JP2015105383A - High strength steel material excellent in impact absorption property - Google Patents

Abstract

PROBLEM TO BE SOLVED: To provide a high strength steel sheet which has high impact absorption energy, in which the occurrence of crack upon being applied with impact load is suppressed and which has good formability such as bendability.SOLUTION: The high strength steel sheet has a chemical composition containing, by mass%, C: over 0.3% and 0.6% or less, Mn:0.4% or more and less than 1.2%, Si: over 0.6% and 1.8% or less, Ti: 0.001% or more and less than 0.01%, Al:0.01 to 0.5%, N:0.001 to 0.015% and the balance Fe with inevitable impurities. The high strength steel sheet has a dual-phase structure which includes a bainite phase and/or a tempered martensite phase having a nano hardness of 9 GPa or less in an amount of 80% or more by an area ratio in total and the balance comprising ferrite, austenite and cementite, and which has an average particle diameter of 10 μm or less of a prior austenite particle. The high strength steel sheet has a tensile strength of 980 MPa or more.

Description

  The present invention relates to a steel material, and more specifically, relates to a high-strength steel material of 980 MPa or higher that is suitable as a material for an impact-absorbing member having a high effective flow stress, in which the occurrence of cracks when an impact load is applied is suppressed.

In recent years, from the viewpoint of protecting the global environment, as a part of reducing CO ₂ emissions from automobiles, the weight reduction of automobile bodies has been demanded, and the strength of automobile steel materials has been increased. This is because the steel material for automobiles can be thinned by improving the strength of the steel material. On the other hand, social demands for improving the collision safety of automobiles are increasing, and it is desired not only to increase the strength of steel materials, but also to develop steel materials that have excellent impact resistance when colliding while traveling. Yes.

Here, since each site | part of the steel material for motor vehicles at the time of a collision receives a deformation | transformation with the high strain rate of several tens (s < ^-1 >) or more, the high strength steel material excellent in the dynamic strength characteristic is requested | required.

  Such high-strength steel materials include high-strength composites such as low-alloy TRIP steel, which has a high static difference (difference between static strength and dynamic strength), and multi-phase structure steel having a second phase mainly composed of martensite. Phase structure steels are known.

  As for the low alloy TRIP steel, for example, Patent Document 1 discloses a work-induced transformation type high strength steel plate (TRIP steel plate) for absorbing automobile collision energy that has excellent dynamic deformation characteristics.

  Moreover, the following invention is disclosed regarding the multiphase-structure steel plate which has the 2nd phase which has martensite as a main body.

  Patent Document 2 discloses an average grain size ds of nanocrystal grains made of fine ferrite grains and having a crystal grain size of 1.2 μm or less, and an average grain diameter dL of microcrystal grains having a crystal grain size exceeding 1.2 μm. Discloses a high-strength steel sheet that satisfies the relationship dL / ds ≧ 3, is excellent in strength and ductility balance, and has a static difference of 170 MPa or more.

  Patent Document 3 discloses a steel plate having a high static ratio, which is composed of a two-phase structure of martensite having an average particle diameter of 3 μm or less and ferrite having an average particle diameter of 5 μm or less.

  Patent Document 4 discloses a cold-rolled steel sheet that contains 75% or more of a ferrite phase having an average particle size of 3.5 μm or less, and the balance is made of tempered martensite and has excellent shock absorption characteristics.

In Patent Document 5, a pre-strained two-phase structure composed of ferrite and martensite is formed, and a cold rolling in which a static difference at a strain rate of 5 × 10 ^{2 to} 5 × 10 ³ / s satisfies 60 MPa or more. A steel sheet is disclosed.

  Furthermore, Patent Document 6 discloses a high-strength hot-rolled steel sheet having excellent impact resistance properties composed of only a hard phase such as 85% or more of bainite and martensite.

Japanese Patent Laid-Open No. 11-80879 JP 2006-161077 A JP 2004-84074 A JP 2004-277858 A JP 2000-17385 A JP-A-11-269606

  However, steel materials that are materials of conventional shock absorbing members have the following problems. That is, in order to improve the impact absorption energy of an impact absorbing member (hereinafter also simply referred to as “member”), the strength of a steel material (hereinafter also simply referred to as “steel material”) that is a material of the impact absorbing member is increased. It is essential.

However, in "Plasticity and processing", Vol. 46, No. 534, pages 641-645, the average load (F _ave ) that determines the impact absorption energy is
F _ave ∝ (σY · t ² ) / 4
σY: Effective flow stress
As disclosed that t is given as a plate thickness, the impact absorption energy greatly depends on the plate thickness of the steel material. Therefore, there is a limit in achieving both a reduction in the thickness of the impact absorbing member and a high impact absorbing performance by simply increasing the strength of the steel material.

  By the way, as disclosed in, for example, International Publication No. 2005/010396 pamphlet, International Publication No. 2005/01097 pamphlet, and further International Publication No. 2005/010398 pamphlet, the impact absorbing energy of the impact absorbing member is in its shape. Also depends heavily on.

  In other words, by optimizing the shape of the shock absorbing member so as to increase the work of plastic deformation, the shock absorbing energy of the shock absorbing member can be dramatically reduced to a level that cannot be achieved simply by increasing the strength of the steel material. There is a possibility that can be increased.

  However, even if the shape of the shock absorbing member is optimized so as to increase the plastic deformation work, if the steel material does not have a deformability capable of withstanding the plastic deformation work, the assumed plastic deformation is completed. Before this, the impact absorbing member is cracked at an early stage. As a result, the amount of plastic deformation work cannot be increased, and the impact absorbing energy cannot be dramatically increased. Moreover, if a crack occurs in the impact absorbing member at an early stage, an unexpected situation such as damage to other members disposed adjacent to the impact absorbing member may occur.

  Conventionally, it has been directed to increase the dynamic strength of steel based on the technical idea that the impact absorption energy of the impact absorbing member depends on the dynamic strength of the steel, but simply increasing the dynamic strength of the steel. Orientation may lead to a significant decrease in deformability. For this reason, even if the shape of the shock absorbing member is optimized so as to increase the work of plastic deformation, the shock absorbing energy of the shock absorbing member cannot always be dramatically increased.

  In addition, since the shape of the shock absorbing member has been studied on the premise that the steel material manufactured based on the above technical idea is used, the optimization of the shape of the shock absorbing member is based on the deformability of the existing steel material from the beginning. Until now, studies have been made to improve the deformability of the steel material and optimize the shape of the shock absorbing member so as to increase the work of plastic deformation.

  Furthermore, the deformability of the steel material decreases significantly with increasing strength. Therefore, there is no precedent for the application of a high-strength material of 980 MPa or more to an impact absorbing member.

  As described above, in order to increase the shock absorbing energy of the shock absorbing member, it is important to optimize the shape of the shock absorbing member together with the characteristics of the steel material so as to increase the work of plastic deformation.

  Regarding steel materials, it is important to increase the effective flow stress so as to increase the work of plastic deformation while suppressing the occurrence of cracking when an impact load is applied.

  In order to propose a high impact absorbing steel excellent in crack resistance, the present inventors diligently studied a method for suppressing the occurrence of impact cracking while increasing impact absorbing energy, and obtained new knowledge listed below. .

[Improvement of shock absorption energy]
(A) In order to increase the impact absorption energy of the steel material, it is effective to increase the yield strength and the work hardening coefficient in the low strain region to improve the effective flow stress.

  (B) In order to improve the yield strength, it is effective to use a bainite phase or a tempered martensite phase as the main phase.

  (C) Increasing the carbon content is effective for improving the work hardening coefficient in the low strain region of the bainite main phase steel.

[Inhibition of impact cracking]
(D) In the impact absorbing material, when cracking occurs when an impact load is applied, not only the energy absorbing ability is lowered, but also causes damage to other parts in the engine.

  (E) However, as the material strength, particularly the yield strength, increases, the sensitivity to cracking (hereinafter also referred to as “impact cracking sensitivity”) at the time of impact heavy load increases.

  (F) In order to suppress the occurrence of embrittlement cracks when an impact load is applied, it is effective to refine the prior austenite grains and improve the fracture toughness.

  (G) Furthermore, it is effective to improve uniform elongation and local ductility (hole expansibility) in order to suppress cracking of the buckled deformation portion when an impact load is applied.

  (H) In order to improve the uniform elongation, it is effective to mix an austenite phase.

  (I) On the other hand, the fresh martensite phase adversely affects the local ductility, so it is necessary to reduce it as much as possible.

  (J) When the amount of Mn in the steel is large, a hard martensite phase is generated due to Mn microsegregation, which adversely affects local ductility. Therefore, it is necessary to reduce the Mn content as much as possible.

  The present invention has been made on the basis of the above new findings, and the gist thereof is as follows.

(1) By mass%, C: more than 0.3% and 0.6% or less, Mn: 0.4% or more and less than 1.2%, Si: more than 0.6% and 1.8% or less, Ti: 0.00. 001% or more and less than 0.01%, Al: 0.01 to 0.5%, N: 0.001 to 0.015%, having a chemical composition consisting of Fe and impurities,
A multiphase structure in which the total amount of bainite phase and / or tempered martensite phase with a nanohardness of 9 GPa or less is 80% or more in area ratio, and the former austenite grains with the balance consisting of ferrite, austenite and cementite have an average particle size of 10 μm or less A steel sheet having a tensile strength of 980 MPa or more.

  In the present invention, the “area ratio” represents the area ratio of each metal structure obtained from a two-dimensional image using the nanoindentation method in an arbitrary cross section.

  (2) By mass%, one or more of Cr: 0.05 to 0.25%, Mo: 0.1 to 0.35%, B: 0.001 to 0.005% should be contained. The steel plate according to (1) above, characterized in that

  (3) The method for manufacturing a steel sheet according to (1) or (2) above, comprising the following steps.

  After heating the slab having the chemical composition described in (1) or (2) above in a temperature range of 1200 ° C. or higher, a finishing hot having a total rolling reduction of 50% or higher in a temperature range of 800 to 900 ° C. A hot rolling process in which a steel sheet is rolled and cooled to a temperature range of 600 to 700 ° C. at an average cooling rate of 500 ° C./second or more and wound up at room temperature to form a hot-rolled steel sheet within 0.4 seconds after the rolling. .

  (4) The method for manufacturing a steel sheet according to (1) or (2) above, comprising the following steps.

  After heating the slab having the chemical composition described in (1) or (2) above in a temperature range of 1200 ° C. or higher, a finishing hot having a total rolling reduction of 50% or higher in a temperature range of 800 to 900 ° C. After rolling and cooling to a temperature range of 600 to 700 ° C. at an average cooling rate of 500 ° C./second or more within 0.4 seconds after the rolling, winding is performed at room temperature, and the temperature range is 500 ° C. or less. A hot rolling process for holding a hot-rolled steel sheet.

  (5) The method for manufacturing a steel sheet according to (1) or (2) above, comprising the following steps.

After heating the slab having the chemical composition described in (1) or (2) above in a temperature range of 1200 ° C. or higher, a finishing hot having a total rolling reduction of 50% or higher in a temperature range of 800 to 900 ° C. A hot rolling step of rolling, cooling at an average cooling rate of 50 ° C./second or more and less than 500 ° C./second, and winding up in a temperature range of 300 to 600 ° C. to obtain a hot-rolled steel sheet;
A cold rolling process in which a hot rolling steel sheet obtained in the hot rolling process is subjected to cold rolling with a total rolling reduction of 30 to 70% to obtain a cold rolling steel sheet; and the cold rolling process obtained in the cold rolling process The steel sheet was heated to a temperature range of 800 to 950 ° C. at an average temperature increase rate of 10 to 100 ° C./second and held for 100 to 500 seconds, and then a temperature of 400 to 560 ° C. at an average cooling rate of 10 to 50 ° C./second. A heat treatment step of performing a heat treatment for cooling to a temperature range and maintaining the temperature range for 50 to 300 seconds.

  (6) The method for manufacturing a steel sheet according to (1) or (2) above, comprising the following steps.

After heating the slab having the chemical composition described in (1) or (2) above in a temperature range of 1200 ° C. or higher, a finishing hot having a total rolling reduction of 50% or higher in a temperature range of 800 to 900 ° C. A hot rolling step of rolling, cooling at an average cooling rate of 50 ° C./second or more and less than 500 ° C./second, and winding at 300 to 600 ° C. to form a hot-rolled steel sheet;
A cold rolling process in which a hot rolling steel sheet obtained in the hot rolling process is subjected to cold rolling with a total rolling reduction of 30 to 70% to obtain a cold rolling steel sheet; and the cold rolling process obtained in the cold rolling process The steel sheet is heated to a temperature range of 800 to 950 ° C. at an average temperature increase rate of 10 to 100 ° C./second and held for 100 to 500 seconds, and then 200 to 330 ° C. at an average cooling rate of 10 to 50 ° C./second. Cool to a temperature range, hold in the temperature range for 50 to 200 seconds, reheat to a temperature range of 400 to 560 ° C. at a temperature increase rate of 5 to 20 ° C./second, and hold in the temperature range for 50 to 200 seconds. Heat treatment process for applying heat treatment.

  The steel sheet according to the present invention is suitable as a material for an impact absorbing member, particularly an automobile impact absorbing member, and as an automobile impact absorbing member, it is particularly used as a material for side members, center pillars, lockers, and the like. preferable.

  According to the present invention, it is possible to suppress or eliminate the occurrence of cracks in the impact absorbing member when an impact load is applied, and to obtain an impact absorbing member having a high effective flow stress. It is possible to dramatically increase the absorbed energy. By applying such an impact absorbing member, it is possible to further improve the collision safety of the product, which is extremely useful in the industry.

FIG. 1 is a diagram showing an annealing cycle defined by the manufacturing method according to the present invention.

The present invention will be described below.
1. Chemical composition (1) C: more than 0.3% and not more than 0.6% C is a basic element for improving the strength of steel materials. C contributes to increasing the strength of the ferrite phase by solid solution strengthening. Furthermore, with the increase in the amount of C, the bainite phase and martensite phase, which are low-temperature transformation phases, are significantly strengthened. Moreover, in order to stabilize a retained austenite phase, it contributes to the improvement of uniform ductility. Furthermore, fine cementite suppresses the coarsening of crystal grains and improves local ductility. In addition, the work hardening coefficient of the main phase is improved. However, if the C content is 0.3% or less, it may be difficult to obtain the effect of the above action. Therefore, the C content is more than 0.3%. On the other hand, when the C content exceeds 0.6%, not only the above effects are saturated, but also cementite, martensite, and austenite are excessively generated, and the occurrence of cracks when an impact load is applied is promoted. Therefore, the C content is more than 0.3% and not more than 0.6%. Desirably, it is 0.40 to 0.57%. More desirably, it is 0.45 to 0.55%.

(2) Mn: 0.4% or more and less than 1.2% Mn improves hardenability and increases the phase ratio of the bainite phase. Moreover, it has the effect | action which strengthens steel by solid solution strengthening and improves yield strength and tensile strength. However, if the Mn content is less than 0.4%, it may be difficult to obtain the effect by the above action. Therefore, the Mn content is 0.4% or more. On the other hand, when the Mn content is 1.2% or more, martensite and austenite are excessively generated, and the local ductility may be significantly reduced. Therefore, the Mn content is less than 1.2%. Preferably it is 1.0% or less.

(3) Si: more than 0.6% and 1.8% or less Si has an effect of improving uniform ductility by suppressing the formation of carbides in bainite and martensite phases and stabilizing austenite. Moreover, it has the effect | action which strengthens steel by solid solution strengthening and improves yield strength and tensile strength. If the Si content is 0.6% or less, it is difficult to obtain the effect of the above action. Therefore, the Si amount is set to exceed 0.6%. On the other hand, if the Si content exceeds 1.8%, excess austenite remains and the impact cracking sensitivity is increased. Therefore, the Si content is 1.8% or less. A desirable range of Si amount is 0.8 to 1.5%.

(4) Ti: 0.001% or more and less than 0.01% Ti refines the solidified structure and brings about austenite grain refinement after the subsequent heat treatment. However, if the Ti content is less than 0.001%, it is difficult to obtain the above effects. Therefore, the Ti content is 0.001% or more. On the other hand, when the Ti content is 0.01% or more, precipitates such as TiC increase and the toughness decreases, so that cracking is likely to occur at the time of collision. Therefore, the Ti content is 0.001% or more and less than 0.01%. Desirably, it is 0.006% or less.

(5) Al: 0.01 to 0.5%
Al has an effect of suppressing impact cracking by suppressing inclusions in the steel due to a deoxidizing effect. However, if the Al content is less than 0.01%, it is difficult to obtain the effect by the above action. On the other hand, if the Al content exceeds 0.5%, the oxides and nitrides are coarsened, which in turn promotes impact cracking. Therefore, the Al content is set to 0.01 to 0.5%.

(6) N: 0.001 to 0.015%
N produces nitrides, thereby suppressing the grain growth of austenite and ferrite and suppressing the impact cracking. However, when the N content is less than 0.001%, it is difficult to obtain the effect by the above action. On the other hand, when the N content exceeds 0.015%, the nitride is coarsened, which promotes impact cracking. Therefore, the N content is 0.001 to 0.015%.

(7) Cr: 0.05 to 0.25%
Cr improves hardenability and increases the bainite phase ratio. Moreover, it has the effect | action which strengthens steel by solid solution strengthening and improves yield strength and tensile strength. Therefore, it can be contained as necessary.

  However, if the Cr content is less than 0.05%, it is difficult to obtain the effect by the above action. Therefore, the Cr content is 0.05% or more. On the other hand, if the Cr content exceeds 0.25%, the martensite phase is excessively generated and the impact cracking sensitivity is increased. Therefore, the Cr content is 0.05 to 0.25%.

(8) Mo: 0.1% to 0.35%
Mo, like Cr, enhances hardenability and increases the bainite phase ratio. Moreover, it has the effect | action which strengthens steel by solid solution strengthening and improves yield strength and tensile strength. Therefore, it can be contained as necessary.

  However, if the Mo content is less than 0.1%, it is difficult to obtain the effect by the above action. On the other hand, if the Mo content exceeds 0.35%, the martensite phase is excessively generated and the impact cracking sensitivity is increased. Therefore, the Mo content is set to 0.1 to 0.35%.

(9) B: 0.001 to 0.005%
B enhances hardenability and, in the case of a material containing a small amount of Ti, brings about refinement of crystal grains. Therefore, it can be contained as necessary.
However, if the B content is less than 0.001%, it is difficult to obtain the effect by the above action. On the other hand, if the content of B exceeds 0.005%, the grain boundary carbides are coarsened and ductility is lowered. Therefore, the B content is set to 0.001 to 0.005%.

(10) P: 0.02% or less P as an impurity weakens the grain boundary and causes deterioration of hot workability. Therefore, the P content is 0.02% or less. The smaller the P content, the better. However, if it is assumed that the P is removed within the range of realistic manufacturing steps and manufacturing costs, the upper limit of P is 0.02%. Desirably, it is 0.015% or less.

(11) S: 0.005% or less S as an impurity weakens the grain boundary and causes deterioration of hot workability and ductility. Therefore, the S content is 0.005% or less. The smaller the S content, the better. However, the upper limit of S is 0.005% on the assumption that the S content is removed within the range of realistic manufacturing processes and manufacturing costs. Desirably, it is 0.002% or less.

2. Steel structure The steel structure according to the present invention is a multi-phase structure in which the average grain size of prior austenite grains is 10 μm or less in order to increase the effective flow stress by obtaining a high yield strength and a work hardening coefficient in a low strain region. The total amount of the bainite phase and / or the tempered martensite phase composing the alloy is 80% or more by area ratio, and the balance is composed of ferrite, austenite, and cementite.

(1) The total amount of bainite phase and / or tempered martensite phase with a nanohardness of 9 GPa or less is 80% or more by area ratio. The bainite phase and / or tempered martensite phase is work-hardened in the yield stress and low strain regions. Increase the coefficient. However, when the total amount of the bainite phase and the tempered martensite phase is less than 80% in terms of area ratio, the yield stress and the work hardening coefficient in the low strain region are low, and the predetermined shock absorption characteristics are not satisfied. Therefore, the total amount of the bainite phase and the tempered martensite phase is 80% or more by area ratio. Desirably, it is 80 to 95%.

  Moreover, the phase whose nano hardness exceeds 9 GPa is hard martensite which is insufficiently tempered and inhibits ductility and toughness. Therefore, the nano hardness is set to 9 GPa or less. Therefore, steel materials containing a hard martensite phase are out of the scope of the invention.

  Here, each phase excluding cementite can be distinguished by nano hardness and nano Young's modulus measurement by the nanoindentation method. That is, a phase having a nano hardness of less than 4.5 GPa is a ferrite phase, a phase having a hardness of 4.5 to 9 GPa is a bainite phase or a tempered martensite phase, and a phase having a nano Young's modulus of 190 GPa or less is austenite. The lower limit is preferably 4.5 GPa because the phase having a nano hardness of less than 4.5 GPa is considered to be a ferrite phase. Since pearlite has a specific lamellar structure, it can be distinguished by SEM and optical microscope observation, and the area ratio is obtained by calculating the area of the lamellar structure in an arbitrary cross section. The cementite phase can be easily distinguished from other phases by the light and dark contrast by SEM observation.

(2) A dual phase structure in which the average grain size of the prior austenite grains composed of ferrite, austenite, and cementite is 10 μm or less, and the balance is composed of ferrite, austenite, and cementite in order to improve uniform ductility. However, in order to combine local ductility, the average crystal grain size of ferrite and austenite is preferably 1.5 μm or less, and the cementite particle size is preferably 200 nm or less.

  Further, when bainite or tempered martensite is the main phase, the prior austenite grains are often the fracture surface unit for brittle fracture. When the average grain size of the prior austenite grains exceeds 10 μm, brittle fracture starting from the prior austenite grain boundaries is likely to occur at the time of collision. Therefore, the average grain size of the prior austenite grains is set to 10 μm or less.

  The prior austenite grain boundaries are specified by variant inverse analysis of the bainite phase or tempered martensite phase, and the prior austenite grain size is determined. However, it is difficult to calculate the prior austenite grain size in steels in which the bainite phase or tempered martensite phase is not the main phase.

  When the pearlite phase is generated, the strength is lowered and the toughness is reduced. Therefore, it is necessary to prevent the pearlite phase from being generated as much as possible. The area ratio of pearlite is desirably 5% or less.

  A hard martensite phase (fresh martensite phase) that is insufficiently tempered reduces local ductility and must be reduced as much as possible. However, since it is difficult to distinguish between a bainite phase, a tempered martensite phase, and a hard martensite phase in high carbon steel, a phase having a nano hardness exceeding 10 GPa by the nanoindentation method is defined as a hard martensite phase. The area ratio of the site is desirably 5% or less.

  The area ratios of ferrite, austenite, and cementite constituting the balance of the multiphase structure are preferably 5 to 10%, 3 to 7%, and 0.5 to 2.5%, respectively, in order to increase local ductility. In order to further increase the local ductility, it is more desirable that each area ratio satisfies the relationship of ferrite> austenite> cementite.

3. Characteristics The steel formed from the steel according to the present invention is characterized in that the effective flow stress is high, the impact absorption energy is high, and at the same time, the occurrence of cracks when an impact load is applied is suppressed. This feature is demonstrated by a high 5% flow stress, a high hole expansion rate, and excellent axial crushing characteristics, as shown in the examples described later.

  Other mechanical properties include a tensile strength of 980 MPa or more, a uniform elongation of 7% or more, high strength and excellent ductility.

4). Manufacturing method (1) Hot rolling process A hot-rolled steel sheet is obtained by hot-rolling to the slab which has the above-mentioned chemical composition and was heat-processed beforehand at 1200 degreeC or more. When the hot-rolled steel sheet is used as it is as a base material for the heat treatment step, it is reheated one or more times in a temperature range of 1200 ° C. or higher in the hot rolling step, and then subjected to a final rough rolling and a final finish rolling of 800 to 900 In the temperature range of ° C., it is carried out by multi-stage rolling with a total rolling reduction of 50% or more, and then cooled within a range of 600 to 700 ° C. within 0.4 seconds at an average cooling rate of 500 ° C./second or more, then room temperature Take up with.

  The slab heating temperature needs to be 1200 ° C. or higher. When the slab heating temperature is less than 1200 ° C., Ti and V precipitates precipitated during casting or slab cooling do not dissolve in a sufficient amount, resulting in insufficient precipitation strengthening and ensuring a tensile strength of 980 MPa or more. I can't.

  Since the solid solution Ti, V delays the ferrite transformation, a structure having a bainite phase of 80% or more can be easily obtained in subsequent cooling after hot rolling.

  A large amount of work strain is accumulated in the austenite by performing hot rolling in which the total rolling reduction in the temperature range of 800 to 900 ° C. is 50% or more. By cooling to 600-700 ° C. at an average cooling rate of 500 ° C./second or more within 0.4 seconds after completion of hot rolling, ferrite transformation occurs while suppressing release of the processing strain. However, in steels containing a large amount of carbon, such as the steel of the present invention, carbon and alloy elements are concentrated at the ferrite transformation interface, and ferrite growth is suppressed. That is, fine ferrite is generated by holding for 0.4 seconds or more in the temperature range of 600 to 700 ° C., and in the subsequent cooling process, both the ferrite interface and the prior austenite interface are used as nucleation sites, and the bainite transformation is performed. Can be advanced.

  By the above hot rolling, a predetermined multiphase structure in which the phase including the bainite phase is 80 area% or more and the prior austenite grain size is 10 μm or less is obtained. However, when the above hot rolling conditions are not satisfied, the prior austenite grains exceed 10 μm, and a desired multiphase structure cannot be obtained, so that the strength, ductility, and impact absorption characteristics are not satisfied.

  When the hot-rolled steel sheet after cooling contains a martensite phase, tempering is performed as necessary to form tempered martensite. However, if the tempering temperature exceeds 500 ° C., the strength and ductility may be lowered. Therefore, tempering needs to be heat-treated at 500 ° C. or less. The holding time is desirably 5 minutes or longer. When the holding time is less than 5 minutes, the transformation strain is not released and the ductility is significantly reduced. On the other hand, the holding time is preferably within 30 minutes. When holding time exceeds 30 minutes, a carbide | carbonized_material will produce | generate at a grain boundary and ductility may fall by what is called temper embrittlement.

  When using a hot-rolled steel sheet as a base material for cold rolling, the hot-rolled base material is preferably a ferrite main phase. Therefore, after reheating once or more in the temperature range of 1200 ° C. or higher in the hot rolling step, the final finish rolling is performed in a temperature range of 800 to 900 ° C. in a temperature range of 800 to 900 ° C., and the total rolling reduction is 50% or more. It is necessary to carry out the multi-stage rolling, then cool at an average cooling rate of 50 ° C./second or more and less than 500 ° C./second, and then perform winding at 300 to 600 ° C. The reason why the cooling rate is set in this range when the hot-rolled steel sheet is used as a base material for cold rolling is to make the cold-rolled base material soft. The reason for setting the coiling temperature in this temperature range is to soften the bainite or martensite phase and prevent cracking during subsequent cold rolling.

(2) Cold rolling step In order to obtain a dual phase structure having a prior austenite grain size of 10 µm or less by subsequent cold rolling and annealing, the hot rolled steel plate is subjected to cold rolling and continuous annealing, When doing, let the reduction rate in cold rolling be 30 to 70%. Processing strain is accumulated by setting the rolling reduction in cold rolling to 30 to 70%, the temperature is raised to a temperature range of 800 ° C. to 950 ° C. at an average heating rate of 10 to 100 ° C./second, and 100 to 500 By holding for 2 seconds, a fine steel structure can be obtained by recrystallization.

(3) Heat treatment process (annealing heat treatment process)
The heat treatment step employed in the production method of the present invention is performed by a thermal cycle as shown in FIG. In the annealing heat treatment after hot rolling or after cold rolling, the temperature is raised to a temperature range of 800 to 950 ° C. at an average temperature rising rate of 10 to 100 ° C./second and held for 100 to 500 seconds, and then a predetermined cooling control is selected. There is a need to. If the cold rolling and annealing heat treatment conditions do not satisfy the above, the prior austenite grains become coarse, and ductility and toughness decrease. Further, in order to obtain a multiphase structure in which the total amount of the bainite phase and / or the tempered martensite phase is 80% or more in area ratio and the balance is composed of ferrite, austenite and cementite, the average cooling rate after the annealing is performed. It is necessary to cool to a temperature of 400 to 560 ° C. at a rate of 10 to 50 ° C./second, and perform an annealing treatment by giving a thermal history of 50 to 300 seconds at a constant temperature in the range of 400 to 560 ° C. In particular, the structure of the bainite main phase can be generated by cooling to a temperature range of 400 to 560 ° C. or less, which is a bainite generation region, at an average cooling rate of 10 to 50 ° C./second or more.

  Furthermore, in order to increase the strength exceeding 1180 MPa, it is necessary to use a bainite phase and a tempered martensite phase composed of a dense lath structure by transforming austenite at a low temperature. For this purpose, after annealing, it is cooled to a temperature of 200 to 330 ° C. at an average cooling rate of 10 to 50 ° C./second, subjected to transformation treatment by giving a thermal history of 50 to 200 seconds, and then increased to 5 to 20 ° C./second. It is necessary to reheat to a temperature range of 400 to 560 ° C. at a temperature rate and give a heat history of 50 to 200 seconds to perform reheating treatment.

  When the heat treatment condition is outside the above condition range, a desired multiphase structure cannot be obtained, and predetermined ductility and impact characteristics cannot be obtained.

  Experiments were performed using slabs having a chemical composition shown in Table 1 (thickness 35 mm, width 160 to 250 mm, and length 70 to 140 mm). In either case, 150 kg of molten steel was vacuum-melted and cast, and then heated at a furnace temperature of 1250 ° C. and hot forged at a temperature of 950 ° C. or higher to form a slab.

  The slab was reheated at 1250 ° C. for 1 hour, then subjected to 4 passes of rough hot rolling using a hot rolling tester, and further subjected to 3 passes of finish hot rolling to obtain a hot rolled steel sheet. Table 2 shows the hot rolling conditions.

  Furthermore, about some hot-rolled steel plates, after performing cold rolling, continuous annealing heat processing was performed using the continuous annealing simulator. These conditions are shown in Table 3.

  In Table 3, when reheating was performed, the temperature was increased from the transformation treatment temperature to the reheating temperature at a temperature increase rate of 5 to 20 ° C./second.

The following investigation was performed on the hot-rolled steel sheet and the cold-rolled steel sheet thus obtained.
That is, 5% flow stress, maximum tensile strength (TS), and uniform elongation (U-El) were obtained by collecting JIS No. 5 tensile test pieces and conducting a tensile test. The hole expansibility was evaluated by reaming the machined holes in order to remove the influence of end face damage.

  An EBSD analysis was performed on a ¼t portion having a cross section parallel to the rolling direction of the steel sheet, and the prior austenite grain size was calculated from variant analysis of bainite and martensite.

  The hardness of the bainite and tempered martensite phases was determined by the nanoindentation method. A 1/4 t part of the plate thickness was polished with emery paper, then mechanochemically polished with colloidal silica, and the processed layer was removed by electrolytic polishing for use in the test. Nanoindentation was performed using a cube corner indenter and an indentation load of 500 μN. The indentation size at this time is 0.5 μm or less in diameter. The total area ratio of the bainite phase and the tempered martensite phase was determined from a two-dimensional image obtained by an atomic force microscope attached to the nanoindentation apparatus.

  Nanohardness measurement was also used for the distinction of ferrite phases. That is, a phase having a nano hardness of less than 4.5 GPa was a ferrite phase. Furthermore, residual austenite was separated and distinguished by measuring the nano Young's modulus by the nano-indentation method. That is, the phase having a nano Young's modulus of 190 GPa or less is austenite. Thus, the area ratio of ferrite and austenite and the average particle diameter of ferrite were obtained. Since pearlite has a unique lamellar structure, it can be distinguished by SEM and optical microscope observation. The area ratio of the pearlite was calculated | required by calculating the area | region of a lamellar structure | tissue in arbitrary cross sections. The cementite phase can be discriminated by light / dark contrast by SEM observation. The area ratio of cementite was obtained by calculating the region of the cementite phase in an arbitrary cross section.

  Furthermore, an octagonal cylindrical member was produced using the steel plate, an axial crush test was performed at an axial collision speed of 64 km / h, and the impact absorption performance was evaluated. In order to remove the surface decarburized layer, the hot-rolled steel sheet was ground on both sides to a thickness of 1.6 mm. The cold rolled steel sheet was subjected to the test as it was. The shape of the cross section perpendicular to the axial direction of the rectangular tube member was a regular octagon. The flat portion width Wp was molded into two types of members having a size of 25.6 mm and 32.0 mm.

  Table 4 shows the structures and evaluation results of invention steel and comparative steel.

  Here, the impact absorption energy index (Epa) is a parameter obtained by calculating the average stress applied to the rectangular tube member at the time of buckling, and normalized by the tensile strength of the steel sheet, and is defined by the following formula (1). .

Here, Load F _ave is the average load applied to the member, L is the circumference of the regular polygon, t is the plate thickness of the steel sheet, and TS is the maximum tensile strength measured in the tensile test described above. is there.

  The stable buckling number is the number of test specimens that were not cracked by the axial crush test. Table 4 shows the total number of specimens and the number of specimens that did not crack (no cracking occurred). Number of test specimens) / (total number of test specimens).

  In general, when the collision energy increases, the possibility of cracking during crushing increases, and as a result, the amount of plastic deformation work cannot be increased and the impact absorption energy may not be increased.

The steel material according to the present invention has a high average load of 0.38 kJ / mm ² or more due to axial crushing. Furthermore, the stable buckling number is 2/2, indicating a good axial crushing characteristic. Therefore, it is suitable for use as a material for impact absorbing members such as side members, center pillars, rockers and the like.

Claims (6)

In mass%, C: more than 0.3% and 0.6% or less, Mn: 0.4% or more and less than 1.2%, Si: more than 0.6% and 1.8% or less, Ti: 0.001% or more Containing less than 0.01%, Al: 0.01-0.5%, N: 0.001-0.015%, having a chemical composition consisting of the balance Fe and impurities,
A multiphase structure in which the total amount of bainite phase and / or tempered martensite phase with a nanohardness of 9 GPa or less is 80% or more in area ratio, and the former austenite grains with the balance consisting of ferrite, austenite and cementite have an average particle size of 10 μm or less A steel sheet having a tensile strength of 980 MPa or more.   It is characterized by containing one or more of Cr: 0.05 to 0.25%, Mo: 0.1 to 0.35%, and B: 0.001 to 0.005% in mass%. The steel plate according to claim 1. The method for manufacturing a steel sheet according to claim 1, comprising the following steps.
After heating the slab having the chemical composition described in claim 1 or 2 in a temperature range of 1200 ° C or higher, finish hot rolling is performed in a temperature range of 800 to 900 ° C with a total rolling reduction of 50% or higher. A hot rolling step in which the steel sheet is cooled to a temperature range of 600 to 700 ° C. at an average cooling rate of 500 ° C./second or more within 0.4 seconds after the rolling, and then wound at room temperature to form a hot-rolled steel sheet. The method for manufacturing a steel sheet according to claim 1, comprising the following steps.
After heating the slab having the chemical composition described in claim 1 or 2 in a temperature range of 1200 ° C or higher, finish hot rolling is performed in a temperature range of 800 to 900 ° C with a total rolling reduction of 50% or higher. Then, within 0.4 seconds after the rolling, after cooling to a temperature range of 600 to 700 ° C. at an average cooling rate of 500 ° C./second or more, winding is performed at room temperature, and kept at a temperature range of 500 ° C. or less. Hot rolling process to make hot rolled steel sheet. The method for manufacturing a steel sheet according to claim 1, comprising the following steps.
After heating the slab having the chemical composition described in claim 1 or 2 in a temperature range of 1200 ° C or higher, finish hot rolling is performed in a temperature range of 800 to 900 ° C with a total rolling reduction of 50% or higher. , A hot rolling step in which the steel sheet is cooled at an average cooling rate of 50 ° C./second or more and less than 500 ° C./second and is wound in a temperature range of 300 to 600 ° C. to obtain a hot-rolled steel sheet;
A cold rolling process in which a hot rolling steel sheet obtained in the hot rolling process is subjected to cold rolling with a total rolling reduction of 30 to 70% to obtain a cold rolling steel sheet; and the cold rolling process obtained in the cold rolling process The steel sheet was heated to a temperature range of 800 to 950 ° C. at an average temperature increase rate of 10 to 100 ° C./second and held for 100 to 500 seconds, and then a temperature of 400 to 560 ° C. at an average cooling rate of 10 to 50 ° C./second. A heat treatment step of performing a heat treatment for cooling to a temperature range and maintaining the temperature range for 50 to 300 seconds. The method for manufacturing a steel sheet according to claim 1, comprising the following steps.
After heating the slab having the chemical composition described in claim 1 or 2 in a temperature range of 1200 ° C or higher, finish hot rolling is performed in a temperature range of 800 to 900 ° C with a total rolling reduction of 50% or higher. , A hot rolling step of cooling at an average cooling rate of 50 ° C./second or more and less than 500 ° C./second and winding up at 300 to 600 ° C. to obtain a hot-rolled steel sheet;
A cold rolling process in which a hot rolling steel sheet obtained in the hot rolling process is subjected to cold rolling with a total rolling reduction of 30 to 70% to obtain a cold rolling steel sheet; and the cold rolling process obtained in the cold rolling process The steel sheet is heated to a temperature range of 800 to 950 ° C. at an average temperature increase rate of 10 to 100 ° C./second and held for 100 to 500 seconds, and then 200 to 330 ° C. at an average cooling rate of 10 to 50 ° C./second. Cool to a temperature range, hold in the temperature range for 50 to 200 seconds, reheat to a temperature range of 400 to 560 ° C. at a temperature increase rate of 5 to 20 ° C./second, and hold in the temperature range for 50 to 200 seconds. Heat treatment process for applying heat treatment.

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