JP6379731B2 - High-strength steel material and manufacturing method thereof - Google Patents

Description

  The present invention relates to a high-strength steel material and a method for producing the same, and more particularly, to a high-strength steel material exceeding 1400 MPa that is suitable as a material for an impact-absorbing member and capable of producing cracks when an impact load is applied and a method for producing the same.

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 in the event of a collision while traveling. .

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.

  As such a high-strength steel material, a low alloy processing-induced transformation type (TRIP type) steel having a high static difference (difference between static strength and dynamic strength), or a duplex phase mainly composed of martensite. High-strength duplex steels such as microstructure steel are known.

  Regarding the low alloy TRIP type steel, for example, Patent Document 1 discloses a TRIP type high strength steel plate having excellent dynamic deformation characteristics.

  Moreover, the following invention is disclosed regarding the multi-phase structure steel plate which has a 2nd phase mainly composed of martensite.

  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 size 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, has an excellent balance between strength and ductility, and has a static difference of 170 MPa or more.

  Patent Document 3 discloses a high-tensile hot-rolled steel sheet for automobiles that has a two-phase structure of martensite having an average particle diameter of 3 μm or less and martensite having an average particle diameter of 5 μm or less and has a high static ratio and excellent impact resistance. It is disclosed. Patent Document 4 discloses a cold-rolled steel sheet containing 75% or more of a ferrite phase having an average particle diameter of 3.5 μm or less and having a remaining structure substantially made of tempered martensite and excellent in impact absorption characteristics. Yes.

Patent Document 5 discloses a cold rolling in which a pre-deformation is applied to form a two-phase structure composed of ferrite and martensite and 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 International Publication No. 2005/010396 International Publication No. 2005/010397 International Publication No. 2005/010398

Kentaro Sato, Takaaki Hira, Akihide Yoshitake, “High-speed deformation characteristics of automobile parts using high-strength steel sheets”, Plasticity and processing, July 2005, Vol. 46, No. 534, p. 641-645

  A steel material (hereinafter also simply referred to as “steel material”), which is a material of a conventional shock absorbing member, has 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”), it is essential to increase the strength of the steel material.

However, in Non-Patent Document 1, the average load (F _ave ) that determines the impact absorption energy is
F _ave ∝ (σY · t ² ) / 4
As disclosed that σY: effective flow stress t: given as a plate thickness, the impact absorption energy largely 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, for example, as disclosed in Patent Documents 7 to 9, the impact absorption energy of the member greatly depends on its shape. In other words, by optimizing the shape of the member so as to increase the amount of plastic deformation work, it is possible to dramatically increase the shock absorption energy of the member to a level that cannot be achieved simply by increasing the strength of the steel material. There is a possibility.

  However, even if the shape of the 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. In addition, the member is cracked at an early stage, and as a result, the amount of plastic deformation cannot be increased, and the shock absorption energy cannot be dramatically increased. In addition, if a crack occurs in a member at an early stage, an unexpected situation such as damage to another member disposed adjacent to the member may occur.

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

  In addition, since the shape of the 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 member has been studied from the beginning based on the deformability of the existing steel material. However, until now, the study itself of improving the deformability of the steel material and optimizing the shape of the member based on it has not been sufficiently performed so as to increase the work of plastic deformation.

  Furthermore, the deformability of the steel material decreases significantly with increasing strength. Therefore, there has been no precedent for the application of high strength materials exceeding 1400 MPa to impact absorbing members.

  The present invention has been made in order to solve the above-described problems. A high-strength steel material having a tensile strength exceeding 1400 MPa and capable of suppressing the occurrence of impact cracking and excellent in shock absorption characteristics and a method for producing the same are provided. The purpose is to provide.

  In order to propose a high impact-absorbing steel with excellent crack resistance, the present inventors have intensively studied a method for suppressing the occurrence of impact cracking while increasing impact absorption energy, and as a result, have obtained the following knowledge. It was.

[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.

  (B) In order to improve the yield strength, it is effective to make the micro metal structure of the steel a single phase of the lower bainite structure.

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

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

  (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 bainite single-phase steel, embrittlement cracks under impact load are often generated at the prior austenite grain boundaries. Therefore, in order to suppress the occurrence of embrittlement cracks, it is effective to make the prior austenite grain size fine and improve the toughness.

  (G) Furthermore, it is effective to improve elongation, local ductility, and toughness in order to suppress cracking of the buckled deformation portion when an impact load is applied in bainite single phase steel.

  (H) In order to improve the elongation of the bainite single phase steel, it is effective to refine the prior austenite and the bainite block size.

  (I) In order to refine the prior austenite grain boundaries, it is effective to add Mo and B together.

  (J) By containing Mo and B in a composite manner, dislocations and movement of grain boundaries are suppressed in the austenite region, contributing to refinement of the austenite grain size. This is an effect resulting from the boride of Mo and B or the interaction of Mo-B.

  (K) Mo and B are also effective in improving hardenability. When the austenite grains are refined, the hardenability decreases, so that ferrite is generated during cooling, and it is difficult to obtain a bainite single phase structure. However, the hardenability is improved by the combined addition of Mo and B, and a bainite single-phase structure in which the prior austenite grain size is extremely fine can be obtained.

  The present invention has been made on the basis of the above-mentioned knowledge, and the gist thereof is the following high-strength steel material and a manufacturing method thereof.

(1) The chemical composition is mass%,
C: more than 0.2% and 0.55% or less,
Si: 0.1 to 0.3%,
Mn: 0.1 to 0.3%,
Mo: 1-4%
Ti: 0.002 to 0.008%,
Al: 0.01-0.2%
B: 0.001 to 0.006%,
N: 0.001 to 0.015%,
V: 0 to 0.3%
Nb: 0 to 0.05%,
Balance: Fe and impurities,
The metal structure is a single phase of a lower bainite structure having a prior austenite particle size of 5 μm or less,
A high-strength steel material having a tensile strength exceeding 1400 MPa.

(2) The chemical composition is mass%,
V: 0.05-0.3%
Nb: 0.01-0.05%
The high-strength steel material according to (1) above, which contains one or more selected from the above.

  (3) After heating the steel having the chemical composition described in the above (1) or (2) to a temperature range of 800 to 900 ° C., to a temperature range of 350 to 450 ° C. at a cooling rate of 10 ° C./s or more. A method for producing a high-strength steel material comprising a heat treatment step of cooling, holding at that temperature for 10 to 100 seconds, and then cooling.

  (4) The steel having the chemical composition described in the above (1) or (2) is subjected to warm working so that the cross-sectional reduction rate is 20% or more in a temperature range of 500 to 750 ° C., and then 800 to A heat treatment step of heating to a temperature range of 900 ° C., then cooling to a temperature range of 350 to 450 ° C. at a cooling rate of 10 ° C./s or more, holding at that temperature for 10 to 100 s, and then cooling; A manufacturing method for high strength steel.

  (5) The steel having the chemical composition described in (1) or (2) above is subjected to cold working so that the cross-section reduction rate is 20% or more, and then heated to a temperature range of 800 to 900 ° C. Then, it is cooled to a temperature range of 350 to 450 ° C. at a cooling rate of 10 ° C./s or higher, held at that temperature for 10 to 100 s, and then cooled, and then a method for producing a high strength steel material.

  According to the present invention, since the occurrence of cracking of the impact absorbing member when an impact load is applied can be suppressed, the impact absorbing energy of the impact absorbing member can be dramatically increased. Therefore, the high-strength steel material according to the present invention is suitable for use as a material for an impact absorbing member, particularly an automobile impact absorbing member.

  Hereinafter, each requirement of the present invention will be described in detail.

1. Chemical composition The reasons for limiting each element are as follows. In the following description, “%” for the content means “% by mass”.

C: More than 0.2% and 0.55% or less C is a basic element that improves the strength of the steel material. C finely disperses carbides in bainite and contributes to increasing the strength of the ferrite phase. In addition, the bainite processing curability is improved. However, when the C content is 0.2% or less, it is difficult to obtain the effect by the above action. Therefore, C needs to be contained exceeding 0.2%. On the other hand, when the C content exceeds 0.55%, cementite is coarsened, and ductility and toughness are lowered. Therefore, the C content is 0.55% or less. The C content is desirably 0.25% or more, and more desirably 0.35% or more. Further, the C content is desirably 0.5% or less, and more desirably 0.45% or less.

Si: 0.1 to 0.3%
Si is an element having the effect of suppressing inclusions in the steel by the deoxidation effect and preventing impact cracking. However, if the Si content is less than 0.1%, it is difficult to obtain the above effects. Therefore, the Si content is 0.1% or more. On the other hand, if the Si content exceeds 0.3%, the prior austenite grain boundaries in the bainite structure are weakened, and impact cracks are more likely to occur. Therefore, the Si content is 0.1 to 0.3%. The Si content is desirably 0.15% or more, and desirably 0.25% or less.

Mn: 0.1 to 0.3%
Mn improves hardenability and facilitates the formation of a bainite phase. However, if the Mn content is less than 0.1%, it is difficult to obtain the above effects. Therefore, the Mn content is 0.1% or more. On the other hand, when the Mn content exceeds 0.3%, a martensite phase is easily generated, and ductility and toughness deteriorate. Therefore, the Mn content is 0.1 to 0.3%. The Mn content is desirably 0.15% or more, and desirably 0.25% or less.

Mo: 1-4%
Mo is an element that improves hardenability and facilitates the formation of a bainite phase by being compounded with B. Furthermore, it suppresses the motion of dislocations in the austenite grain boundaries and austenite, thereby contributing to refinement of the austenite grains. However, if the Mo content is less than 1%, it is difficult to obtain the above effects. Therefore, the Mo content is 1% or more. On the other hand, if the Mo content exceeds 4%, coarse carbides are generated at the prior austenite grain boundaries, and the local ductility and toughness are reduced. Therefore, the Mo content is 1 to 4%. The Mo content is desirably 1.5% or more, and desirably 3.5% or less.

Ti: 0.002 to 0.008%
Ti is an element that produces carbon / nitrides such as TiC and TiN, and has an effect of suppressing the coarsening of crystal grains by a pinning effect on the growth of crystal grains. Further, Ti has an effect of reducing the solute N. In the steel containing B as will be described later, if solid solution N is present, B and N are combined to form BN, so that the effect of containing B is not exhibited. Therefore, Ti is an effective element for indirectly exerting the effect of B. However, if the Ti content is less than 0.002%, it is difficult to obtain the above effects. Therefore, the Ti content is set to 0.002% or more. On the other hand, when the Ti content exceeds 0.008%, coarse TiN is generated, and local ductility and toughness are lowered. Therefore, the Ti content is set to 0.002 to 0.008%. The Ti content is desirably 0.003% or more, and desirably 0.006% or less.

Al: 0.01 to 0.2%
Al has the effect of suppressing inclusions in the steel due to the deoxidation effect and preventing impact cracking. However, if the Al content is less than 0.01%, it is difficult to obtain the above effects. On the other hand, if the Al content exceeds 0.2%, the oxides and nitrides are coarsened, and impact cracking is promoted. Therefore, the Al content is set to 0.01 to 0.2%. The Al content is desirably 0.02% or more, and desirably 0.1% or less.

B: 0.001 to 0.006%
B is an element that improves the hardenability and facilitates the formation of a bainite phase by being contained in a composite with Mo. However, if the B content is less than 0.001%, it is difficult to obtain the above effects. Therefore, the B content is 0.001% or more. On the other hand, if the B content exceeds 0.006%, coarse carbides are generated at the prior austenite grain boundaries, and the local ductility and toughness deteriorate. Therefore, the B content is 0.001 to 0.006%. The B content is desirably 0.003% or more and desirably 0.005% or less.

N: 0.001 to 0.015%
N is an element that has the effect of suppressing austenite and ferrite grain growth and impact cracking by forming nitrides. However, when the N content is less than 0.001%, it is difficult to obtain the above effects. On the other hand, when the N content exceeds 0.015%, the nitride becomes coarse, and on the contrary, impact cracking is promoted. Therefore, the N content is 0.001 to 0.015%. The N content is desirably 0.002% or more, and desirably 0.010% or less.

V: 0 to 0.3%
V, like Mo, is an element that improves the hardenability and facilitates the formation of the bainite phase by being contained in a composite with B. Furthermore, it suppresses the motion of dislocations in the austenite grain boundaries and austenite, thereby contributing to refinement of the austenite grains. Therefore, you may contain V as needed. However, if the V content exceeds 0.3%, coarse VC and VN are generated, and ductility and toughness are reduced. Therefore, the V content is 0 to 0.3%. The V content is desirably 0.25% or less. Moreover, when obtaining said effect, it is desirable to make V content 0.05% or more.

Nb: 0 to 0.05%
Nb, like Mo, is an element that improves hardenability and facilitates the formation of a bainite phase by being compounded with B. Furthermore, it suppresses the motion of dislocations in the austenite grain boundaries and austenite, thereby contributing to refinement of the austenite grains. Therefore, you may contain Nb as needed. However, if the Nb content exceeds 0.05%, coarse NbC and NbN are generated, and ductility and toughness are reduced. Therefore, the Nb content is 0 to 0.05%. The Nb content is preferably 0.04% or less. In order to obtain the above effect, the Nb content is preferably 0.01% or more, and more preferably 0.015% or more.

P: 0.02% or less P is contained as an impurity, weakens the grain boundary, and causes deterioration of hot workability. When the P content exceeds 0.02%, the deterioration of hot workability becomes significant. Therefore, the P content is 0.02% or less. The P content is desirably 0.015% or less. A lower P content is desirable, but excessive reduction leads to a significant cost increase. From the viewpoint of a realistic manufacturing process and manufacturing cost, the P content is preferably 0.001% or more.

S: 0.005% or less S is contained as an impurity, weakens the grain boundary, and causes deterioration of hot workability and ductility. When the S content exceeds 0.005%, the hot workability and the ductility deteriorate significantly. Therefore, the S content is 0.005% or less. The S content is desirably 0.004% or less. The smaller the S content, the better. However, excessive reduction leads to a significant cost increase. From the viewpoint of a realistic manufacturing process and manufacturing cost, the S content is preferably 0.0005% or more.

  The high-strength steel material of the present invention has a chemical composition comprising the above-described elements from C to S, the balance Fe and impurities.

  Here, “impurities” are components that are mixed due to various factors of raw materials such as ores and scraps and manufacturing processes when industrially manufacturing steel sheets, and are permitted within a range that does not adversely affect the present invention. Means something.

2. Metal structure: Lower bainite structure single phase The high strength steel according to the present invention has a lower bainite structure single phase metal structure in order to improve ductility and toughness while having a tensile strength exceeding 1400 MPa. . The lower bainite structure is a structure made of plate-like ferrite, in which cementite particles are generated in the ferrite plate. The lower bainite structure is characterized by the fact that cementite is aligned in the same direction, and can be easily distinguished from other structures (martensite, ferrite, pearlite, upper bainite, etc.) by observation using a scanning electron microscope or transmission electron microscope. can do.

  In the present invention, the “single phase of the lower bainite structure” allows a structure other than the lower bainite such as cementite, ferrite, and austenite to be included within a total range of less than 5%.

Former austenite grain size: 5 μm or less The lower bainite structure is in the low temperature region of the bainite formation temperature range, so carbon redistribution occurs in the structure due to bainite transformation, and the hard martensite structure in the high carbon region May generate. When hard martensite is generated in the lower bainite structure, local ductility and toughness may be lowered. However, it has been clarified that by reducing the prior austenite grain boundaries, the formation of martensite is suppressed and a homogeneous lower bainite structure can be obtained.

  When the prior austenite particle size exceeds 5 μm, the effect of suppressing the formation of martensite cannot be obtained. Therefore, the prior austenite grain size is 5 μm or less. The prior austenite particle size is desirably 3 μm or less. The prior austenite grain size is desirably as small as possible, but it is not practical to make it excessively fine, so it is desirable to make it 1 μm or more.

3. Manufacturing method Although there is no restriction | limiting in particular about the manufacturing method of the high strength steel materials which concern on this invention, For example, it can manufacture by giving the heat processing shown below with respect to steel which has said chemical composition.

3-1 Metal structure of steel before heat treatment The metal structure of steel to be subjected to heat treatment (hereinafter also referred to as “heat-treating steel”) is a single-phase structure selected from ferrite, pearlite, cementite, and tempered martensite. Or it is desirable to set it as the multiphase structure in which 2 or more types are mixed. The structure is preferably fine-grained. In the case of a composite structure of ferrite and pearlite or a composite structure of ferrite and cementite, the ferrite particle diameter is preferably controlled to 10 μm or less. In the case of tempered martensite, it is desirable that the prior austenite grain size be 10 μm or less. The crystal grain size is more preferably 5 μm or less, and further preferably 3 μm or less. The steel for heat treatment that satisfies the above structural requirements may be produced by any process of hot rolling, hot forging, and heat treatment.

3-2 Heat Treatment Step As described above, the high-strength steel material according to the present invention can be manufactured by subjecting the heat treatment steel to a heat treatment including the following steps. Each step in the heat treatment process will be described in detail below.

a) Heating step In the heat treatment step of the present invention, heating is performed to a temperature range of 800 to 900 ° C. In order to austenize the structure, the heating temperature is desirably 800 ° C. or higher. On the other hand, when the heating temperature exceeds 900 ° C., the crystal grains become coarse and adversely affect ductility and toughness. Therefore, the heating temperature is desirably in the range of 800 to 900 ° C. In addition, you may add hot processing after a heating as needed.

b) Cooling step After heating, cooling is performed to a temperature range of 350 to 450 ° C at a cooling rate of 10 ° C / s or more, and then cooling is stopped. When the cooling rate is less than 10 ° C./s, ferrite precipitates and there is a possibility that the lower bainite structure single phase is not obtained. On the other hand, if the cooling stop temperature is not in the range of 450 to 350 ° C., the transformation temperature range of bainite or martensite is not achieved, and ferrite is precipitated.

c) Holding step After cooling to the cooling stop temperature, the temperature is held for 10 to 100 s. Within this holding time, bainite or martensite transformation is performed. If the holding time is less than 10 s, the bainite transformation is not completed, and elongation, local ductility, and toughness may be deteriorated. On the other hand, if the holding time exceeds 100 s, the carbides may become coarse and adversely affect toughness.

d) Others In the heat treatment step of the present invention, after the heating, cooling, and holding steps, a tempering treatment within 1 h may be performed in a temperature range of 400 to 600 ° C.

3-3 Warm working step If the steel is warm-worked before it is heat-treated, the subsequent heat-treatment tends to make the austenite crystal grains finer, and the lower austenite grain size is finer lower bainite. It becomes easy to obtain a steel material having a metal structure composed of a structure. In the case where the warm working is performed before the heat treatment, it is desirable that the cross-section reduction rate is 20% or more in a temperature range of 500 to 750 ° C. When the warm working temperature exceeds 750 ° C. or when the cross-section reduction rate is less than 20%, the effect of warm working may not be exhibited. On the other hand, it is desirable that the warm working temperature is low, but in practice it is preferably 500 ° C. or higher, and more preferably 600 ° C. or higher. In addition, the higher the cross-sectional reduction rate, the higher is desirable, but the upper limit is desirably 50% in view of the performance of the equipment. The warm working can be performed by a general process such as plate rolling, hole rolling, warm pressing or the like.

3-4 Cold working step When cold working is performed on the above steel before the heat treatment, austenite crystal grains are likely to be refined in the subsequent heat treatment, and the lower austenite grain size of the lower bainite is fine. It becomes easy to obtain a steel material having a metal structure composed of a structure. In the case where the cold working is performed before the heat treatment, it is desirable to carry out under the condition that the cross-sectional reduction rate is 20% or more. The higher the cross-sectional reduction rate, the higher is desirable, but the upper limit is preferably 50% in view of the performance of the equipment. The cold working can be performed by a general process such as rolling, pressing, or bending.

  EXAMPLES Hereinafter, although an Example demonstrates this invention more concretely, this invention is not limited to these Examples.

  After casting steel having the chemical composition shown in Table 1 by vacuum melting 150 kg of molten steel, the steel is heated at a furnace temperature of 1250 ° C., hot forged at a temperature of 950 ° C. or higher, and further shown below. The steel for heat treatment was produced by processing.

  For test numbers 1 to 4, hot rolling was performed by 6-pass multi-pass rolling, and finish rolling was performed at 920 ° C. with a total rolling reduction of 43%, followed by air cooling to produce steel for heat treatment. In addition, for test numbers 5 to 8 and 30, hot rolling was performed by 6-pass multi-pass rolling, and finish rolling was performed at 780 ° C. with a total rolling reduction of 65%, followed by water cooling to produce a heat-treating steel.

  For test numbers 9 to 16, austenitizing heat treatment at 850 ° C. for 30 minutes, followed by water cooling, followed by tempering treatment at 680 ° C. for 2 hours to produce heat-treated steel. Among them, for test numbers 13 to 16, warm rolling was performed at 650 ° C. with a total rolling reduction of 40% after the tempering treatment.

  For test numbers 17 to 20, hot rolling was performed by 6-pass multi-pass rolling, and finish rolling with a total reduction ratio of 43% was added at 850 ° C., and then air-cooled to room temperature. Thereafter, cooling with a total reduction of 40% was performed. Hot rolling was performed to obtain a steel for heat treatment. And about test numbers 21-23, after hot-rolling by multipass rolling of 6 passes, after finishing rolling of 65% of total rolling reduction at 780 degreeC, it is air-cooled to room temperature, Then, total rolling reduction of 40 % Of the steel was cold-rolled to obtain a heat-treating steel.

  For test No. 24, hot rolling was performed by multi-pass rolling of 4 passes, and after finishing rolling at 920 ° C. with a total rolling reduction of 35%, air cooling was performed, and then cold rolling with a total rolling reduction of 25% was performed. To give steel for heat treatment. For test numbers 25 and 26, hot rolling was performed by 6-pass multi-pass rolling, finish rolling with a total reduction ratio of 53% was added at 830 ° C., water cooling was performed, and then the total reduction ratio was 30 at 650 ° C. % And 40% warm rolling were performed to obtain heat-treated steel.

  For test numbers 27 and 28, hot rolling was performed by multi-pass rolling of 6 passes, and finish rolling with a total rolling reduction of 65% was added at 780 ° C., followed by water cooling, followed by 40% warm at 650 ° C. Rolling was performed to obtain steel for heat treatment. For test number 29, austenitizing heat treatment at 850 ° C. for 30 minutes, followed by water cooling, followed by tempering treatment at 680 ° C. for 2 hours. Then, 40% warm rolling was performed at 650 ° C. to obtain a steel for heat treatment.

  Specimens were cut out from the above steel for heat treatment and then mirror polished. Thereafter, the steel was corroded with nital or picric acid, and was observed with respect to 5 visual fields of each sample using a scanning electron microscope (magnification 3000 times), and the metal structure was identified and the average particle diameter was measured. Here, the average particle size means the ASTM nominal particle size obtained in accordance with ASTM-E112. The ASTM nominal particle size can be calculated by multiplying the average particle section by 1.13.

  The plate | board thickness of the steel for heat processing obtained at said process is all 2 mm. Each steel for heat treatment was heat treated to produce a steel material. The heat treatment includes a heating step, a cooling step, and a holding step, and each processing condition is as shown in Table 2. Thereafter, the characteristics of each steel material were evaluated. Evaluation items of characteristics are tensile characteristics, toughness (Charpy impact characteristics) and bending characteristics. The characteristic evaluation method will be described below.

<Tensile properties>
By taking a JIS No. 13 tensile test piece and conducting a tensile test, the values of 0.2% yield strength (MPa), maximum tensile stress (MPa), yield ratio, elongation (%), and width drawing (%) were measured. .

<Toughness>
Toughness was evaluated by Charpy impact test. After laminating 5 pieces of the above-mentioned 2 mm steel plate samples and screwing them, the cross section was processed so as to have a 10 mm square to produce a laminate. Thereafter, a V notch having a depth of 2 mm and an angle of 45 ° was processed on the side surface of the laminate so that the crack cross section was parallel to the 10 mm square cross section to obtain a Charpy impact test piece. Impact characteristics were evaluated by impact energy at a temperature of 0 ° C.

<Bending characteristics>
Bending characteristics were evaluated based on the presence or absence of cracks after bending at bending radii of 5 mm and 7 mm. The presence or absence of cracks was determined using a stereomicroscope.

In the present invention, 0.2% proof stress: 1200 MPa or more, yield ratio: 0.85 or more, elongation: 10% or more, width drawing: 10% or more, Charpy value: 30 J / cm ² or more, bendability: R = 7 mm, Those satisfying all the conditions that there was no crack in 5 mm were judged to be “excellent in impact absorption characteristics”.

  As can be seen from Table 2, test numbers 1, 5, 13 and 28, which are comparative examples, do not provide the necessary tensile strength because the metal structure of the steel material is not a single phase of the lower bainite structure. The% yield strength and yield ratio were also poor. Test No. 2 resulted in slightly inferior Charpy impact properties and bending properties because the metal structure of the steel material was not a single phase of the lower bainite structure and the prior austenite grain size was also coarse. Test Nos. 3, 4, 7-9, 12, 16 and 23 have a steel structure having a martensite single phase and a coarse austenite grain size, so 0.2% proof stress, yield ratio, elongation, All Charpy impact properties and bending properties were inferior.

  In Test Nos. 10 and 24, although the metal structure was a single phase of the lower bainite structure, the prior austenite grain size was also coarse, resulting in poor Charpy impact properties and bending properties. Test Nos. 11 and 15 resulted in inferior yield ratio, elongation, Charpy impact property and bending property because the metal structure of the steel material was a martensite single phase and the prior austenite grain size was also coarse.

  In Test No. 25, the contents of Si and Mo are out of the specified range, and in Test No. 26, the contents of Mn and Mo are out of the specified range. Since it was not a phase, the yield ratio, elongation, Charpy impact property and bending property were inferior. In Test Nos. 29 and 20, since the B content is out of the specified range, and the metal structure of the steel material is not a single phase of the lower bainite structure, the necessary tensile strength cannot be obtained, 0.2% proof stress and The yield ratio was also bad.

  On the other hand, Test Nos. 6, 14, 17-22 and 27, which are examples of the present invention, have a tensile strength exceeding 1400 MPa, and all of 0.2% proof stress, yield ratio, elongation, Charpy impact property and bending property. It is clear that it has good shock absorption characteristics.

  According to the present invention, since the occurrence of cracking of the impact absorbing member when an impact load is applied can be suppressed, the impact absorbing energy of the impact absorbing member can be dramatically increased. Therefore, the high-strength steel material according to the present invention is suitable for use as a material for an impact absorbing member, particularly an automobile impact absorbing member.

Claims (5)

Chemical composition is mass%,
C: more than 0.2% and 0.55% or less,
Si: 0.1 to 0.3%,
Mn: 0.1 to 0.3%,
Mo: 1-4%
Ti: 0.002 to 0.008%,
Al: 0.01-0.2%
B: 0.001 to 0.006%,
N: 0.001 to 0.015%,
V: 0 to 0.3%
Nb: 0 to 0.05%,
Balance: Fe and impurities,
The metal structure is a single phase of a lower bainite structure having a prior austenite particle size of 5 μm or less,
A high-strength steel material having a tensile strength exceeding 1400 MPa. The chemical composition is mass%,
V: 0.05-0.3%
Nb: 0.01-0.05%
The high-strength steel material of Claim 1 containing 1 or more types selected from these. A method for producing the high-strength steel material according to claim 1 or 2,
The steel having the chemical composition according to claim 1 or 2 is heated to a temperature range of 800 to 900 ° C, and then cooled to a temperature range of 350 to 450 ° C at a cooling rate of 10 ° C / s or more. A method for producing a high-strength steel material, comprising a heat treatment step of holding at a temperature for 10 to 100 seconds and then cooling. A method for producing the high-strength steel material according to claim 1 or 2,
The steel having the chemical composition according to claim 1 or claim 2 is subjected to warm working so that the cross-sectional reduction rate is 20% or more in a temperature range of 500 to 750 ° C, and then a temperature of 800 to 900 ° C. High-strength steel material comprising a heat treatment step of heating to a temperature range, then cooling to a temperature range of 350 to 450 ° C. at a cooling rate of 10 ° C./s or more, holding at that temperature for 10 to 100 seconds, and then cooling. Method. A method for producing the high-strength steel material according to claim 1 or 2,
The steel having the chemical composition according to claim 1 or claim 2 is subjected to cold working so that the cross-sectional reduction rate is 20% or more, and then heated to a temperature range of 800 to 900 ° C. A method for producing a high-strength steel material, comprising a heat treatment step of cooling to a temperature range of 350 to 450 ° C at a cooling rate of ° C / s or more, holding the temperature for 10 to 100 seconds, and then cooling.