JP2009108343A - High strength steel plate and manufacturing method thereof - Google Patents

Abstract

The present invention provides a high-strength steel sheet suitable for use as a material for various reinforcing members for automobile bodies, automobile seat members, and the like because the tensile strength is 780 MPa or more and excellent bendability.
SOLUTION: C: 0.05 to 0.17%, Mn: 2.0 to 3.0%, B: 0.0005% to 0.01%, P: 0.03% or less, S: 0.0. 01% or less, N: 0.01% or less, Si: 0.005-0.5%, sol. Al: 0.01% to 0.1% of both are contained in a total of 0.5% or less, and one or two of Ti and Nb are added to 0.05% ≦ [Ti] + [Nb] / 2 ≦ 0.2% of content is included, the chemical composition is composed of the remaining Fe and impurities, does not contain unrecrystallized ferrite, and has a steel structure in which the average grain size of ferrite and bainite is 3.5 μm or less. The minimum value of B intensity by glow discharge emission spectrometry using argon sputtering at a depth of 2 to 10 μm from the surface is glow discharge emission spectrometry using argon sputtering at a depth of 80 μm from the surface. Is a high-strength steel sheet having a B strength of 60% or less and a tensile strength of 780 MPa or more. It manufactures by performing a continuous annealing on the conditions which reduce B on the steel plate surface after cold rolling.
[Selection figure] None

Description

  The present invention relates to a high-strength steel sheet suitable for use as a raw material for various reinforcing members for automobile bodies and seat members for automobiles, and a method for producing the same.

  As a part of weight reduction for improving the safety and fuel consumption of automobiles, high-strength steel sheets with excellent workability have been increasingly attracting attention as materials for automobile bodies and parts. In recent years, the demand for higher strength has become more severe. For example, high strength steel sheets having a tensile strength of 780 MPa or more have been used. An extremely high level of bendability is often required for such high-strength steel sheets. In particular, recently, high-strength steel plates are being used even for automobile parts having bending portions with a small bending radius, such as seat rails, and there is a tendency for higher bendability than ever before.

  Regarding improvement of bendability of high-strength steel sheets, as disclosed in Patent Document 1, the hardness of the low-temperature transformation generation phase is reduced to reduce the hardness difference from the soft ferrite phase, and a uniform tempered structure is obtained. It is considered effective.

  Furthermore, Patent Document 2 states that by using a martensite single-phase structure that is the ultimate uniform structure, stretch flangeability that requires local deformability as well as bendability can be improved.

On the other hand, Patent Document 3 states that both a bendability and a high strength can be achieved by forming a ductile structure mainly composed of ferrite only on the surface.
Japanese Patent Laid-Open No. 62-13533 JP 2002-161336 A Japanese Patent Laid-Open No. 10-130782

  In a high-strength steel sheet containing a large amount of alloying elements, the chemical composition partially varies due to solidification segregation, and component segregation is likely to occur. Therefore, in the invention disclosed by Patent Document 1 containing a large amount of Ti, Nb, etc., it is extremely difficult to precisely control the hardness of the ferrite phase and the low temperature transformation phase for the entire steel sheet. In addition, not only the bendability itself deteriorates due to the unevenness of the structure caused by the segregation of components, but also when the parts are subjected to strong processing such as bending, the irregularities on the surface of the bent portion become prominent. Since this tends to be the starting point of cracking, the collision characteristics of this part may deteriorate.

  In the invention disclosed in Patent Document 2, the shape of the steel sheet becomes poor due to rapid cooling and transformation for making the steel structure into a martensite single phase, and the flatness of the steel sheet is impaired. Moreover, in this invention, since a welding heat affected zone is easy to soften, it is difficult to use it for an automobile part to be welded. For this reason, it is not possible to provide a high-strength steel sheet excellent in bendability suitable for automobile parts only by controlling the structure of the base material.

  Since the distortion of the bent portion is concentrated on the surface, the bendability is governed by the ductility near the surface. For this reason, making the ferrite structure only ductile as in the invention disclosed in Patent Document 3 is suitable for automobile parts and is effective in improving the bendability of a high-strength steel sheet. However, in the present invention, a large amount of scale is generated by decarburization of the surface of the steel plate necessary for making the surface into a ferrite structure, so that not only the surface properties of the steel plate are impaired, but also the generated scale is removed. Therefore, the plate thickness is lowered and the accuracy of the plate thickness cannot be ensured. Therefore, in addition to the control of the base metal structure, the ductility of the surface of the steel sheet must be improved by means other than decarburization.

  An object of the present invention is a high-strength steel plate suitable for use as a material for various reinforcing members of automobile bodies, automobile seat members, and the like, and a method for producing the same, since the tensile strength is 780 MPa or more and has excellent bendability. Is to provide. In the present specification, “excellent bendability” means that the minimum bending radius of the 180 ° bending test is 1.5 t or less. Therefore, unless otherwise specified, the bendability in this specification is evaluated based on such physical properties.

  As a result of intensive studies on the composition, structure, and production conditions in order to provide a high-strength steel sheet having the above-mentioned characteristics, the present inventors have made non-recrystallization by making the composition and production conditions appropriate. Argon sputtering at a depth position of 2 to 10 μm from the surface by reducing the microstructure of the ferrite and bainite average particle size to 3.5 μm or less and reducing the surface B by a continuous annealing process. By setting the minimum value of B intensity by glow discharge emission spectrometry using NO to 60% or less of the B intensity by glow discharge emission spectrometry using argon sputtering at a depth of 80 μm from the surface, the tensile strength is reduced. Knowing that a high-strength steel sheet having a bendability of 780 MPa or more and excellent in bendability can be produced, the present invention has been completed.

In the present invention, C: 0.05% or more and 0.17% or less (in the present specification, “%” relating to composition means “% by mass” unless otherwise specified), Mn: 2.0% More than 3.0% or less, B: 0.0005% or more and 0.01% or less, P: 0.03% or less, S: 0.01% or less, N: 0.01% or less, and Si: 0.005% to 0.5%, sol. Al: 0.01% or more and 0.1% or less in total, 0.5% or less in total, and further containing one or two of Ti and Nb in a range satisfying the following formula (1), Furthermore, as an optional additive element, Cr: 1.0% or less, Mo: 0.5% or less, Ni: 0.5% or less, Cu: 0.5% or less, and V: 0.2% or less 1 selected from the group consisting of one or more selected, Ca: 0.003% or less, Mg: 0.003% or less, REM: 0.01% or less, and Zr: 0.01% or less Containing a seed or two or more, having a chemical composition consisting of the balance Fe and impurities, not containing unrecrystallized ferrite, and having a steel structure in which the average grain size of ferrite and bainite is 3.5 μm or less, Glow discharge generation using argon sputtering at a depth of 2 to 10 μm from the surface The minimum value of B strength by optical spectroscopic analysis (hereinafter abbreviated as “GDS analysis”) is 60% or less of B strength by GDS analysis at a depth of 80 μm from the surface, and the tensile strength is 780 MPa or more. It is a featured high strength steel plate.
0.05% ≦ [Ti] + [Nb] /2≦0.2% (1)
However, in the formula (1), [Ti] and [Nb] indicate the contents (%) of Ti and Nb, respectively.

  Here, "GDS analysis" is a glow discharge tube having a special structure, causing discharge using the sample as a cathode, and existing on this surface while scraping the surface of the sample using a sputtering phenomenon caused by argon ions. Means glow discharge optical emission spectrometry for measuring elements. As analysis conditions, GDS (manufactured by Horiba, JY-5000RF) is used, and the anode diameter is 4 mmφ, the RF output is 35 W, and the argon pressure is 600 Pa. In addition, the depth is calculated by setting the sputtering speed to 0.1 μm / s.

From another viewpoint, the present invention is a method for producing a high-strength steel sheet, comprising the following steps (A) to (C).
(A) The steel material having the chemical composition of the high-strength steel sheet according to the present invention described above is subjected to hot rolling after being set to 1100 ° C. or higher and 1300 ° C. or lower, and this hot rolling is performed in a temperature range of 800 ° C. or higher and 950 ° C. or lower. The hot rolling process which complete | finishes and winds in the temperature range of 450 degreeC or more and 700 degrees C or less to make a hot-rolled steel plate.
(B) A cold rolling step in which the hot rolled steel sheet is pickled and then cold rolled to form a cold rolled steel sheet. And (C) Austenite The cold rolled steel sheet is heated to _{N 2} concentration temperature range of not lower than Ac ₃ transformation point temperature range over a period of 20 seconds 600 ° C. to Ac ₃ transformation point in atmosphere of more than 95 vol% After maintaining in the austenite single-phase structure state for 60 seconds or more in an atmosphere having an N ₂ concentration of 95% by volume or more in an atmosphere having a single-phase structure, the average cooling rate from the Ar ₃ transformation point to 550 ° C. becomes 5 ° C./second or more. Continuous annealing process for cooling under cooling conditions.

  According to the present invention, since the tensile strength is 780 MPa or more and the bendability is excellent, for example, a high-strength steel sheet suitable for use as a raw material for various reinforcing members of automobile bodies, automobile seat members, and the like, and a method for producing the same are provided. As a result, it is possible to promote weight reduction of the vehicle body and its components.

Hereinafter, the best mode for carrying out the high-strength steel sheet and the manufacturing method thereof according to the present invention will be described in detail with reference to the accompanying drawings.
The reasons for limiting the chemical composition of the high-strength steel sheet according to the present invention are as follows.

(C: 0.05% or more and 0.17% or less)
C is an element contributing to strength improvement, and is contained at least 0.05% in order to ensure a tensile strength of 780 MPa or more. However, if it exceeds 0.17%, it becomes a non-uniform structure and the bendability deteriorates. Therefore, the C content is set to 0.05% or more and 0.17% or less. Preferably they are 0.05% or more and 0.10% or less.

(Mn: 2.0% to 3.0%)
Mn is an element that contributes to strength improvement, and is contained at least 2.0% in order to make the steel sheet have a tensile strength of 780 MPa or more. However, if the content exceeds 3.0%, the structure becomes uneven and the bendability deteriorates. Therefore, the Mn content is set to 2.0% to 3.0%.

(B: 0.0005% to 0.01%)
B is one of important elements, and not only makes the structure uniform and fine, but also suppresses the occurrence of microcracks by the effect of increasing the strength of grain boundaries and heterogeneous interfaces, and improves bendability. For this purpose, it contains at least 0.0005%. However, if the content exceeds 0.01%, the above-described effects are saturated and not only economically wasteful, but borides are formed at the grain boundaries and the bendability deteriorates. Therefore, the B content is set to 0.0005% or more and 0.01% or less.

  In addition, the boride which exists in the surface vicinity of a steel plate becomes a crack generation source by concentration of distortion, and degrades the bendability of a steel plate. For this reason, in the present invention, in order to reduce the amount of boride existing in the vicinity of the surface and ensure bendability, the B concentration in the vicinity of the surface is reduced as described later.

(P: 0.03% or less)
P is an impurity in the present invention, and if P is excessively contained, it becomes a non-uniform structure and bendability deteriorates. Therefore, the P content is 0.03% or less. Preferably it is 0.015% or less.

(S: 0.01% or less)
S exists as a sulfide in steel, and this sulfide becomes a stress concentration source and the bendability deteriorates. For this reason, it is desirable that the S content is as low as possible. However, if the S content is 0.01% or less, the bendability is not adversely affected. Therefore, the S content is set to 0.01% or less. Preferably it is 0.005% or less.

(N: 0.01% or less)
N is an impurity in the present invention. When N is excessively contained, coarse nitrides are precipitated and workability is deteriorated. For this reason, it is desirable that the N content is as low as possible. However, if the N content is 0.01% or less, the workability is not adversely affected. Therefore, the N content is 0.01% or less. Preferably it is 0.005% or less.

(Si: 0.005% to 0.5%, sol.Al: 0.01% to 0.1%, 0.5% or less in total)
Since Si is an element that contributes to the improvement of strength without significantly degrading the bendability, it is contained in an amount of 0.005% or more. However, when it contains exceeding 0.5%, chemical conversion property will deteriorate. For this reason, Si content shall be 0.005% or more and 0.5% or less.

  Al is an element added for deoxidation of steel, and effectively acts to improve the cleanliness of the steel. In order to remove silicate inclusions and improve bendability, sol. Al content of 0.01% or more. However, if the content exceeds 0.1%, oxide inclusions increase and the surface properties deteriorate. For this reason, sol. Al content shall be 0.01% or more and 0.1% or less. Preferably they are 0.01% or more and 0.06% or less.

Furthermore, Si and sol. When both Al content exceeds 0.5% in total, the Ac ₃ transformation point is increased, Mn oxide is generated during continuous annealing in the austenite single phase region, and the surface properties of the steel sheet deteriorate. For this reason, Si content and sol. The total Al content is 0.5% or less.

(Ti, Nb: range satisfying 0.05% ≦ [Ti] + [Nb] /2≦0.2%)
Ti and Nb are both important elements, and improve the strength without significantly degrading the bendability by precipitation strengthening and crystal grain refinement. Therefore, it contains one or both of Ti and Nb.

  In order to improve bendability by refining crystal grains, Ti and / or Nb is contained so that the value of ([Ti] + [Nb] / 2) is 0.05% or more. Here, [Ti] and [Nb] indicate the contents (%) of Ti and Nb, respectively. However, when the value of ([Ti] + [Nb] / 2) exceeds 0.2%, precipitates in the steel become coarse and the strength is lowered. Therefore, the content of Ti and / or Nb is set such that the value of ([Ti] + [Nb] / 2) is 0.05% or more and 0.2% or less.

As will be described later, by setting the value of ([Ti] + [Nb] / 2) to 0.05% or more, performing heating above the Ac ₃ transformation point temperature, and annealing holding for 60 seconds or more in the austenite phase A high-strength steel sheet having excellent bendability can be obtained.

(Cr: 1.0% or less, Mo: 0.5% or less, Ni: 0.5% or less, Cu: 0.5% or less, V: 0.2% or less)
Cr, Mo, Ni, Cu and V are all optional additional elements contained as necessary. In order to ensure a tensile strength of 780 MPa or more without significantly degrading the bendability, it is preferable to contain one or more of Cr, Mo, Ni, Cu and V. However, Cr content is over 1.0%, Mo content is over 0.5%, Ni content is over 0.5%, Cu content is over 0.5%, V content is 0.2% If it is super, the above-described effect is saturated and it is economically useless. For this reason, Cr: 1.0% or less, Mo: 0.5% or less, Ni: 0.5% or less, Cu: 0.5% or less, V: 0.2% or less.

  In order to reliably obtain the above-described effects, Cr: 0.05% or more, Mo: 0.05% or more, Ni: 0.03% or more, Cu: 0.03% or more, V: 0.005% or more It is preferable to do.

(Ca: 0.003% or less, Mg: 0.003% or less, REM: 0.01% or less, Zr: 0.01% or less)
Ca, Mg, REM and Zr are also optional additional elements contained as necessary, and one or two of Ca, Mg, REM and Zr are used to improve bendability by controlling the form of inclusions. It is preferable to contain the above. However, if the Ca content exceeds 0.003%, the Mg content exceeds 0.003%, the REM content exceeds 0.01%, and the Zr content exceeds 0.01%, the above-described effects are saturated. It is economically wasteful. For this reason, Ca: 0.003% or less, Mg: 0.003% or less, REM: 0.01% or less, Zr: 0.01% or less.

In order to reliably obtain the above-described effects, it is preferable to set Ca: 0.0005% or more, Mg: 0.0005% or more, REM: 0.0005% or more, and Zr: 0.0002% or more.
The balance other than the components described above is Fe and impurities. As impurities, in addition to the above-described P, S, and N, for example, O: 0.006% or less can be allowed.

  The high-strength steel sheet according to the present invention does not contain non-recrystallized ferrite, has a steel structure in which the average grain size of ferrite and bainite is 3.5 μm or less, and has an argon sputter at a depth of 2 to 10 μm from the surface. Distribution of B concentration in the plate thickness direction in which the minimum value of B intensity by glow discharge optical emission spectrometry using Gd is 60% or less of G intensity by GDS analysis at a depth of 80 μm from the surface (in this specification, (Referred to as “cross-section B concentration distribution”). These will be described.

(Section B concentration distribution)
When the B concentration near the surface, that is, at a depth of 2 to 10 μm from the surface (minimum value of B intensity by GDS analysis) increases, not only the bendability deteriorates due to the formation of boride, but also bending after bending. The workability at the time of applying is significantly deteriorated. For this reason, the high-strength steel sheet according to the present invention has a minimum value of B intensity by GDS analysis at a depth of 2 to 10 μm from the surface, which is 60% or less of the B intensity by GDS analysis at a depth of 80 μm from the surface. The cross-section B concentration distribution is as follows.

Moreover, since the high strength steel plate according to the present invention has such a cross-sectional B concentration distribution, only the surface is softened and thus has excellent bendability.
In order to have such a cross-sectional B concentration distribution, as will be described later, B is reduced from the vicinity of the surface of the steel sheet during continuous annealing.

(Non-recrystallized ferrite)
Non-recrystallized ferrite is ferrite in which processing strain due to cold rolling remains. If non-recrystallized ferrite is present, non-uniform deformation is promoted and bendability deteriorates. Therefore, the structure does not contain unrecrystallized ferrite. In order not to include unrecrystallized ferrite, it is necessary to heat to the Ac ₃ transformation point or higher as described later. The presence or absence of non-recrystallized ferrite can be measured by SEM. In the present invention, an elongated ferrite having an aspect ratio of 5 or more is defined as non-recrystallized ferrite.

(Average particle size of ferrite and bainite)
If the average particle size of ferrite and bainite becomes too large, microcracks are likely to occur at the interface between ferrite or bainite that is a soft phase and martensite that is a hard phase, and crack propagation is not suppressed. Deteriorates. For this reason, the average particle diameter of ferrite and bainite is set to 3.5 μm or less. This average particle diameter means the average value of the particle diameters of all ferrite crystal grains and all bainite crystal grains at the measurement site. The grain size of bainite means a packet diameter composed of several laths. This average particle diameter is determined by the cross line segment method defined in JIS G 0552.

  The high-strength steel sheet according to the present invention is a cold-rolled annealed steel sheet, and is subjected to recrystallization annealing to make the austenite single phase completely after the crystal grains are stretched in the rolling direction by cold rolling. There is no significant difference between the size of the crystal grains in the rolling direction and the size of the crystal grains in the direction perpendicular to the rolling direction. However, as a precaution, in the present invention, the average particle diameter of ferrite and bainite is an average value of a value measured in the rolling direction and a value measured in the direction perpendicular to the rolling.

  The area ratio of ferrite and bainite in the metal structure of the high-strength steel sheet according to the present invention is not particularly specified, but when the area ratio of ferrite and bainite is less than 70% in total, the elongation of uniaxial tensile deformation is significantly reduced, Molding other than bending may be difficult. On the other hand, if the area ratios of ferrite and bainite are 70% or more in total, it is possible to ensure an elongation of 6% or more in the tensile test, to enable stretch forming, and to secure a wide range of uses of the steel sheet. For this reason, it is preferable that the area ratios of ferrite and bainite are 70% or more in total.

The high-strength steel sheet according to the present invention has the above chemical composition, steel structure, and cross-section B concentration distribution, and has a tensile strength of 780 MPa or more.
Next, the manufacturing method of the high strength steel plate which concerns on this invention is demonstrated in order of a hot rolling process (process A), a cold rolling process (process B), and a continuous annealing process (process C).

(Process A)
The molten steel having the above-described chemical composition is melted by a known melting method using a converter, an electric furnace or the like, and is made into a steel material such as a slab by a continuous casting method. Instead of the continuous casting method, an ingot-making method or a thin slab casting method may be used.

  In order to use this steel material for hot rolling, it is set to 1100 ° C. or higher and 1300 ° C. or lower. Remelting TiC and NbC during heating of the steel material, making TiC and NbC fine during annealing after cold rolling, and reducing the bendability by making the ferrite and bainite of the steel structure fine as described above. In order to prevent this, the steel material is heated to 1100 ° C. or higher. However, heating the steel material above 1300 ° C. not only saturates these effects, but also increases scale loss. For this reason, the heating temperature of a steel raw material shall be 1100 degreeC or more and 1300 degrees C or less.

  Next, this steel material is hot-rolled to obtain a hot-rolled steel sheet. When the steel material is a continuous cast slab, hot rolling is performed by direct feed rolling in which the steel material is charged in a heating furnace without being cooled to room temperature and then heated and rolled immediately after performing slight heat retention. Either rolling or rolling in which the steel material is once cooled to room temperature and then heated may be used.

  In this invention, the completion | finish temperature of this hot rolling shall be 800 degreeC or more and 950 degrees C or less. If the end temperature of hot rolling is less than 800 ° C., deformation resistance during hot rolling increases and productivity decreases. On the other hand, when the end temperature of hot rolling exceeds 950 ° C., most of Ti or Nb in the steel precipitates as carbides in the hot-rolled steel plate during the subsequent cooling, and then cold rolling can be performed. become unable.

  Furthermore, in this invention, the coiling temperature of the steel plate after hot rolling shall be 450 degreeC or more and 700 degrees C or less. When the coiling temperature is less than 450 ° C., hard bainite and martensite are generated, and subsequent cold rolling cannot be performed. On the other hand, when the coiling temperature exceeds 700 ° C., scale generation is promoted and the unevenness of the steel sheet becomes remarkable, and the surface properties of the steel sheet after cold rolling and continuous annealing deteriorate.

  The hot rolling conditions other than those described above may be well-known and commonly used conditions, and such conditions are obvious to those skilled in the art, and thus further description of the hot rolling conditions is omitted.

(Process B)
The hot-rolled steel sheet manufactured by the process A is pickled by a normal method and then cold-rolled to obtain a cold-rolled steel sheet. The rolling reduction in this cold rolling is preferably 30% or more in order to make the ferrite and bainite of the steel structure fine as described above during annealing after cold rolling.

  Conditions for cold rolling other than those described above may be well-known and conventional conditions, and such conditions are obvious to those skilled in the art, and thus further description regarding the conditions for cold rolling is omitted.

(Process C)
The cold-rolled steel sheet manufactured by the process B is continuously annealed under heating conditions, soaking conditions and cooling conditions described below.

[Heating conditions]
The cold-rolled steel sheet, over N ₂ concentration is 600 ° C. or higher Ac ₃ temperature range below the transformation point or more for 20 seconds in an atmosphere of more than 95% by volume, heated to a temperature above Ac ₃ transformation point to a single phase of austenite Heat under conditions.

By heating the temperature range of 600 ° C. or more and the Ac ₃ transformation point or less in an atmosphere having an N ₂ concentration of 95% by volume or more over 20 seconds, the B concentration near the surface of the steel sheet is lowered to reduce the B concentration distribution in the cross section described above. The ductility near the surface of the steel sheet is improved, and the workability of the steel sheet is improved.

  In the present invention, the cold rolled steel sheet containing Ti and / or Nb is once made into an austenite single phase structure by heating during continuous annealing, whereby the structure is made uniform and fine. Moreover, by containing Ti and / or Nb and having a uniform and fine structure, bendability is improved while having high strength.

If the heating temperature is lower than the Ac ₃ transformation point, unrecrystallized ferrite remains and a band-like structure is formed, so that the bendability is significantly deteriorated. Therefore, the heating temperature of the cold-rolled steel sheet to Ac ₃ transformation point or more. Since the continuous annealing furnace and the heating temperature exceeds 900 ° C. is easily damaged, the heating temperature is preferably set to Ac ₃ transformation point temperature or higher 900 ° C. or less.

[Soaking conditions]
The cold-rolled steel sheet heated to the Ac ₃ transformation point or higher under the heating conditions described above is held in an austenite single-phase structure state for at least 60 seconds in an atmosphere having an N ₂ concentration of 95% by volume or higher. When the holding time is less than 60 seconds, not only the above-described cross-section B concentration distribution cannot be obtained, but also B and C are unevenly distributed under the influence of Mn segregation, and the structure of the steel sheet after annealing is It becomes non-uniform. By setting the holding time to 60 seconds or more, not only can the above-mentioned cross-sectional B concentration distribution be obtained, but also B and C that are non-uniformly distributed due to the influence of Mn segregation will be uniformly distributed, and annealing is performed. The structure of the later steel sheet is homogenized.

What is necessary is just to hold the state of an austenite single-phase structure, and it is not always necessary to hold it above the Ac ₃ transformation point temperature. For example, after heating above Ac ₃ transformation temperature, it is also possible to slow cooling to a temperature range below Ac ₃ transformation temperature at Ar ₃ transformation point temperature (= ferrite precipitation start temperature) or more, in this case The holding time includes a time during which the temperature stays in a temperature range not lower than the Ar ₃ transformation point temperature and lower than the Ac ₃ transformation point temperature. That is, the sum of the stay time in this temperature range and the stay time in the temperature range equal to or higher than the Ac ₃ transformation point temperature is the retention time in the state of the austenite single phase structure.

  Usually, when the steel sheet is held for a long time in the austenite single phase state, grain growth becomes remarkable and a fine grain structure cannot be obtained. However, in the present invention, since the contents of Ti and / or Nb and B are made relatively high, it is possible to suppress grain growth even when held in an austenite single-phase state for a long time, whereby ferrite and bainite. A fine grain structure having an average particle diameter of 3.5 μm or less can be obtained.

[Cooling conditions]
The cold-rolled steel sheet is heated and maintained under the above-described conditions, then rapidly cooled from the austenite single phase state, and cooled under the cooling condition that the average cooling rate is 5 ° C./second or more from the Ar ₃ transformation point to 550 ° C. To do. If the average cooling rate from the Ar ₃ transformation point to 550 ° C. is less than 5 ° C./second, it is difficult to ensure a tensile strength of 780 MPa or more.

In order to improve the flatness of the steel sheet after annealing, it is preferable to set the cooling stop temperature to 200 ° C. or higher and hold it in a temperature range of 200 ° C. or higher and 500 ° C. or lower for 50 seconds or longer.
Furthermore, it is preferable to perform temper rolling with an elongation of 0.1% to 1%. Yield point elongation can be suppressed by this temper rolling.

  Moreover, when corrosion resistance is required for the high-strength steel sheet according to the present invention, the surface may be subjected to molten metal plating or electroplating. The plating type is not particularly defined, but is usually zinc or a zinc alloy. However, other platings such as aluminum and aluminum alloys are possible.

  Thus, according to the present invention, the average particle diameter of ferrite and bainite has a uniform steel structure and a cross-section B concentration distribution of 3.5 μm or less, a tensile strength of 780 MPa or more and excellent bendability. Therefore, for example, a high-strength steel plate suitable for use as a material for various reinforcing members of an automobile body or a seat member of an automobile can be manufactured.

An ingot obtained by melting a test steel having the chemical composition shown in Table 1 was forged so as to be a slab having a thickness of 20 mm.
The slab is heated to the slab heating temperature shown in Table 2 and then hot-rolled, and the hot rolling is finished at the finish rolling temperature shown in Table 2, followed by winding at a cooling rate of about 20 ° C./second. The steel sheet is cooled to a temperature of 650 ° C., simulated for winding and held at this temperature for 30 minutes, and then cooled to room temperature at a cooling rate of 20 ° C./hour to obtain a hot-rolled steel sheet having a thickness of 2.4 mm. Obtained.

After this hot-rolled steel sheet was pickled, it was cold-rolled to a sheet thickness of 1.2 mm at a reduction rate of 50% to obtain a cold-rolled steel sheet.
The cold-rolled steel sheet was subjected to a heat treatment simulating continuous annealing (heating atmosphere with N ₂ concentration of 97%). First, it is heated to 600 ° C. at a temperature rising rate of 10 ° C./second, and is heated from 600 ° C. to the Ac ₃ transformation point for the heating time shown in Table 2 (if the temperature is below the Ac ₃ transformation point), After maintaining the annealing time shown in Table 2 at the annealing temperature shown in No. 2, it was cooled to 700 ° C. at the cooling rate after annealing shown in Table 2.

  Rapid cooling was started from 700 ° C., and cooling was performed at the cooling rate and cooling stop temperature shown in Table 2. The rapidly cooled steel sheet was kept at the rapid cooling stop temperature for 300 seconds, and then cooled to room temperature at a rate of 10 ° C./second.

After cooling, temper rolling with an elongation of 0.2% was performed. 1-44 were obtained. The Ar ₃ point (ferrite precipitation start temperature) in Table 2 was determined by thermal expansion analysis when this heat treatment was performed.

The obtained test steel plate No. 1-4, the Ac ₃ transformation point (described in Table 1) and the retention time of the austenite single phase structure state (“γ single phase structure time” in Table 2) were measured, and GDS analysis, structure observation, tensile test and A bending test (the above results are shown in Table 3) was performed. The test methods are listed below.

(Measurement of Ac ₃ transformation point temperature)
A specimen was collected from each unrolled cold-rolled steel sheet, and the temperature of the Ac ₃ transformation point was determined by analyzing the change in expansion coefficient when heated from room temperature to 1000 ° C. at 10 ° C./s.
(Measurement of retention time of austenite single phase structure)
A specimen is taken from each unheat-treated cold-rolled steel sheet, and the change in the expansion coefficient when the heat treatment is performed under the same annealing conditions and cooling conditions as shown in Table 2 makes it possible to maintain the austenite single-phase structure state. Was measured.

(GDS analysis)
Based on the analysis conditions described above, Fe and B were measured at a sputtering rate of 0.1 μm / second, and sampled at intervals of 0.02 seconds.

(Tissue observation)
Test steel plate No. 1 to 44, a test piece having a plate thickness cross section in the rolling direction and a test piece having a plate thickness cross section in the direction perpendicular to the rolling direction are prepared. Each structure is photographed with an optical microscope or an electron microscope, and the ferrite phase ( The fraction of each phase of the bainite phase (shown as B in Table 3) and the average crystal grain size combining these phases were measured. The particle size was measured in accordance with the JIS G 0552 cross-line method for the total thickness of the sheet thickness section in the rolling direction and the sheet thickness section in the direction perpendicular to the rolling direction. Expressed by value. The presence or absence of non-recrystallized ferrite, a field of view of 0.04 mm ² when observed with an electron microscope, the aspect ratio and without the case where five or more elongated ferrite is not confirmed, there a case where elongated ferrite is confirmed It was.

(Tensile test)
Test steel plate No. JIS No. 5 tensile test pieces having the longitudinal direction of 1 to 44 as the longitudinal direction were taken, and the tensile properties (tensile strength TS, yield strength YS, elongation El) were investigated. Moreover, the yield ratio (YR) which is the value of YS / TS was calculated.

(Bending test)
Test steel plate No. A bending test piece (width 40 mm, length 160 mm, plate thickness 1.2 mm) having a longitudinal direction perpendicular to the rolling direction is taken from 1 to 44, and subjected to a 180 ° bending test with a 4.8 mm steel plate sandwiched between cracks. The presence or absence was confirmed visually.

  For a test piece without cracks, the thickness of the sandwiched steel sheet is 4.2 mm thinner than the previous one by 4.2 mm, and the presence of cracks is similarly confirmed by performing a 180 ° bending test. The thickness of the sandwiched steel sheet was decreased every 0.6 mm and changed to 3.6 mm, 3.0 mm, 2.4 mm, 1.8 mm, 1.2 mm, and 0.6 mm, and the 180 ° bending test was sequentially performed. . And even if it did the 180 degree bending test which pinched | interposed the 0.6 mm steel plate, when there was no crack, the adhesion bending which did not pinch the steel plate was performed.

  Then, by dividing the plate thickness of the steel plate in which no cracks are recognized after the bending test by 1.2 mm, which is the plate thickness of the bending test piece, the minimum bending radius (the limit bending (in Table 3) t) ") was calculated.

  In Tables 2 and 3, the test steel plate No. 1-4, 7, 10, 11, 13, 15, 16, 18-22, 24, 25, 27, 28, 30, 33, 35-37, 39-41, 43 and 44 are all the present invention. This is an example of the present invention that satisfies all the conditions specified in (1).

  These test steel plates of the present invention do not contain unrecrystallized ferrite, have a fine and uniform steel structure in which the average grain size of ferrite and bainite is 3.5 μm or less, and have a depth of 2 to 10 μm from the surface. Since the minimum value of the B intensity by the GDS analysis at the depth position has a cross-section B concentration distribution which is 60% or less of the B intensity by the GDS analysis at a depth position of 80 μm from the surface, the ductility near the surface is high. For this reason, the test steel sheets of these examples of the present invention also have excellent bending characteristics such as a minimum bending radius of 1.5 t or less, despite having a high tensile strength of 780 MPa or more.

  The test steel plate No. No. 12 has a total area ratio of 70% or more for ferrite and bainite, so that the bendability is good and a high strength of 780 MPa or more can be obtained, but the elongation is small and the ductility is poor.

  In contrast, the test steel plate No. 5, 6, 8, 9, 12, 14, 17, 23, 26, 29, 31, 32, 34, 38, and 42 are comparative examples that do not satisfy at least one of the conditions defined in the present invention.

Test steel plate No. In No. 5, since the B content is below the lower limit of the range defined in the present invention, the effect of suppressing microcracks is small and the bendability is poor.
Test steel plate No. No. 6 has a low tensile strength because the Mn content is below the lower limit of the range defined in the present invention.

Test steel plate No. No. 8, since the heating time at the 600 to Ac ₃ transformation point in continuous annealing is less than the lower limit of the range defined in the present invention, B near the surface cannot be sufficiently reduced, and a desired cross-sectional B concentration distribution can be obtained. Since it was not obtained, the bending characteristics were unsatisfactory.

  Test steel plate No. 9 is because the heating temperature of the slab in the hot rolling is below the lower limit of the range specified in the present invention, TiC and NbC in the slab cannot be sufficiently re-dissolved, and the structure after annealing becomes coarse, The bendability is poor.

Test steel plate No. In No. 12, since the annealing temperature in continuous annealing is below the lower limit of the range defined in the present invention, unrecrystallized ferrite remains, a band-like structure is formed, and the bendability is poor.
Test steel plate No. No. 14 has a low tensile strength because the C content is below the lower limit of the range defined in the present invention.

Test steel plate No. No. 17, since the heating time at the 600 to Ac ₃ transformation point in continuous annealing is less than the lower limit of the range defined in the present invention, B near the surface cannot be sufficiently reduced, and a desired cross-sectional B concentration distribution is obtained. Since it was not obtained, the bending characteristics were unsatisfactory.

Test steel plate No. In No. 23, since the value of (Ti + Nb / 2) exceeds the upper limit of the range defined in the present invention, precipitates in the steel are coarsened, so that the tensile strength is low.
Test steel plate No. In No. 26, since the B content exceeds the upper limit of the specified range in the present invention, borides at the grain boundaries are precipitated and the bendability is unsatisfactory.

Test steel plate No. No. 29 has poor bendability because the Mn content exceeds the upper limit of the specified range in the present invention.
Test steel plate No. In No. 31, since the annealing temperature in continuous annealing is below the lower limit of the range defined in the present invention, unrecrystallized ferrite remains, a band-like structure is formed, and the bendability is poor.

Test steel plate No. Since No. 32 does not hold the austenite single phase structure in the continuous annealing for the time in the specified range in the present invention, the bendability is poor.
Test steel plate No. Since the value of (Ti + Nb / 2) is lower than the lower limit of the range defined in the present invention, the crystal grains are not refined and the bendability is poor.

Test steel plate No. No. 38 has poor bendability because the C content exceeds the upper limit of the specified range in the present invention.
Furthermore, the test No. In No. 42, since the value of (Ti + Nb / 2) is below the lower limit of the range defined in the present invention, the crystal grains are not refined and the bendability is poor.

Claims (4)

In mass%, C: 0.05 to 0.17%, Mn: 2.0 to 3.0%, B: 0.0005 to 0.01%, P: 0.03% or less, S: 0.01 %: N: 0.01% or less, Si: 0.005-0.5%, sol. Al: 0.01 to 0.1% of both are contained in a total amount of 0.5% or less, and one or two of Ti and Nb are contained within the range satisfying the following formula (1), and the balance is Fe And a chemical composition comprising impurities, not containing unrecrystallized ferrite, having a steel structure in which the average grain size of ferrite and bainite is 3.5 μm or less, and argon at a depth of 2 to 10 μm from the surface. The minimum value of B intensity by glow discharge emission spectrometry using sputtering is 60% or less of the B intensity by glow discharge emission spectroscopy using argon sputtering at a depth of 80 μm from the surface, and the tensile strength is 780 MPa. A high-strength steel sheet characterized by the above.
0.05% ≦ [Ti] + [Nb] /2≦0.2% (1)
In the formula (1), [Ti] and [Nb] indicate the contents (mass%) of Ti and Nb, respectively.   Further, the chemical composition is, in mass%, Cr: 1.0% or less, Mo: 0.5% or less, Ni: 0.5% or less, Cu: 0.5% or less, and V: 0.2% or less. The high-strength steel sheet according to claim 1, comprising one or more selected from the group consisting of:   The chemical composition is 1% selected from the group consisting of Ca: 0.003% or less, Mg: 0.003% or less, REM: 0.01% or less, and Zr: 0.01% or less in terms of mass%. The high-strength steel sheet according to claim 1 or 2 containing seeds or two or more kinds. The manufacturing method of the high strength steel plate characterized by including the following process (A)-(C):
(A) Hot rolling is performed after setting the steel material having the chemical composition described in any one of claims 1 to 3 to 1100 to 1300 ° C, and the heat is applied in a temperature range of 800 to 950 ° C. A hot rolling step in which the hot rolling is finished and wound into a hot rolled steel sheet at a temperature range of 450 to 700 ° C;
(B) was performed pickling the hot-rolled steel sheet, cold rolling the cold-rolled steel sheet by performing cold rolling; a and (C) the cold-rolled steel sheet, N ₂ concentration is more than 95 vol% Atmosphere in 600 ° C. to Ac ₃ temperature range of the transformation point over 20 seconds to a temperature range of not lower than Ac ₃ transformation point to an austenite single-phase structure, N ₂ concentration austenite single-phase structure in an atmosphere of more than 95 vol% A continuous annealing step in which, after holding in this state for 60 seconds or more, cooling is performed under cooling conditions in which the average cooling rate is 5 ° C./second or more from the Ar ₃ transformation point to 550 ° C.