JP2007070660A - High strength thin steel sheet having excellent formability, and method for producing the same - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a TRIP type high strength thin steel sheet having a structure more uniform than that of the conventional one, and having excellent formability, and to provide a method for producing the same. <P>SOLUTION: The high strength thin steel sheet having excellent formability has a steel composition comprising, by mass, 0.05 to 0.25% C, ≤2.0% Si, 0.8 to 3% Mn, 0.0010 to 0.1% P, 0.0010 to 0.05% S, 0.0010 to 0.010% N and 0.01 to 2.0% Al, and the balance iron with inevitable impurities, and in which Mn micro-segregation in the range of 1/8t to 3/8t of the sheet thickness t lies in the range satisfying formula (1), and its structure comprises retained austenite having a mean carbon content of ≥0.9% by ≥3%: 0.10≥σ/Mn (1); wherein, Mn denotes its additional amount, and σ denotes the standard deviation in Mn micro-segregation measurement. <P>COPYRIGHT: (C)2007,JPO&INPIT

Description

  TECHNICAL FIELD The present invention relates to a high-strength thin steel sheet excellent in formability, which is suitable for undercarriage parts and structural materials such as automobiles that are mainly pressed and used, and a method for producing the same.

  Known as high-strength thin steel sheets having both formability and high strength are steel sheets having a composite structure having a ferrite / martensite structure, steel sheets containing retained austenite, and the like. A composite steel is a steel sheet in which martensite is dispersed in a ferrite ground, has a low yield ratio, high tensile strength, and excellent elongation characteristics. In order to obtain higher elongation, there is a steel sheet containing retained austenite. Residual austenite is generated in the structure, and this retained austenite induces transformation during processing deformation and exhibits excellent ductility. For example, Patent Documents 1 and 2 disclose high-strength steel sheets having excellent formability of TRIP type made of ferrite, bainite, and retained austenite.

However, in the conventional continuous casting, the average cooling rate in the middle part of the slab (1/4 t position of the slab having the thickness t) was as small as about 0.1 ° C./sec. Mn microsegregation was large. This micro-segregation portion is elongated during rolling to form a Mn band, and this portion has a low Ms point, so that the retained austenite is unevenly distributed in the TRIP type steel sheet. As a result, stress was concentrated at the interface between martensite and ferrite, which was induced by cold working, and fracture was likely to occur. Thus, in the conventional TRIP-type high-strength thin steel sheet, the structure non-uniformity caused by the Mn band has been a factor that hinders formability, particularly local ductility.
JP 2001-329340 A JP 2002-12948 A

  An object of the present invention is to provide a TRIP-type high-strength thin steel sheet having a uniform structure and excellent formability as compared with the conventional one and a method for producing the same.

The high-strength thin steel sheet excellent in formability of the present invention made to solve the above problems is
In mass%
C: 0.05-0.25%, Si: 2.0% or less, Mn: 0.8-3%, P: 0.0010-0.1%, S: 0.0010-0.05%, N: 0.0010 to 0.010%, Al: 0.01 to 2.0%, steel composition consisting of the balance iron and inevitable impurities,
Mn microsegregation in the range of 1 / 8t to 3 / 8t of the plate thickness t is in the range satisfying the formula (1),
It contains 3% or more of retained austenite having an average carbon content of 0.9% or more.
0.10 ≧ σ / Mn (1)
Here, Mn is an addition amount, and σ is a standard deviation in Mn microsegregation measurement.

In the above-described invention, during the steel composition,
Cr: 0.01-5%, Mo: 0.01-5%, Ni: 0.01-5%, Cu: 0.01-5%, Co: 0.01-5%, W: 0.01 Can contain -5% of one or more,
Further during the steel composition
One or more of Ti, Nb, Zr, Hf, Ta, V can be contained alone or in total 0.001-1%,
Further during the steel composition
0.0001-0.0050% of B can be contained,
Further during the steel composition
One or more of Mg, Ca, Y, and REM can be contained in an amount of 0.0001 to 0.5%.

Moreover, the manufacturing method of the high strength thin steel sheet excellent in formability of the present invention is as follows:
A method for producing a high-strength thin steel sheet, wherein the high-strength thin steel sheet according to claim 1 is produced from a slab,
The slab in the middle of cooling after casting is cooled as it is or 1100 ° C after cooling between the liquidus temperature and the solidus temperature at an average cooling rate at 1/4 t of the slab thickness t of 100 ° C / min or more. Reheat to above,
Subsequently, hot rolling was performed at a finishing temperature of 850 to 970 ° C., and thereafter cooling was performed at an average cooling rate of 10 to 100 ° C./sec to a temperature range of 700 to 600 ° C., and then retained in the same temperature range for 1 to 5 seconds. Then, it is cooled again at an average cooling rate of 10 to 100 ° C./sec and wound up at a temperature of 300 ° C. or higher and 450 ° C. or lower to form a hot-rolled steel sheet.

Moreover, the manufacturing method of the high strength thin steel sheet excellent in formability of the present invention is as follows:
A method for producing a high-strength thin steel sheet, wherein the high-strength thin steel sheet according to claim 1 is produced from a slab,
The slab that is in the process of being cooled after casting is cooled as it is or 1100 after cooling the temperature range from the liquidus temperature to the solidus temperature at an average cooling rate at 1/4 t of the slab thickness t of 100 ° C./min or more. Reheat to over ℃,
Hot-rolling is performed at a finishing temperature of 850 to 970 ° C., and then cooled to a temperature range of 650 ° C. or lower at an average cooling rate of 10 to 100 ° C./sec. None,
The hot-rolled steel sheet is subjected to cold rolling with a reduction rate of 40% or more after pickling,
After annealing at a maximum temperature of 0.1 × (Ac ₃ −Ac ₁ ) + Ac ₁ or more and Ac ₃ + 50 ° C. or less, 350 ° C. or more and 500 ° C. or less at an average cooling rate of 0.1 to 200 ° C./sec. The steel sheet is cooled to a temperature range, and subsequently kept in the same temperature range for 10 seconds or more and 1000 seconds or less to form a cold-rolled steel sheet.

In the high-strength thin steel sheet of the present invention, the Mn microsegregation is remarkably smaller than that of the prior art, so that a Mn band in which the Mn segregation is extended in the rolling direction is less likely to occur. Therefore, since the band-like structure resulting from the Mn band can be avoided, the formability is superior to the conventional high-strength thin steel sheet.
The method for producing a high-strength thin steel sheet according to the present invention is to produce a hot-rolled steel sheet with an increased cooling rate during solidification, thereby reducing the Mn microsegregation by making the solidification structure finer than a normal slab. Can do. Therefore, since the Mn band is small and the structure is uniform, it is possible to manufacture a high-strength thin steel sheet that is more excellent in formability than conventional ones.
Moreover, since the manufacturing method of the high-strength thin steel plate of this invention rolls and anneals said hot-rolled steel plate and manufactures a cold-rolled steel plate, Mn microsegregation is smaller than before and a structure | tissue is uniform. Therefore, it is possible to produce a high-strength thin steel sheet that has better formability than before.
In the present invention, casting may be performed by any method as long as the cooling rate during solidification can be higher than 100 ° C./min. For example, in continuous casting, the thickness of the slab may be reduced, in ingot casting, the size of the ingot may be reduced, or a surface layer portion having a high cooling rate may be cut out from a normal slab and used.

The high strength thin steel sheet having excellent formability according to the present invention is characterized in that the microsegregation of Mn in the range of 1 / 8t to 3 / 8t of the sheet thickness t satisfies the formula (1).
0.10 ≧ σ / Mn (1)
Here, Mn is an addition amount, and σ is a standard deviation in Mn microsegregation measurement. The standard deviation σ was obtained from Mn concentration distribution data obtained by performing line analysis in the plate thickness direction on a sample having a plate thickness polished using EPMA (X-ray microanalyzer).

When σ is 0.10 <σ / Mn, variation in Mn concentration is large, and microsegregation of Mn is not sufficiently small. For this reason, since the microsegregation of Mn is extended in the rolling direction to form a relatively large Mn band, the structure cannot be made uniform, which can cause ductility deterioration. Therefore, the microsegregation of Mn must satisfy the relationship of 0.10 ≧ σ / Mn. When the demand for formability is high, it is desirable that the microsegregation satisfies the formula (2). This is because the structure can be made more uniform and the moldability can be improved.
0.05 ≧ σ / Mn (2)
This condition needs to be satisfied in the range of 1 / 8t to 3 / 8t of the plate thickness t with slow cooling.
In addition, a high strength thin steel plate means a high strength thin steel plate or a high strength thin steel plate.

  Further, the high strength thin steel sheet having excellent formability according to the present invention is characterized in that the structure contains 3% or more of retained austenite having an average carbon content of 0.9% or more. That is, the high-strength thin steel sheet contains 3% to 20% of retained austenite metastable in the composite structure of ferrite and bainite. In order to cause the TRIP phenomenon and improve the moldability, 3% or more of retained austenite is essential. On the other hand, if the retained austenite exceeds 20%, a large amount of martensite is present, which may cause a problem in secondary formability. The retained austenite must have an average carbon content of 0.9% or more. If it is less than 0.9%, the stability of austenite is insufficient to reflect the TRIP phenomenon in ductility.

The reason for limiting the chemical components of the high-strength thin steel sheet of the present invention will be described below.
C is an austenite stabilizing element and is an important element for producing retained austenite. C moves from ferrite to austenite in the two-phase coexistence temperature range and the bainite transformation temperature range, and increases its stability. As a result, stable austenite remains after cooling to room temperature, which leads to a large elongation. If the C content is less than 0.05%, retained austenite having an appropriate stability cannot be obtained. On the other hand, if it exceeds 0.25%, a large amount of retained austenite is obtained, but the weldability is lowered. Therefore, the range of C in the present invention is 0.05 to 0.25%.

  Si is an important element for stabilizing the retained austenite. By suppressing the precipitation of carbides during bainite transformation, 0.9% or more of C is concentrated in the untransformed austenite phase, and the Ms point. Is reduced to below room temperature. It is also effective as a deoxidizing component, and in order to exert such effects, it is necessary to add Si in an amount of 0.01% or more. Since the processability also decreases, the upper limit is made 2.0%.

  Mn is necessary for stabilizing austenite and enhancing the hardenability of the steel to increase the strength. For this purpose, it is necessary to add 0.8% or more of Mn. However, if it exceeds 3.0%, the elongation is lowered, and the Mn band becomes prominent and the workability is lowered, so the upper limit is made 3.0%. The upper limit is preferably 2.0% from the viewpoint of securing moldability.

  When P is contained in a large amount, it segregates to the grain boundary, so that the local ductility is deteriorated. In addition, the weldability is deteriorated. Therefore, the upper limit is made 0.1%. In addition, since reducing P unnecessarily leads to a cost increase at the time of refining in the steelmaking stage, the lower limit is made 0.0010%.

  S is an element that forms MnS and significantly deteriorates local ductility and weldability. Therefore, the upper limit is made 0.05%. Further, the lower limit is made 0.0010% due to the problem of refining costs.

  N, like C, contributes to the stabilization of austenite. For this purpose, it is necessary to contain 0.0010% or more. However, if N is contained in excess of 0.010%, ductility and weldability deteriorate, so the upper limit is made 0.010%.

  Al is important as a deoxidizer. It is also an important additive element for promoting bainite. For this purpose, Al needs to be added in an amount of 0.01% or more. On the other hand, even if Al is added excessively, the above effect is saturated and the steel is embrittled, so 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.

  Cr, Mo, Ni, Cu, Co, and W improve the hardenability and increase the strength of the steel, but the effect is small when the content is less than 0.01%. On the other hand, even if added over 5.0%, the effect of increasing the strength is saturated and the ductility is lowered.

  Ti, Nb, Zr, Hf, Ta, and V precipitate fine nitrides and carbides to strengthen the steel, but the effect is small at less than 0.001%. On the other hand, adding over 1% not only saturates the effect, but also reduces ductility.

  B increases the hardenability in a small amount. For this purpose, it is necessary to add 0.0001% or more, but even if added over 0.0050%, the effect is not only saturated but also the ductility is lowered. Further, composite addition with Ti is effective.

  Mg, Ca, Y, and REM (rare earth elements) improve the ductility by controlling the shape of sulfides and oxides. For this purpose, it is necessary to add one or more of these elements alone or in total to 0.0001% or more. However, excessive addition degrades moldability, so the upper limit is made 0.5%.

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

Below, the manufacturing method of the high intensity | strength thin steel plate which concerns on this invention is demonstrated.
In producing the high strength thin steel sheet of the present invention, the cast slab is cooled at an average cooling rate of 100 ° C./min or more between the liquidus temperature and the solidus temperature. Here, the average cooling rate refers to the average cooling rate in the middle part of the slab (the position of 1/4 t of the slab of thickness t). In the present invention, casting may be performed by any method as long as the cooling rate during solidification can be higher than 100 ° C./min. For example, in continuous casting, the thickness of the slab may be reduced, in ingot casting, the size of the ingot may be reduced, or a surface layer portion having a high cooling rate may be cut out from a normal slab and used. For example, when the thickness of the continuous cast slab is changed, the thickness of the slab is preferably 100 to 30 mm. This is because when the thickness exceeds 100, the slab cannot be cooled at a sufficiently high cooling rate. When the thickness is less than 30 mm, the casting speed increases, causing fluctuations in the molten metal surface, breakout, etc., and stable slab casting. This is because it becomes difficult.

  In addition, when the average cooling rate between the liquidus temperature and the solidus temperature is less than 100 ° C./min, the molten steel cannot be rapidly solidified, and Mn microsegregation is reduced to 0.10. It cannot be as small as satisfying the relationship of ≧ σ / Mn. Therefore, the said average cooling rate shall be 100 degrees C / min or more. Desirably, cooling is performed at an average of 200 ° C./min or more between the liquidus temperature and the solidus temperature. Thereby, the microsegregation of Mn can be made smaller.

  The slab after cooling can be directly subjected to hot rolling. Alternatively, when it is cooled to less than 1100 ° C., it can be reheated to 1100 ° C. or higher and 1300 ° C. or lower. This is because the deformation resistance in hot rolling is large at a temperature below 1100 ° C., and the generation of scale is large at temperatures exceeding 1300 ° C., and the surface properties of the steel sheet cannot be made favorable.

  Next, the slab is hot-rolled at a finishing temperature of 850 to 970 ° C. If the finishing temperature is less than 850 ° C., (α + γ) two-phase region rolling may occur, and the shape of the plate may be impaired. If it exceeds 970 ° C., the austenite grain size becomes coarse, and the strength and ductility decrease. It is.

  After hot rolling, the steel is first cooled to 700 to 600 ° C. at an average cooling rate of 10 to 100 ° C./sec, then stopped for 1 to 5 seconds, air cooled, and again at an average cooling rate of 10 to 100 ° C./sec. Secondary cooling is performed and winding is performed at a temperature of 300 ° C. or higher and 450 ° C. or lower. When the cooling temperature after hot rolling is 700 ° C. or higher, the formation of ferrite in the subsequent air cooling is slow. On the other hand, when the temperature is lower than 600 ° C., pearlite harmful to moldability is easily generated at an early stage. After cooling, it is retained for 1 to 5 seconds, but if the retention time is less than 1 second, ferrite cannot be sufficiently precipitated, and a desired amount of ferrite can be precipitated by cooling to 5 seconds. is there.

After stopping, it is cooled again at an average cooling rate of 10 to 100 ° C./sec and wound up at a temperature of 300 ° C. or higher and 450 ° C. or lower. This is because harmful pearlite is generated at a cooling rate of less than 10 ° C./sec, and generation of pearlite can be sufficiently suppressed by cooling at 100 ° C./sec. Further, when the coiling temperature exceeds 450 ° C., carbide formation is promoted, and when it is lower than 300 ° C., it becomes difficult to form bainite. Therefore, the coiling temperature is set to 300 ° C. or more and 450 ° C. or less. By performing bainite transformation, the carbon concentration in the austenite can be increased, and stable retained austenite can be left even at room temperature.
As described above, after cooling the slab at high speed, by controlling the temperature and performing hot rolling to take up, the microsegregation of Mn is small, the main phase is ferrite, and 3% or more of retained austenite is contained. A high-strength thin steel sheet having a uniform structure and excellent formability can be produced.

  Moreover, the high-strength thin steel sheet excellent in formability of the present invention can be produced as follows. That is, the casting slab having the above-described chemical component is cooled as it is or 1100 ° C. or more after cooling between the liquidus temperature and the solidus temperature with an average cooling rate of the slab intermediate part being 100 ° C./min or more. Reheat to. The reason for controlling the temperature in cooling the slab is as described above.

  Next, hot rolling is performed at a finishing temperature of 850 to 970 ° C., and then cooling is performed at an average cooling rate of 10 to 100 ° C./sec. A hot-rolled steel sheet. The reasons for limiting the finishing temperature are as described above. This is because when the cooling temperature after hot rolling is higher than 650 ° C., layered pearlite is easily generated. Further, even if the cooling rate is less than 10 ° C./sec, layered pearlite is likely to be generated, and if it exceeds 100 ° C./sec, it is difficult to control the coiling temperature. The reason why the coiling temperature is set to 650 ° C. or lower is that, at a temperature higher than this, lamellar pearlite is easily generated and it is difficult to obtain a uniform hot rolled structure.

The hot-rolled steel sheet produced as described above is cold-rolled at a reduction rate of 40% or more after pickling, and the maximum temperature is 0.1 × (Ac ₃ -Ac ₁ ) + Ac ₁ or more, Ac ₃ + 50 ° C. or less. After being annealed at a temperature of, it is cooled to a temperature range of 350 to 500 ° C. at an average cooling rate of 0.1 to 200 ° C./sec, and subsequently held in the same temperature range for 10 to 1000 seconds, thereby being excellent in moldability. High strength thin steel sheet can be manufactured.

In the production of a cold-rolled steel sheet, if the rolling reduction is less than 40%, crystal grains after annealing cannot be made fine, so the rolling reduction is set to 40% or more.
Further, the maximum temperature of the _{annealing, 0.1 × (Ac 3 -Ac 1} ) + Ac 1 or more, is required to be _Ac 3 + 50 ° C. or less. When the maximum temperature is less than 0.1 × (Ac ₃ −Ac ₁ ) + Ac ₁ (° C.), the amount of austenite obtained at the annealing temperature is small, so that a desired amount of retained austenite may be left in the steel sheet. Can not. Further, increasing the annealing temperature promotes the formation of grain boundary oxide layers and the coarsening of crystal grains, and increases the manufacturing cost, so that the upper limit of the annealing temperature is set to Ac ₃ + 50 ° C. or lower.

  Cooling after annealing is important for promoting the transformation from the austenite phase to the ferrite phase and concentrating C in the untransformed austenite phase to stabilize the austenite. When this cooling rate is less than 0.1 ° C./sec, pearlite is generated, so the lower limit of this cooling rate is 0.1 ° C./sec. On the other hand, when the cooling rate exceeds 200 ° C./sec, the ferrite transformation cannot sufficiently proceed, so the cooling rate after annealing is 0.1 to 200 ° C./sec.

Cooling temperature shall be 350-500 degreeC. This is because martensite is likely to be generated when the temperature is lower than 350 ° C., and it is difficult to generate bainite when the temperature exceeds 500 ° C.
And a steel plate is hold | maintained for 10 to 1000 seconds in the temperature range. This is because if it is less than 10 seconds, sufficient bainite cannot be generated, and the target amount of bainite can be generated by holding up to 1000 seconds. If it exceeds 1000 seconds, carbides are generated.

  After cooling the slab at a high speed as described above, a hot-rolled steel sheet is produced by controlling the temperature, and by cold-rolling and annealing the hot-rolled steel sheet, Mn microsegregation is small, and ferrite is the main phase. A high-strength thin steel sheet having a uniform structure containing 3% or more of retained austenite and excellent in formability can be produced.

Hereinafter, the present invention will be described in detail based on examples.
Steels having chemical components shown in Table 1 that were melted in a converter or a laboratory were cast. At this time, the cooling rate between the liquidus temperature and the solidus temperature at 1/4 t of the slab was changed as shown in Tables 2 and 3. These slabs were subjected to hot rolling to produce hot rolled steel sheets and cold rolled steel sheets. Table 2 shows the manufacturing conditions and material characteristics of the hot-rolled steel sheet, and Table 3 shows the manufacturing conditions and material characteristics of the cold-rolled steel sheet.

In Tables 2 and 3, the volume fraction of retained austenite and its carbon concentration were experimentally determined by X-ray analysis as described in JP-A-11-193435. That is, the volume fraction Vγ of retained austenite can be calculated from the data obtained using the Mo—Kα ray by the following equation.
Vγ = (2/3) {100 / (0.7 × α (211) / γ (220) +1)} + (1/3) {100 / (0.78 × α (211) / γ (311) +1)}
However, α (211), γ (220), α (211), and γ (311) indicate surface strength.
The carbon concentration Cγ of retained austenite is obtained by calculating the lattice constant (unit: angstrom) from the reflection angles of the (200) plane, (220) plane, and (311) plane of austenite by X-ray analysis using Cu-Kα ray. It can be calculated according to the formula.
Cγ = (lattice constant−3.572) /0.033

In Table 1, the upper steel grade with H is steel used for the production of hot-rolled steel sheets, and the lower steel grade with C is steel used for the production of cold-rolled steel sheets.
First, the test results of hot-rolled steel sheet production will be described with reference to Tables 1 and 2.
Steel types AH to GH are steels whose chemical components are within the scope of the present invention. On the other hand, the steel type CAH has a Mn higher than the range of the present invention. For this reason, martensite was generated in the structure, and the strength was high but the elongation was extremely low as shown in Process No. 22.
Steel grade CBH is higher in Cr, Mo and Mg, and steel grade CCH is higher in Ti and Nb than the scope of the present invention. For this reason, as shown in treatment numbers 23 and 24, cracks frequently occurred during hot rolling.

  For treatment numbers 3, 5, 8, 14, and 18, the steel grade has a chemical component that is within the scope of the present invention, but in the cooling of the slab during casting, between the liquidus temperature and the solidus temperature. The cooling rate is significantly lower than 100 ° C./min. For this reason, the index σ / Mn of Mn microsegregation was larger than 0.1, and a Mn band was formed and the structure became non-uniform, resulting in a hot-rolled steel sheet with low elongation.

Process No. 9 has a low hot rolling finishing temperature, a small average cooling rate until winding, and a winding temperature higher than the range of the present invention. For this reason, residual austenite cannot be left, and the strength of the steel sheet is low and the elongation is small.
In the process number 12, the cooling rate after hot rolling is small, the primary cooling stop temperature is low, and the retention time after cooling is longer than the range of the present invention. For this reason, pearlite was produced | generated and a retained austenite could not be produced | generated, As a result, it became the steel plate inferior to an intensity | strength and an elongation balance.
The thing of the process number 21 has the low heating temperature before hot rolling. Moreover, the average cooling rate until winding is small, and winding temperature is high exceeding the range of this invention. As a result, pearlite was generated in the structure, and sufficient bainite could not be generated, resulting in low strength.

  Compared to the comparative examples as described above, the chemical composition of the test steel is appropriate for the treatment numbers 1, 2, 4, 6, 7, 10, 11, 13, 15, 16, 17, 19, and 20. Since the slab cooling conditions and hot rolling conditions were also within the scope of the present invention, the microphase segregation of Mn was small and the main phase was a ferrite structure, and an appropriate amount of retained austenite could be secured. . As a result, a high-strength thin steel sheet having an excellent balance between strength and ductility could be produced.

Next, the test results of cold-rolled steel sheet production will be described with reference to Tables 1 and 3.
Steel types AC to HC are steels whose chemical components are within the scope of the present invention. On the other hand, the steel type CAC has a Mn higher than the range of the present invention. For this reason, martensite was generated in the structure, and as shown in treatment number 53, the strength was high but the elongation was extremely low.
Further, the steel type CBC is higher in the range of Mo and Cu, the steel type CCC is higher in the range of C, Ti and Nb, and the steel type CDC is higher in the range of B and Mg. For this reason, as shown in treatment numbers 54, 55 and 56, cracks frequently occurred during hot rolling, and cold rolling was impossible.

  For treatment numbers 33, 35, 38, 46 and 51, the steel grade has a chemical component within the scope of the present invention, but in the cooling of the slab during casting, between the liquidus temperature and the solidus temperature. The cooling rate is significantly lower than 100 ° C./min. For this reason, the index σ / Mn of Mn microsegregation was larger than 0.1, and as a result of the formation of the Mn band, the structure became non-uniform, resulting in a cold-rolled steel sheet with poor strength and elongation balance.

In the process No. 39, the heating temperature before hot rolling and the finishing temperature of hot rolling are low, the average cooling rate until winding is small, and the rolling reduction of cold rolling is low. For this reason, a retained austenite with a high carbon concentration cannot be generated, and the crystal grains are coarse, so that the elongation of the steel sheet is low.
The thing of the process number 40 has the cooling rate after annealing smaller than the range of this invention. For this reason, pearlite was generated during cooling, and the retained austenite could not be left, resulting in a steel sheet having low strength.

The thing of the process number 41 has the low maximum temperature of annealing, and the cooling stop temperature of annealing is high. For this reason, residual austenite cannot be left, and therefore the strength and elongation of the steel sheet are low.
In the case of treatment number 49, the hot rolling finishing temperature is low and the winding temperature is higher than the range of the present invention. Moreover, the maximum temperature of annealing is low and the holding time after a cooling stop is long. For this reason, the retained austenite cannot be left, and the steel sheet has low strength and small elongation.

  Compared to the comparative examples as described above, the treatment numbers 31, 32, 34, 36, 37, 42 to 45, 47, 48, 50, and 52 are suitable for the chemical composition of the test steel, and the slab Since the cooling conditions, hot rolling conditions, cold rolling and annealing conditions of the present invention were within the scope of the present invention, the micro segregation of Mn was small and the structure containing 3% or more of retained austenite in the ferrite-bainite structure was uniform and formability It was possible to produce a high-strength thin steel sheet with excellent resistance.

It is a conceptual diagram which shows one manufacturing method of the high strength thin steel plate which concerns on this invention.

Claims (7)

In mass%
C: 0.05-0.25%, Si: 2.0% or less, Mn: 0.8-3%, P: 0.0010-0.1%, S: 0.0010-0.05%, N: 0.0010 to 0.010%, Al: 0.01 to 2.0%, and having a steel composition composed of the balance iron and inevitable impurities,
Mn microsegregation in the range of 1 / 8t to 3 / 8t of the plate thickness t is in the range satisfying the formula (1),
A high-strength thin steel sheet with excellent formability characterized by containing 3% or more of retained austenite having an average carbon content of 0.9% or more in the structure.
0.10 ≧ σ / Mn (1)
Here, Mn is an addition amount, and σ is a standard deviation in Mn microsegregation measurement. Further during the steel composition
Cr: 0.01-5%, Mo: 0.01-5%, Ni: 0.01-5%, Cu: 0.01-5%, Co: 0.01-5%, W: 0.01 The high-strength thin steel sheet having excellent formability according to claim 1, containing ˜5% of one kind or two or more kinds. Further during the steel composition
One or two or more of Ti, Nb, Zr, Hf, Ta, and V are contained alone or in total in an amount of 0.001 to 1%, and the moldability according to claim 1 or 2 is excellent. High strength thin steel sheet. Further during the steel composition
The high strength thin steel sheet having excellent formability according to any one of claims 1 to 3, wherein B is contained in an amount of 0.0001 to 0.0050%. Further during the steel composition
The high-strength thin steel sheet having excellent formability according to any one of claims 1 to 4, comprising 0.0001 to 0.5% of one or more of Mg, Ca, Y, and REM. . A method for producing a high-strength thin steel sheet, wherein the high-strength thin steel sheet according to claim 1 is produced from a slab,
The slab in the middle of cooling after casting is cooled as it is or 1100 ° C after cooling between the liquidus temperature and the solidus temperature at an average cooling rate at 1/4 t of the slab thickness t of 100 ° C / min or more. Reheat to above,
Subsequently, hot rolling was performed at a finishing temperature of 850 to 970 ° C., and thereafter cooling was performed at an average cooling rate of 10 to 100 ° C./sec to a temperature range of 700 to 600 ° C., and then retained in the same temperature range for 1 to 5 seconds. After that, it is cooled again at an average cooling rate of 10 to 100 ° C./sec and wound at a temperature of 300 ° C. or higher and 450 ° C. or lower to produce a hot-rolled steel plate. Method. A method for producing a high-strength thin steel sheet, wherein the high-strength thin steel sheet according to claim 1 is produced from a slab,
The slab that is in the process of being cooled after casting is cooled as it is or 1100 after cooling the temperature range from the liquidus temperature to the solidus temperature at an average cooling rate at 1/4 t of the slab thickness t of 100 ° C./min or more. Reheat to over ℃,
Hot-rolling is performed at a finishing temperature of 850 to 970 ° C., and then cooled to a temperature range of 650 ° C. or lower at an average cooling rate of 10 to 100 ° C./sec. None,
The hot-rolled steel sheet is subjected to cold rolling with a reduction rate of 40% or more after pickling,
After annealing at a maximum temperature of 0.1 × (Ac ₃ −Ac ₁ ) + Ac ₁ or more and Ac ₃ + 50 ° C. or less, 350 ° C. or more and 500 ° C. or less at an average cooling rate of 0.1 to 200 ° C./sec. A method for producing a high-strength thin steel sheet having excellent formability, wherein the steel sheet is cooled to a temperature range and subsequently kept in the same temperature range for 10 seconds to 1000 seconds to form a cold-rolled steel sheet.

Images