is small, and the slab may be manufactured according to a conventional method.
The slab before hot rolling is set to 1100 ° C. or higher as it is after continuous casting or by reheating. On the other hand, when the temperature exceeds 1300 ° C., scale formation becomes large and the surface properties of the steel sheet cannot be made favorable, so that the temperature is desirably 1300 ° C. or lower. If heating temperature is less than 1100 degreeC, melt | dissolution of carbonitride will be inadequate and will cause a fall of intensity | strength and local ductility. Thereafter, precipitation of carbonitride occurs between rough rolling at a low finishing temperature and before finish rolling, causing a decrease in strength, so the rough rolling is finished at 1050 ° C. or higher.

Next, the slab is hot rolled at a finishing temperature of Ar ₃ or higher and 970 ° C. or lower. This is because when the finishing temperature is less than Ar ₃ , (α + γ) two-phase rolling occurs, resulting in a decrease in ductility. When the finishing temperature exceeds 970 ° C., the austenite grain size becomes coarse, the ferrite phase fraction decreases, and the ductility This is because of a decrease.

For the cooling after hot rolling, the average cooling rate in the temperature range of 800 ° C. or higher is 10 ° C./second after finish rolling. If the cooling rate is too low, ferrite transformation and pearlite transformation occur before reaching the temperature range at which interfacial precipitation occurs, and the strength decreases.
After cooling, it is wound up in a temperature range of 200 ° C. or higher and lower than 450 ° C. to obtain a hot rolled steel sheet. In addition, after hot rolling, surface treatment such as plating is performed as necessary.
As described above, the high-strength thin steel sheet of the present invention can be manufactured.

Hereinafter, the present invention will be described in detail based on examples.
Steel having the component composition shown in Table 1 was manufactured, the steel pieces after cooling and solidification were reheated to 1200 ° C, rough rolling was finished at 1080 ° C, and hot rolling was performed under the conditions shown in Tables 2 and 3. Carried out. The average cooling rate up to 800 ° C. was 50 ° C./second.
While observing the structure of the obtained steel sheet, the characteristics were examined.

Observation of the precipitate was performed in STEM mode using a FE-TEM with an acceleration voltage of 200 kV by preparing a thin film sample.
The measurement of the interplanar interfacial precipitation and the occurrence rate of the region having an interplanar spacing of 20 nm to 60 nm (interfacial interfacial precipitation occurrence) are performed by randomly extracting ferrite grains and interphase in the extracted crystal grains. Interfacial precipitation is observed with a STEM to measure the face spacing, and the proportion of the region that satisfies the face spacing is measured. This measurement was performed for at least 50 particles, and the average was calculated to calculate the incidence.

Tensile properties were evaluated by C direction tension of JIS No. 5 tensile test pieces.
The local ductility is determined by performing a tensile test using a JIS No. 5 test piece, measuring the plate thickness t and the plate width w of the cross section of the sample after breaking, and from the original plate thickness t0 and the original plate width w0.
RA =-(ln (t / t0) + ln (w / w0))
For each direction of L, C, 45 °,
Local ductility = (RA (L direction) + RA (C direction) + 2 × RA (45 direction)) / 4
I asked for it.

Of steels A to g, a and b, which are comparative steels, do not satisfy the C upper limit and the lower limit, respectively. c, d, and e do not satisfy the upper limits of S, Si, and Mn, respectively. The amount of addition of Si and Al is below the lower limit for f. g does not satisfy the upper limits of Al and Ti. h exceeds the upper limit of Nb. For i, Nb and Ti are both out of range.
Tables 2 and 3 show the production conditions and characteristics of steel plates A1 to g1 obtained using steels A to g. Among the examples using the inventive steels of Table 1, A2 and B2 have a winding temperature range. The austenite fraction is out of range. D2 has a finishing temperature of Ar3 or lower and an austenite fraction outside the range. In G2, the air cooling time is below the lower limit, and the ferrite fraction and the interphase precipitation rate are low. K2 and K3, respectively, the air cooling start temperature does not satisfy the upper limit and the lower limit, the former is the ferrite fraction, interphase interface precipitation occurrence rate, precipitate size and density, the latter is the ferrite fraction, interphase interface precipitation occurrence rate, Density is out of range. In P2, the cooling rate from 800 ° C. to 600 ° C. is lower than the lower limit, and the precipitate size and density are out of the range.
Although these materials are shown in FIGS. 1 and 2, only the steel plate of the inventive steel has both excellent elongation and local ductility, both of which show extremely high values compared to the steel plate of the comparative steel, The goal has been achieved.

Claims (7)

In mass%
C: 0.08% or more, 0.30% or less,
Si: 0.005% or more, 2.0% or less,
Al: 0.010% or more, 2.0% or less,
Mn: 0.3% or more, 3.0% or less,
P: 0.08% or less,
S: 0.010% or less,
N: 0.010% or less,
Furthermore, Nb: 0.01% or more and 0.10% or less, Ti: 0.01% or more and 0.20% or less of one or two kinds of steel composition, the balance iron and unavoidable impurities and , a high strength thin steel sheet having a steel structure containing 30% or more of ferrite phase and less than 3% of retained austenite phase,
The carbonitride the ferrite phase are precipitated by phase interfacial precipitation, spacing of the deposition surface of the interphase surface precipitation at 40% or more regions of the ferrite phase Ri der than 60nm or less 20 nm, the interphase interfaces precipitation high strength thin steel sheet average carbonitrides size in the column is excellent in elongation and local elongation, wherein less der Rukoto 6 nm. Precipitate density in the plane of the interphase surface precipitation 1 × 10 ⁸ pieces / mm ² or more, excellent in elongation and local elongation of claim 1, wherein the 5 × 10 is ⁹ / mm ² or less High strength thin steel sheet. Further in the steel composition,
V: 0.005% or more, 0.10% or less, Mo: 0.02% or more, 0.5% or less, Cr: 0.1% or more, 5.0% or less, W: 0.01% or more, The high-strength thin steel sheet excellent in elongation and local ductility according to claim 1 or 2 , characterized by containing one or more of 5.0% or less. Further during the steel composition
Ca, Mg, Zr, one or two or more of REM, in mass%, 0.0005% or more, according to any one of claims 1 to 3, characterized in that it contains more than 0.05% High-strength thin steel sheet with excellent elongation and local ductility. Further in the steel composition,
Cu: 0.04% or more, 2.0% or less, Ni: 0.02% or more, 1.0% or less, B: 0.0003% or more, 0.007% or less A high-strength thin steel sheet excellent in elongation and local ductility according to any one of claims 1 to 4 . A manufacturing method for manufacturing the high-strength thin steel sheet according to any one of claims 1 to 5 ,
During cooling after hot rolling, air cooling is performed for 1.5 seconds or more in a temperature range of 780 ° C. or lower and 620 ° C. or higher, and other than air cooling in a temperature range of 800 ° C. or lower and 600 ° C. or higher. A method for producing a high-strength thin steel sheet excellent in elongation and local ductility, wherein cooling is performed so that the average cooling rate of the region is 15 ° C./second or more, and winding is performed at a temperature of 200 ° C. or more and less than 450 ° C. . In the manufacturing method for producing the high strength thin steel sheet according to claim 6 , the slab temperature before hot rolling is set to 1100 ° C or higher as it is or after re-heating after continuous casting, and then rough rolling is performed at 1050 ° C. When the hot rolling finish temperature is Ar ₃ or higher and 970 ° C. or lower and cooling is performed after hot rolling, the temperature range of 800 ° C. or higher is further averaged at 10 ° C./second or higher. A method for producing a high-strength thin steel sheet excellent in elongation and local ductility, characterized by cooling at a speed.

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