High strength thin steel sheet having high hydrogen embrittlement resisting property and high workability

Abstract

The present invention provides a high strength thin steel sheet that has high hydrogen embrittlement resisting property and high workability. The high strength thin steel sheet having high hydrogen embrittlement resisting property has a metallurgical structure after stretch forming process to elongate 3% comprises: (i) 1% or more residual austenite; 80% or more in total of bainitic ferrite and martensite; and 9% or less (may be 0%) in total of ferrite and pearlite in terms of proportion of area to the entire structure, wherein the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher, or (ii) 1% or more residual austenite in terms of proportion of area to the entire structure; mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher; mean length of minor axes of the residual austenite grains is 1 μm or less; minimum distance between the residual austenite grains is 1 μm or less; and the steel has tensile strength of 1180 MPa or higher.

Description

BACKGROUND OF THE INVENTION

1. Field of the Invention

The present invention relates to a high strength thin steel sheet that has high hydrogen embrittlement resisting property (particularly the hydrogen embrittlement resisting property after being subjected to forming process) and high workability, especially to a high strength thin steel sheet that has high resistance against fractures due to hydrogen embrittlement such as season crack and delayed fracture that pose serious problems for steel sheets having tensile strength of 1180 MPa or higher, and has high workability.

2. Description of the Related Art

There are increasing demands for the steel sheet, that is pressed or bent into a form of a high-strength component of automobile or industrial machine, to have both high strength and high ductility at the same time. In recent years, there are increasing needs for high strength steel sheets having strength of 1180 MPa or higher, as the automobiles are being designed with less weight. A type of steel sheet that is regarded as promising to satisfy these needs is TRIP (transformation induced plasticity) steel sheet.

The TRIP steel sheet includes residual austenite structure and, when processed to deform, undergoes considerable elongation due to induced transformation of the residual austenite (residual γ) into martensite by the action of stress. Known examples of the TRIP steel include TRIP type composite-structure steel (TPF steel) that consists of polygonal ferrite as the matrix phase and residual austenite; TRIP type tempered martensite steel (TAM steel) that consists of tempered martensite as the matrix phase and residual austenite; and TRIP type bainitic steel (TBF steel) that consists of bainitic ferrite as the matrix phase and residual austenite. Among these, the TBF steel has long been known (described, for example, in NISSIN STEEL TECHNICAL REPORT, No. 43, December 1980, pp 1-10), and has such advantages as the capability to readily provide high strength due to the hard bainitic ferrite structure, and the capability to show outstanding elongation because fine residual austenite grains can be easily formed in the boundary of lath-shaped bainitic ferrite in the bainitic ferrite structure. The TBF steel also has such an advantage related to manufacturing, that it can be easily manufactured by a single heat treatment process (continuous annealing process or plating process).

In the realm of high strength of 1180 MPa upward, however, the TRIP steel sheet is known to suffer a newly emerging problem of delayed fracture caused by hydrogen embrittlement, similarly to the conventional high strength steel. Delayed fracture refers to the failure of high-strength steel under stress, that occurs as hydrogen originating in corrosive environment or the atmosphere infiltrates and diffuses in microstructural defects such as dislocation, void and grain boundary, and makes the steel brittle. This results in decreases in ductility and toughness of the metallic material.

It has been well known that the high strength steel that is widely used in the manufacture of PC steel wire and line pipe experiences hydrogen embrittlement (pickling embrittlement, plating embrittlement, delayed fracture, etc.) caused by the infiltration of hydrogen into the steel when tensile strength of the steel becomes 980 MPa or higher. Accordingly, most of technologies of improving hydrogen embrittlement resisting property have been developed aiming at steel members such as bolt. “New Development in Elucidation of Delayed Fracture” (published by The Iron and Steel Institute of Japan in January, 1997), for example, describes that it is effective in improving the resistance against delayed fracture to add element such as Cr, Mo or V that demonstrates resistance against temper softening to the metal structure that is based on tempered martensite as the major phase. This technology is intended to cause the delayed fracture to take place within grains instead of in the grain boundaries, thereby to constrain the fracture from occurring, by precipitating alloy carbide and making use thereof as the site for trapping hydrogen.

Thin steel sheets having strength higher than 780 MPa have rarely been used for the reason of workability and weldability. Also hydrogen embrittlement has rarely been regarded as a problem for thin steel sheets where hydrogen that has infiltrated therein is immediately released due to the small thickness. For these reasons, much efforts have not been dedicated to counter the hydrogen embrittlement. In recent years, however, higher strength is required of the reinforcement members such as bumper, impact beam and seat rail, etc., in order to meet the requirement of weight reduction of the automobile and to improve the collision safety. Automobile components that are shaped by pressing or bending process such as pillar are also required to have higher strength. As a result, there have been increasing demands for high strength steel sheet having strength of 980 MPa or higher for the manufacture of these parts. This makes it necessary to improve hydrogen embrittlement resisting property of the high strength steel sheet.

Use of the technology addressed to the bolt steel described above may be considered for improving the hydrogen embrittlement resisting property of the high strength steel sheet. However, in the case of “New Development in Elucidation of Delayed Fracture” (published by The Iron and Steel Institute of Japan in January, 1997), for example, 0.4% or higher of C content and much alloy elements are contained, and therefore application of this technology to a thin steel sheet compromises the workability required of the thin steel sheet. The technology also has a drawback related to the manufacturing process, since it takes several hours or longer period of heat treatment to cause the alloy carbide to precipitate. Therefore, improvement of the hydrogen embrittlement resisting property of a thin steel sheet requires it to develop a novel technology.

It is relatively easy to achieve a high strength with quench-hardened (tempered) martensite steel that has been commonly used as a high-strength steel. However, improvement of the workability without variability essentially requires it to provide a tempering process which makes it necessary to strictly control the temperature and duration of the process. This also sometimes increases the possibility of tempering embrittlement and makes it difficult to reliably improve workability. Although there is a steel of composite structure of martensite and ferrite or the like developed to improve ductility, such a steel has a high notch sensitivity due to mixed presence of hard phase and soft phase, thus making it difficult to achieve sufficient improvement of hydrogen embrittlement resisting property.

Hydrogen-induced delayed fracture is believed to occur in such a steel that contains martensite, because hydrogen is concentrated in grain boundaries of prior austenite thereby to form voids or other defects that become the starting points of the fracture. Common practice that has been employed to decrease the sensitivity to delayed fracture is to diffuse fine grains of carbide or the like uniformly as the site for trapping hydrogen, thereby to decrease the concentration of diffusive hydrogen. However, even when a large number of carbide grains or the like are diffused as the trap site for hydrogen, there is a limitation to the hydrogen trapping capability and delayed fracture attributable to hydrogen cannot be fully suppressed.

Japanese Unexamined Patent Publication (Kokai) No. 11-293383 describes a technology to improve the hydrogen embrittlement resisting property of steel sheet, where hydrogen-induced defects can be suppressed by having oxides that include Ti and Mg exist as the main components in the structure. However, this technology is intended for thick steel sheets and, although consideration is given to delayed fracture after welding with a large input heat, no consideration is given to the environment (for example, corrosive environment, etc.) in which automobile parts manufactured by using thin steel sheets are used.

Japanese Unexamined Patent Publication (Kokai) No. 2003-166035 describes that it is made possible to improve the ductility and delayed fracture resistance after being subjected to forming process, by controlling the mutual relationships between 1) the form (standard deviation and mean grain size) in which oxide, sulfide, composite crystallization product or composite precipitate of Mg is dispersed, 2) volumetric proportion of residual austenite and 3) strength of the steel sheet. However, it is difficult to improve the hydrogen embrittlement resisting property in such an environment as hydrogen is generated through corrosion of the steel sheet simply through the trapping effect achieved by controlling the form of precipitate.

It has been a common practice in the past to reduce the residual austenite that was believed to have an adverse effect on the hydrogen embrittlement resisting property. In recent years, however, the effect of residual austenite on the improvement of hydrogen embrittlement resisting property has been recognized and accordingly much attention has been paid to the TRIP steel that contains residual austenite.

Tomohiko HOJO et. al “Hydrogen Embrittlement of High Strength Low Alloy TRIP Steel (Part 1: Hydrogen Absorbing Characteristic and Ductility”, The Society of Materials Science, Japan, proceedings of 51^st academic lecture meeting, 2002, vol. 8, pp 17-18 and Tomohiko HOJO et. al “Influence of Austempering Temperature on Hydrogen Embrittlement of High, for example, describe investigations into the hydrogen embrittlement resisting property of the TRIP steel. It is pointed out that, among the TRIP steels, TBF steel has particularly high hydrogen absorbing capacity, and observation of a fracture surface of the TBF steel shows the restriction of quasi cleavage fracture due to storage of hydrogen. However, the TBF steels reported in the documents described above show delayed fracture characteristic of about 1000 seconds at the most in terms of the time before crack occurrence measured in cathode charging test, indicating that these steels are not meant to endure the harsh operating environment such as that of automobile parts over a long period of time. Moreover, since the heat treatment conditions reported in the documents described above involve heating temperature being set higher, there are such problems as low efficiency of practical manufacturing process. Thus it is strongly required to develop a new species of TBF steel that provides high production efficiency as well. Also there has been such a problem that press forming operation leads to lower hydrogen embrittlement resisting property.

As described above, there have been virtually no TRIP steels containing residual austenite that have been developed so as to demonstrate high workability when processed to form parts, by taking measures to counter hydrogen embrittlement after the forming process in consideration of the harsh operating environment such as that of automobile parts over a long period of time.

SUMMARY OF THE INVENTION

The present invention has been made with the background described above, and has an object of providing a high strength thin steel sheet that shows high hydrogen embrittlement resisting property in a harsh operating environment over a long period of time after the process of forming the steel sheet into a part, and has improved workability and tensile strength of 1180 MPa or higher.

In order to achieve the object described above, the present inventors conducted a research on a steel sheet that shows high hydrogen embrittlement resisting property after the forming process, and demonstrates improved workability which is the characteristic property of the TRIP steel sheet during the forming process. Through the research, it was found that it is very important to control the metallurgical structure after the forming process in order to achieve high hydrogen embrittlement resisting property after the forming process. Specifically, it was found that it is important that the metal structure after the stretch forming process is constituted from:

1% or more of residual austenite;

80% or more in total of bainitic ferrite and martensite; and

9% or less (may be 0%) in total of ferrite and pearlite in terms of the proportion of area to the entire structure, wherein the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher.

A first high strength thin steel sheet having high hydrogen embrittlement resisting property according to the present invention comprises higher than 0.25 and up to 0.60% of C (contents of components given in terms of percentage in this patent application all refer to percentage by weight), 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less P, 0.02% or less S and 1.5% or less (higher than 0%) of Al, while iron and inevitable impurities making up the rest, wherein the metallurgical structure comprises:

1% or more residual austenite;

80% or more in total of bainitic ferrite and martensite; and

9% or less (may be 0%) in total of ferrite and pearlite in the proportion of area to the entire structure, and wherein the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher, and the steel has tensile strength of 1180 MPa or higher.

The present inventors also conducted another research from a point of view that was different from that of the former research, and found that high hydrogen embrittlement resisting property after the forming process can be achieved by controlling the metal structure after the forming process as follows. It is important that the metal structure after the forming process comprises:

1% or more residual austenite;

the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher.

mean length of minor axes of the residual austenite grains is 1 μm or less; and minimum distance between the residual austenite grains is 1 μm or less.

When the metal structure is controlled as described above, hydrogen embrittlement resisting property of the high strength thin steel sheet can be sufficiently improved without adding much alloy elements. The phrase “after the forming process” means the state of the steel sheet after being stretched with an elongation ratio of 3%. Specifically, the steel sheet is subjected to uniaxial stretching of 3% at the room temperature (the stretching process of 3% elongation may hereinafter be referred to simply as “processing”).

A second high strength thin steel sheet having high hydrogen embrittlement resisting property according to the present invention comprises higher than 0.25 and up to 0.60% of C, 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less P, 0.02% or less S, 0.5% or less (higher than 0%) Al, while iron and inevitable impurities making up the rest, wherein the metal structure after the stretch forming process of 3% elongation comprises:

1% or more residual austenite;

the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher;

mean length of minor axes of the residual austenite grains is 1 μm or less;

minimum distance between the residual austenite grains is 1 μm or less; and

tensile strength is 1180 MPa or higher.

According to the present invention, it is made possible to manufacture, with a high level of productivity, a high strength thin steel sheet having tensile strength of 1180 MPa or higher that neutralizes hydrogen that infiltrates from the outside after the steel sheet has been formed into a part thereby to maintain satisfactory hydrogen embrittlement resisting property, and demonstrates high workability during the forming process. Use of the high strength thin steel sheet makes it possible to manufacture high strength parts that hardly experience delayed fracture, such as bumper, impact beam and other reinforcement members and other automobile parts such as seat rail, pillar, etc.

BRIEF DESCRIPTION OF THE DRAWINGS

FIG. 1 is a schematic perspective view of a part used in pressure collapse test in Example 1.

FIG. 2 is a side view schematically showing the setup of pressure collapse test in Example 1.

FIG. 3 is a schematic perspective view of a part used in impact resistance test in Example 1.

FIG. 4 is a sectional view along A-A in FIG. 3.

FIG. 5 is a side view schematically showing the setup of impact resistance test in Example 1.

FIG. 6 is a photograph of TEM observation (magnification factor 15000) of No. 101 (inventive steel) of Example 1.

FIG. 7 is a photograph of TEM observation (magnification factor 15000) of No. 120 (comparative steel) of Example 1.

FIG. 8 is a photograph of TEM observation (magnification factor 15000) of No. 201 (inventive steel) of Example 2.

FIG. 9 is a photograph of TEM observation (magnification factor 15000) of No. 220 (comparative steel) of Example 2.

FIG. 10 is a graph showing the relationship between the mean axis ratio of the residual austenite grains and hydrogen embrittlement risk index.

FIG. 11 is a diagram schematically showing the minimum distance between residual austenite grains.

FIG. 12 is a photograph of TEM observation (magnification factor 15000) of No. 301 (inventive steel) of Example 3.

FIG. 13 is a photograph of TEM observation (magnification factor 60000) of No. 301 (inventive steel) of Example 3.

FIG. 14 is a photograph of TEM observation (magnification factor 15000) of No. 313 (comparative steel) of Example 3.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

First Embodiment

The first high strength thin steel sheet according to the present invention is constituted from higher than 0.25 and up to 0.60% of C (contents of components given in terms of percentage in this patent application all refer to percentage by weight), 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less P, 0.02% or less S, 1.5% or less (higher than 0%) of Al, 1.0% or less (higher than 0%) of Mo and 0.1% or less (higher than 0%) of Nb, while iron and inevitable impurities making up the rest, and is characterized in that:

(i) the metal structure after the forming process contains: 1% or more residual austenite;

80% or more in total of bainitic ferrite and martensite; and

9% or less (may be 0%) in total of ferrite and pearlite in terms of the proportion of area to the entire structure, and

the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher; and

(ii) the steel contains a specified amount of Mo and/or Nb.

The requirements described above have reasons as follows.

(Metal Structure after Stretch Forming by 3% Elongation)

Metal structure after stretch forming process by 3% elongation was specified because, in various experiments conducted for the actual processing conditions in the manufacture of a part, best correlation between the result of laboratory test and the actual occurrence of cracks in the part was observed when the part was processed by stretch forming with an elongation ratio of 3%.

The phrase “after the forming process” means the state of the steel sheet after being stretch formed with elongation of 3%. Specifically, the steel sheet is subjected to elongation of 3% by uniaxial stretching at the room temperature (the stretch forming process of 3% elongation may hereinafter be referred to simply as “process”).

(1% or More Residual Austenite in the Area Proportion to the Entire Structure)

It is necessary that the metal contains 1% or more residual austenite in the area proportion to the entire structure after the process of forming the part, in order to achieve high hydrogen embrittlement resisting property in harsh operating environment over an extended period of time after forming the part. Content of the residual austenite is preferably 2% or higher, and more preferably 3% or higher. Since the desired level of high strength cannot be obtained when an excessive amount of residual austenite is contained after processing, it is recommended to set an upper limit of 20% (more preferably 15%) to the residual austenite content.

(Mean Axis Ratio (Major Axis/Minor Axis) of the Residual Austenite Grains: 5 or Higher)

Lath-shaped grains of residual austenite after the process have far higher capacity of trapping hydrogen than carbide. When the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher, in particular, it was found that hydrogen that infiltrates from the outside through atmospheric corrosion can be substantially neutralized thereby to achieve remarkable achievement in hydrogen embrittlement resisting property. The mean axis ratio of the residual austenite grains is preferably 10 or higher, and more preferably 15 or higher.

The residual austenite refers to a region that is observed as FCC (face centered cubic lattice) by the FE-SEM/EBSP method which will be described later. Measurement by the EBSP may be done, for example, by measuring a measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness at measuring intervals of 0.1 μm. The measuring surface is prepared by electrolytic polishing in order to prevent the residual austenite from transforming. Then the test piece is set in the lens barrel of an FE-SEM equipped with an EBSP detector (of which details will be described later) and is irradiated with electron beam. An EBSP image projected onto a screen is captured by a high sensitivity camera (VE-1000-SIT manufactured by Dage-MTI Inc.) and is sent to a computer. The computer carries out image analysis and generates color mapping of the FCC phase through comparison with a structural pattern simulated with a known crystal system (FCC (face centered cubic lattice) phase in the case of residual austenite). Area proportion of the region that is mapped as described above is taken as the area proportion of the residual austenite. This analysis was carried out by means of hardware and software of OIM (Orientation Imaging Microscopy™) system of TexSEM Laboratories Inc.

The mean axis ratio was determined by measuring the major axis and minor axis of residual austenite crystal grains existing in each of three arbitrarily chosen fields of view in the observation by means of TEM (transmission electron microscope) with magnification factor of 15000, and averaging the ratios of major axis to minor axis.

(80% or More in Total of Bainitic Ferrite and Martensite)

In order to decrease the number of intergranular fracture initiating points in the steel thereby to surely decrease the concentration of diffusive hydrogen to a harmless level and achieve a high strength, it is desirable to form the matrix phase of the steel structure after processing from a binary phase structure of bainitic ferrite and martensite with the bainitic ferrite acting as the main phase, instead of the single phase structure of martensite that is generally used for high strength steels.

In the single phase structure of martensite, a carbide (for example, film-like cementite) is likely to precipitate in the grain boundaries, thus making intergranular fracture likely to occur. In the case of the binary structure of bainitic ferrite and martensite with the bainitic ferrite acting as the main phase, in contrast, the bainitic ferrite is a hard phase and therefore it is easy to increase the strength of the entire structure as in the case of the single phase of martensite. The hydrogen embrittlement resisting property can also be improved as much hydrogen is trapped in the dislocations. It also has such an advantage that coexistence of the bainitic ferrite and the residual austenite which will be described later prevents the generation of carbide that acts as the intergranular fracture initiating points, and it becomes easier to create the lath-shaped residual austenite in the boundaries of lath-shaped bainitic ferrite.

Accordingly, it is required in the present invention that the binary structure of bainitic ferrite and martensite occupy 80% or more, preferably 85% or more and more preferably 90% or more of the entire structure after the stretch forming processing to elongate by 3%. Upper limit of the proportion may be determined by the balance with other structure (residual austenite), and is set to 99% when the other structures (ferrite, etc.) than the residual austenite is not contained.

The bainitic ferrite referred to in the present invention is plate-shaped ferrite having a lower structure of high density of dislocations. It is clearly distinguished from polygonal ferrite that has lower structure including no or very low density of dislocations, by SEM observation as follows.

Area proportion of bainitic ferrite structure is determined as follows. A test piece is etched with Nital etchant. A measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness is observed with SEM (scanning electron microscope) (magnification factor of 1500) thereby to determine the area proportion.

Bainitic ferrite is shown with dark gray color in SEM photograph (bainitic ferrite, residual austenite and martensite may not be distinguishable in the case of SEM observation), while polygonal ferrite is shown black in SEM photograph and has polygonal shape that does not include residual austenite and martensite inside thereof.

The SEM used in the present invention is a high-resolution FE-SEM (Field Emission type Scanning Electron Microscope XL30S-FEG manufactured by Philips Inc.) equipped with an EBSP (Electron Back Scattering Pattern) detector, that has a merit of being capable of analyzing the area observed by the SEM at the same time by means of the EBSP detector. EBSP detection is carried out as follows. When the sample surface is irradiated with electron beam, the EBSP detector analyzes the Kikuchi pattern obtained from the reflected electrons, thereby to determine the crystal orientation at the point where the electron beam has hit upon. Distribution of orientations over the sample surface can be measured by scanning the electron beam two-dimensionally over the sample surface while measuring the crystal orientation at predetermined intervals. The EBSP detection method has such an advantage that different structures that are regarded as the same structure in the ordinary microscopic observation but have different crystal orientations can be distinguished by the difference in color tone.

(9% or Less (may be 0%) in Total of Ferrite and Pearlite)

The steel sheet after the processing may be constituted either from only the structures described above (namely, a mixed structure of bainitic ferrite+martensite and residual austenite), or may include other structure such as ferrite (the term ferrite used herein refers to polygonal ferrite, that is a ferrite structure that includes no or very few dislocations) or pearlite to such an extent that the effect of the present invention is not compromised. Such additional components are structures that can inevitably remain in the manufacturing process of the present invention, of which concentration is preferably as low as possible, within 9%, preferably less than 5% and more preferably less than 3% according to the present invention.

In order to maintain high hydrogen embrittlement resisting property after the forming process, for example, large content of residual austenite of 5% or more may be contained in the steel sheet prior to the forming process, or large amount of fine residual austenite grains may be dispersed in the structure. Alternatively, forming process conditions may be controlled so as to make the residual austenite less likely to transform (for example, form the part by bending operation or control the forming temperature and/or stretching speed). The most desirable means of improving the workability and hydrogen embrittlement resisting property at the same time while maintaining the content of residual austenite before and after the processing substantially constant within an appropriate range and maintaining other properties (high strength, etc.) is to satisfy the following requirements (A) and (B).

(A) Increase C content in the composition and increase the concentration of C in the residual austenite.

Although residual austenite transforms into martensite when the steel sheet is deformed (processed), high content of C in the residual austenite stabilizes it so that further transformation becomes unlikely to occur. Thus residual austenite can be retained after the forming process, thereby maintaining the high hydrogen embrittlement resisting property.

According to the present invention, higher than 0.25% of C is contained in order to achieve the effects described above. C is also an element required to achieve a high strength of 1180 MPa or higher, and 0.27% or more, preferably 0.30% or more C is contained. However, in order to ensure corrosion resistance, concentration of C is limited within 0.6%, preferably 0.55% or lower and more preferably 0.50% or lower in the present invention.

It is recommended to increase the C content in the steel sheet as described above, thereby to maintain the concentration of C in the residual austenite (CγR) of 0.8% or higher. Controlling the value of CγR to 0.8% or higher enables it to effectively improve the elongation property, which is preferably 1.0% or higher and more preferably 1.2% or higher. While it is preferable that CγR is as high as possible, it is considered that in practice there is an upper limit of around 1.6%.

(B) Form the residual austenite in fine lath-shaped grains.

Residual austenite formed in fine lath-shaped grains does not undergo excessive transformation during the forming process, thus enabling it to maintain the residual austenite.

Some of the TRIP steels of the prior art have unsatisfactory hydrogen embrittlement resisting property despite sufficient content of residual austenite. The reason may be that, since residual austenite existing in the TRIP steel of the prior art generally has block shape of size on micrometer order, it can easily transform into martensite when being stressed and may act as the starting point of mechanical destruction. Through a research conducted by the present inventors, it was found that residual austenite formed in lath shape is more stable and less likely to transform into martensite than the residual austenite of the prior art that has block shape, given the same amount of deformation. This difference may be caused by the difference in the way in which the stress is applied and in the difference in spatial restriction, although not fully elucidated. Stabilization of residual austenite during processing has no influence on the lowering of workability of TRIP steel sheet due to induced transformation. According to the present invention, induced transformation proceeds efficiently and high workability can be achieved without hardly reducing the residual austenite, when the residual austenite is formed into fine lath shape as described above.

Lath-shaped grains of residual austenite having mean axis ratio (major axis/minor axis) of 5 or higher (preferably 10 or higher, and more preferably 15 or higher) minimizes the decrease of residual austenite during processing and makes it possible to easily achieve mean axis ratio (major axis/minor axis) of 5 or higher after processing, put the hydrogen absorbing capability of the residual austenite into full play and greatly improve hydrogen embrittlement resisting property. While no upper limit of the mean axis ratio is specified for the consideration of improvement in hydrogen embrittlement resisting property, the residual austenite grains are required to have certain level of thickness in order to achieve the TRIP effect during processing. Thus it is preferable to set an upper limit to 30, more preferably to 20 or less.

According to a preferred embodiment of the present invention, Mo and Nb are added for the purpose of reducing the size of the residual austenite grains. Mo has the effects of strengthening the grain boundary so as to suppress hydrogen embrittlement from occurring, in addition to reducing the size of the residual austenite grains. Mo also has the effect of improving the hardenability of the steel sheet. It is recommended to add 0.005% or more of Mo in order to achieve these effects. More preferably 0.1% or more of Mo is added. However, since the effects described above reach saturation when the Mo content exceeds 1.0%, resulting in economical disadvantage, Mo content is limited to 0.8% or less and more preferably to 0.5% or less.

Nb, in cooperation with Mo, acts very effectively to decrease the grain size of the structure. Nb also has the effect of increasing the strength of the steel sheet. It is recommended to add 0.005% or more of Nb in order to achieve these effects. More preferably 0.01% or more of Nb is added. However, since the effects described above reach saturation when an excessive Nb content is included, resulting in economical disadvantage, Nb content is limited to 0.1% or less and more preferably to 0.08% or less.

In order to readily obtain the structure described above after processing, it is recommended to make the steel sheet constituted from 80% or more (preferably 85% or more, and more preferably 90% or more) in total of bainitic ferrite and martensite, and 9% or less (preferably less than 5%, and more preferably less than 3% containing 0%) in total of ferrite and pearlite making up the rest of the residual austenite before processing. This is because it is preferable that the steel sheet has high hydrogen embrittlement resisting property prior to the processing as well as after the processing, and this constitution makes it easier to achieve the specified strength.

While this embodiment is characterized in that metal structure is controlled after processing, it is necessary to control the other components as described below, in order to form the metal structure and efficiently improve hydrogen embrittlement resisting property and strength thereby to ensure ductility required for the thin steel sheet.

<Si: 1.0 to 3.0%>

Si is an important element that effectively suppresses the residual austenite from decomposing and carbide from being generated, and is also effective in enhancing substitution solid solution for hardening the material. In order to make full use of these effects, it is necessary to include Si in a concentration of 1.0% or higher, preferably 1.2% or higher and more preferably 1.5% or higher. However, excessively high content of Si leads to conspicuous formation of scales due to hot rolling and makes it necessary to remove flaws, thus adding up to the manufacturing cost and resulting in economical disadvantage. Therefore Si content is controlled within 3.0%, preferably within 2.5% and more preferably within 2.0%.

<Mn: 1.0 to 3.5%>

Mn is an element required to stabilize austenite and obtain desired residual austenite. In order to make full use of this effect, it is necessary to add Mn in concentration of 1.0% or higher, preferably 1.2% or higher, and more preferably 1.5% or higher. However, adding an excessive amount Mn leads to conspicuous segregation and poor workability. Therefore upper limit to the concentration of Mn is set to 3.5% and more preferably to 3.0% or less.

<P: 0.15% or Lower (Higher than 0%)>

P intensifies intergranular fracture due to intergranular segregation, and the content thereof is therefore preferably as low as possible. Upper limit to the concentration of P is set to 0.15%, preferably 0.1% or less and more preferably to 0.05% or less.

<S: 0.02% or Lower (Higher than 0%)>

S intensifies the absorption of hydrogen into the steel sheet in corrosive environment, and the content thereof is therefore preferably as low as possible. Upper limit to the concentration of S is set to 0.02%.

<Al: 1.5% or Less (Higher than 0%)> (In the Case of Inventive Steel 1)

<Al: 0.5% or Less (Higher than 0%)> (In the Case of Inventive Steel 2)

0.01% or higher content of Al may be included for the purpose of deoxidation. In addition to deoxidation, Al also has the effects of improving the corrosion resistance and improving hydrogen embrittlement resisting property.

The mechanism of improving the corrosion resistance is supposedly based on the improvement of corrosion resistance of the matrix phase per se and the effect of formation rust generated by atmospheric corrosion, while the effect of the formation rust presumably has greater contribution. This is supposedly because the formation rust is denser and better in protective capability than ordinary iron rust, and therefore depresses the progress of atmospheric corrosion so as to decrease the amount of hydrogen generated by the atmospheric corrosion, thereby to effectively suppress the occurrence of hydrogen embrittlement, and hence the delayed fracture.

While details of the mechanism of improvement of the hydrogen embrittlement resistance by Al is not known, it is supposed that condensing of Al on the surface of the steel makes it difficult for hydrogen to infiltrate into the steel, and the decreasing diffusion rate of hydrogen in the steel makes it difficult for hydrogen to migrate so that hydrogen embrittlement becomes less likely to occur. In addition, stability of lath-shaped residual austenite improved by the addition of Al is believed to contribute to the improvement of hydrogen embrittlement resisting property.

In order to effectively achieve the effects of Al in improving the corrosion resistance and improving the hydrogen embrittlement resisting property, Al content is controlled to 0.2% or higher, preferably 0.5% or higher.

However, Al content must be controlled within 1.5% in order to prevent inclusions such as alumina from increasing in number and size so as to ensure satisfactory workability, ensure the generation of fine residual austenite grains, suppress corrosion from proceeding from the inclusion containing Al as the starting point, and prevent the manufacturing cost from increasing. In view of the manufacturing process, it is preferable to control so that A3 point is not higher than 1000° C.

As the Al content increases, inclusions such as alumina increase and workability becomes poorer. In order to suppress the generation of the inclusions such as alumina and make a steel sheet having higher workability, Al content is restricted within 0.5%, preferably within 0.3% and more preferably within 0.1%.

While constituent elements (C, Si, Mn, P, S, Al, Mo, Nb) of the steel of this embodiment is as described above with the rest substantially being Fe, it may include inevitable impurities introduced into the steel depending on the stock material, production material, manufacturing facility and other circumstances, containing 0.001% or less of N (nitrogen) In addition, other elements as described below may be intentionally added to such an extent that does not adversely affect the effects of the present invention.

<B: 0.0002 to 0.01%>

B is effective in increasing the strength of the steel sheet, and it is preferable that 0.0002% or more (more preferably 0.0005% or more) B is contained. However, an excessive content of B leads to poor hot processing property. Therefore, it is preferable to control the concentration of B to within 0.01% (more preferably within 0.005%).

<At Least One Selected from Among Ca: 0.0005% to 0.005%, Mg: 0.0005% to 0.01% and REM: 0.0005% to 0.01%)

Ca, Mg and REM (rare earth element) are effective in suppressing an increase in hydrogen ion concentration, that is, a decrease in pH in the atmosphere of the interface due to corrosion of the steel sheet surface, thereby to improve the corrosion resistance of the steel sheet. These elements are also effective in controlling the form of sulfide contained in the steel and improve the workability of the steel. In order to achieve the effects described above, it is recommended to add each of Ca, Mg and REM in concentration of 0.0005% or higher. However, since excessive contents of these elements leads to poor workability, it is preferable to keep the concentration of Ca to 0.005% or less and concentration of Mg and REM each within 0.01%.

Second Embodiment

The second high strength thin steel sheet according to the present invention is constituted from higher than 0.25% and up to 0.60% of C (contents of components given in terms of percentage in this patent application all refer to percentage by weight), 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less of P, 0.02% or less of S, 1.5% or less (higher than 0%) of Al while iron and inevitable impurities constitute the rest, wherein: (i) the structure after the forming process comprises:

1% or more residual austenite;

mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher;

80% or more in total of bainitic ferrite and martensite; and

9% or less (may be 0%) in total of ferrite and pearlite in the proportion of area to the entire structure, and

(ii) the steel contains specified amount of Cu and/or Ni.

The requirements (i) have the reasons as described above.

The requirement (ii) described above has the reason as follows.

Specific measures were studied to retain residual austenite after processing, control the shape of the residual austenite grains, improve the hydrogen trapping capability and reliably reduce the concentration of diffusive hydrogen in the steel sheet to a harmless level by: (a) sufficiently suppressing the generation of hydrogen from the steel sheet in corrosive environment; and (b) suppressing hydrogen that has been generated from infiltrating the steel sheet.

It was found that it is very effective to include 0.003 to 0.5% of Cu and/or 0.003 to 1.0% of Ni in achieving the objectives of (a) and (b), and that the effect of improving hydrogen embrittlement resisting property through control of the structure can be achieved further by containing these elements.

Specifically, presence of Cu and Ni improves the corrosion resistance of the steel, and effectively suppresses the generation of hydrogen due to corrosion of the steel sheet. These elements also have the effect of promoting the generation of iron oxide, α-FeOOH, that is believed to be particularly stable thermodynamically and have protective property among various forms of rust generated in the atmosphere. By assisting the generation of this rust, it is made possible to suppress hydrogen that has been generated from infiltrating into the spring steel thereby to sufficiently improve the hydrogen embrittlement resisting property to endure in harsh corrosive environment. This effect can be achieved particularly satisfactorily when Cu and Ni are contained at the same time.

In order to achieve the effects described above, concentration of Cu, if added, should be 0.003% or higher, preferably 0.05% or higher and more preferably 0.1% or higher. Concentration of Ni, if added, should be 0.003% or higher, preferably 0.05% or higher and more preferably 0.1% or higher.

Since excessively high concentration of either Cu or Ni is detrimental to workability, it is preferable to limit the Cu content to 0.5% or lower and limit the Ni content to 1.0% or lower.

In order to achieve high hydrogen embrittlement resisting property after the forming process by retaining the predetermined amount of residual austenite after the forming process as in (i) described above, for example, 5% or more residual austenite may be contained in the steel sheet prior to the forming process, or large amount of fine residual austenite grains may be dispersed in the structure. Alternatively, forming process conditions may be controlled so as to make the residual austenite less likely to transform (for example, form the part by bending operation or control the forming temperature and/or stretching speed). The most desirable means of improving the workability and hydrogen embrittlement resisting property at the same time while maintaining the content of residual austenite before and after the processing substantially constant within an appropriate range and maintaining other properties (high strength, etc.) is to satisfy the requirements (A) and (B) described previously.

While this embodiment is characterized in that metal structure is controlled after processing and predetermined amount of Cu and/or Ni are added, it is necessary to control the other components as described below, in order to readily form the metal structure and efficiently improve hydrogen embrittlement resisting property and strength thereby to ensure ductility required for the thin steel sheet.

While constituent elements (C, Si, Mn, P, S, Al, Cu and/or Ni) of the steel of this embodiment are as described above with the rest substantially being Fe, it may include inevitable impurities introduced into the steel depending on the stock material, production material, manufacturing facility and other circumstances, containing 0.001% or less of N (nitrogen). In addition, other elements as described below may be intentionally added to such an extent that does not adversely affect the effects of the present invention.

<Ti and/or V: 0.003 to 1.0% in total>

Ti has the effect of assisting in the generation of protective rust, similarly to Cu and Ni. The protective rust has a very valuable effect of suppressing the generation of β-FeOOH that appears in chloride environment and has adverse effect on the corrosion resistance (and hence on the hydrogen embrittlement resisting property). Formation of such a protective rust is promoted particularly by adding Ti and V (or Zr). Ti renders the steel high corrosion resistance, and also has the effect of cleaning the steel.

V is effective in increasing the strength of the steel sheet and decreasing the size of crystal grains, in addition to having the effect of improving hydrogen embrittlement resistance through cooperation with Ti, as described previously.

In order to fully achieve the effect of Ti and/or V described above, it is preferable to add Ti and/or V in total concentration of 0.003% or higher (more preferably 0.01% or higher). For the purpose of improving hydrogen embrittlement resisting property, in particular, it is preferable to add more than 0.03% of Ti, more preferably 0.05% or more of Ti. However, the effects described above reach saturation when an excessive amount of Ti is added, resulting in economical disadvantage. Excessive V content also increases the precipitation of much carbonitride and leads to poor workability and lower hydrogen embrittlement resisting property. Therefore, it is preferable to control the total concentration of Ti and/or V to within 1.0%, more preferably within 0.5%.

<Zr: 0.003 to 1.0%>

Zr is effective in increasing the strength of the steel sheet and decreasing the crystal grain size, and also has the effect of improving hydrogen embrittlement resisting property through cooperation with Ti. In order to sufficiently achieve these effects, it is preferable that 0.003% or more of Zr is contained. However, excessive Zr content increases the precipitation of carbonitride and leads to poor workability and lower hydrogen embrittlement resisting property. Therefore, it is preferable to control the concentration of Zr to within 1.0%.

<Mo: 1.0% or Less (Higher than 0%)>

Mo has the effects of stabilizing austenite so as to retain the residual austenite, and suppress the infiltration of hydrogen thereby to improve hydrogen embrittlement resisting property. Mo also has the effect of improving the hardenability of the steel sheet. In addition, Mo strengthens the grain boundary so as to suppress hydrogen embrittlement from occurring. It is recommended to add 0.005% or more Mo in order to achieve these effects. More preferably 0.1% or more Mo is added. However, since the effects described above reach saturation when the Mo content exceeds 1.0%, resulting in economical disadvantage, Mo content is limited to 0.8% or less and more preferably to 0.5% or less.

<Nb: 0.1% or Less (Higher than 0%)>

Nb is very effective in increasing the strength of the steel sheet and decreasing the grain size of the structure. Nb achieves these effects particularly effectively in cooperation with Mo. In order to achieve these effects, it is recommended to include 0.005% or more of Nb. More preferably 0.01% or more of Nb is added. However, since the effects described above reach saturation when an excessive Nb content is included, resulting in economical disadvantage, Nb content is limited to 0.1% or less and more preferably to 0.08% or less.

<B: 0.0002 to 0.01%>

B is effective in increasing the strength of the steel sheet, and it is preferable that 0.0002% or more (more preferably 0.0005% or more) B is contained in order to achieve these effects. However, an excessive content of B leads to poor hot processing property. Therefore, it is preferable to control the concentration of B within 0.01% (more preferably within 0.005%).

<At Least One Kind Selected from Among a Group Consisting of Ca: 0.0005% to 0.005%, Mg: 0.0005% to 0.01% and REM: 0.0005% to 0.01%)

Ca, Mg and REM (rare earth element) are effective in suppressing an increase in hydrogen ion concentration, that is, a decrease in pH in the atmosphere of the interface due to corrosion of the steel sheet surface, thereby to improve the corrosion resistance of the steel sheet. It is also effective in controlling the form of sulfide in the steel and improving the workability of the steel. In order to achieve the effects described above, it is recommended to add each of Ca, Mg and REM in concentration of 0.0005% or higher. However, since excessive contents of these elements leads to poor workability, it is preferable to keep the concentrations of Ca within 0.005%, Mg and REM each within 0.01%.

Third Embodiment

A third high strength thin steel sheet according to the present invention is constituted from higher than 0.25 and up to 0.60% of C (contents of components given in terms of percentage in this patent application all refer to percentage by weight), 1.0 to 3.0% of Si, 1.0 to 3.5% of Mn, 0.15% or less of P, 0.02% or less of S, 1.5% or less (higher than 0%) of Al, while iron and inevitable impurities making up the rest,

wherein: (iii) the structure satisfies the following requirements after forming:

1% or more residual austenite; the mean axis ratio (major axis/minor axis) of the residual austenite grains is 5 or higher;

mean length of minor axes of the residual austenite grains is 1 μm or less; and

minimum distance between residual austenite grains is 1 μm or less.

When the metal structure is controlled as described above, hydrogen embrittlement resisting property of the high strength thin steel sheet can be sufficiently improved without adding much alloy elements.

The phrase “after the forming process” means the state of the steel sheet after being stretch formed with elongation of 3%. Specifically, the steel sheet is subjected to elongation of 3% by uniaxial stretching at the room temperature (the stretch forming process of 3% elongation may hereinafter be referred to simply as “process”).

The requirements for the residual austenite of the present invention will now be described in detail below.

<1% or More Residual Austenite>

<Mean Axis Ratio (Major Axis/Minor Axis) of the Residual Austenite Grains is 5 or Higher>

It is necessary that the metal structure contains 1% or more residual austenite in terms of area proportion to the entire structure after processing, in order to achieve high hydrogen embrittlement resisting property in harsh operating environment over an extended period of time after forming the part. Residual austenite contributes not only to the improvement of hydrogen embrittlement resisting property as described above, but also to the improvement of total elongation as has been known in the prior art. Content of the residual austenite is preferably 2% or higher, and more preferably 3% or higher. Since the desired level of high strength cannot be obtained when an excessive amount of residual austenite is contained, it is recommended to set an upper limit of 15% (more preferably 10%) to the residual austenite content.

Lath-shaped grains of residual austenite after processing have far higher capacity of trapping hydrogen than carbide. FIG. 1 is a graph showing the relationship between the mean axis ratio of the residual austenite grains measured by a method to be described later and hydrogen embrittlement risk index (measured by a method to be described later in an example, lower value of this index means better hydrogen embrittlement resisting property). From FIG. 1, it can be seen that hydrogen embrittlement risk index sharply decreases when the mean axis ratio (major axis/minor axis) of the residual austenite grains increases beyond 5. This is supposedly because, when the mean axis ratio of the residual austenite grains becomes 5 or higher, intrinsic capability of the residual austenite to absorb hydrogen is put into full play, so that the residual austenite attains far higher capacity of trapping hydrogen than carbide and substantially neutralizes the hydrogen that infiltrates from the outside through atmospheric corrosion thereby to achieve remarkable achievement in hydrogen embrittlement resisting property. The mean axis ratio of the residual austenite grains is preferably 10 or higher, and more preferably 15 or higher.

<Mean Length of Minor Axes of the Residual Austenite Grains is 1 μm or Less>

According to the present invention, it has been found that hydrogen embrittlement resisting property can be effectively improved by dispersing fine grains of residual austenite of lath shape. Specifically, hydrogen embrittlement resisting property can be surely improved by dispersing the lath-shape grains of residual austenite having sizes of 1 μm or less (submicrometer order). This is supposedly because surface area of the residual austenite grains (interface) increases resulting in larger hydrogen trapping capability, when larger number of fine lath-shape grains of residual austenite having smaller mean length of minor axis are dispersed. Mean length of minor axes of the residual austenite grains is preferably 0.5 μm or less, more preferably 0.25 μm or less.

According to the present invention, hydrogen trapping capability of the fine lath-shape grains of residual austenite can be made far greater than that in the case of dispersing carbide, and thereby to substantially neutralize hydrogen that infiltrates from the outside through atmospheric corrosion, even when the same proportion by volume of residual austenite is contained, by controlling the mean axis ratio and mean length of minor axes of the residual austenite grains as described above.

<Minimum Distance Between Residual Austenite Grains is 1 μm or Less>

According to the present invention, it has been found that hydrogen embrittlement resisting property can be improved further by controlling the minimum distance between adjacent residual austenite grains, in addition to the above. Specifically, hydrogen embrittlement resistance can be surely improved when the minimum distance between residual austenite grains is 1 μm or less. This is supposedly because propagation of cracks is suppressed so that the structure demonstrates higher resistance against fracture, when a large number of fine lath-shape grains of residual austenite are dispersed in proximity to each other. Minimum distance between adjacent residual austenite grains is preferably 0.8 μm or less, and more preferably 0.5 μm or less.

The residual austenite refers to a region that is observed as FCC (face centered cubic lattice) by the FE-SEM/EBSP method which will be described later. Measurement by the EBSP may be done, for example, by measuring a measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness at measuring intervals of 0.1 μm. The measuring surface is prepared by electrolytic polishing in order to prevent the residual austenite from transforming. Then the test piece is set in the lens barrel of an FE-SEM equipped with the EBSP detector (of which details will be described later) and is irradiated with electron beam. An EBSP image projected onto a screen is captured by a high sensitivity camera (VE-1000-SIT manufactured by Dage-MTI Inc.) and is sent to a computer. The computer carries out image analysis and generates color mapping of the FCC phase through comparison with a structural pattern simulated with a known crystal system (FCC (face centered cubic lattice) phase in the case of residual austenite). Area proportion of the region that is mapped as described above is taken as the area proportion of the residual austenite. This analysis was carried out by means of hardware and software of OIM (Orientation Imaging Microscopy™) system of TexSEM Laboratories Inc.

The mean axis ratio, mean length of minor axes and minimum distance between residual austenite grains were determined as follows. The mean axis ratio of the residual austenite grains was determined by measuring the major axis and minor axis of residual austenite crystal grain existing in each of three arbitrarily chosen fields of view in the observation by means of TEM (transmission electron microscope) with magnification factor of 15000, and averaging the ratios of major axis to minor axis. The mean length of minor axes of the residual austenite grains was determined by averaging the lengths of minor axes measured as described above. The minimum distance between adjacent residual austenite grains was determined by measuring the distance between adjacent residual austenite grains that were aligned in the direction of major axis as shown as (a) in FIG. 11 (distance (b) in FIG. 11 is not regarded as the minimum distance between the grains) by observing with TEM (magnification factor of 15000) in each of three arbitrarily chosen fields of view and averaging the distances measured in the three fields of view.

In order to decrease the number of intergranular fracture initiating points in the steel thereby to surely decrease the concentration of diffusive hydrogen to a harmless level and achieve a high strength, it is desirable to form the matrix phase of the steel sheet after processing from a binary phase structure of bainitic ferrite and martensite with the bainitic ferrite acting as the main phase, instead of the single phase structure of martensite that is generally used for high strength steels.

In the case of single phase structure of martensite, a carbide (for example, film-like cementite) is likely to precipitate in the grain boundaries, thus making intergranular fracture likely to occur. In the case of the binary phase structure of bainitic ferrite and martensite with the bainitic ferrite acting as the main phase, in contrast, the bainitic ferrite is a hard phase and therefore it is easy to increase the strength of the entire structure as in the case of the single phase of martensite. The hydrogen embrittlement resisting property can also be improved as much hydrogen is trapped in the dislocations. It also has such an advantage that coexistence of the bainitic ferrite and residual austenite which will be described later prevents the generation of carbide that acts as the intergranular fracture initiating points, and it becomes easier to create the lath-shaped residual austenite in the boundaries of lath-shaped bainitic ferrite.

Accordingly, it is required in the present invention that the binary phase structure of bainitic ferrite and martensite occupy 80% or more, preferably 85% or more and more preferably 90% or more of the entire structure after the stretch forming processing to elongate by 3%. Upper limit of the proportion may be determined by the balance with other structure (residual austenite), and is set to 99% when other structure (ferrite, etc.) than the residual austenite is not contained.

The bainitic ferrite referred to in the present invention is plate-shaped ferrite having a lower structure of high density of dislocations. It is clearly distinguished from polygonal ferrite that has lower structure including no or very low density of dislocations, by SEM observation as follows.

Area proportion of bainitic ferrite structure is determined as follows. A test piece is etched with Nital etchant. A measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness is observed with SEM (scanning electron microscope) (magnification factor of 1500) thereby to determine the area proportion.

Bainitic ferrite is shown with dark gray color in SEM photograph (bainitic ferrite, residual austenite and martensite may not be distinguishable in the case of SEM observation), while polygonal ferrite is shown black in SEM photograph and has polygonal shape that does not contain residual austenite and martensite inside thereof.

The SEM used in the present invention is a high-resolution FE-SEM (Field Emission type Scanning Electron Microscope XL30S-FEG manufactured by Philips Inc.) equipped with an EBSP (Electron Back Scatter diffraction Pattern) detector, that has a merit of being capable of analyzing the area observed by the SEM at the same time by means of the EBSP detector. EBSP detection is carried out as follows. When the sample surface is irradiated with electron beam, the EBSP detector analyzes the Kikuchi pattern obtained from the reflected electrons, thereby to determine the crystal orientation at the point where the electron beam has hit upon. Distribution of orientations over the sample surface can be measured by scanning the electron beam two-dimensionally over the sample surface while measuring the crystal orientation at predetermined intervals. The EBSP detection method has such an advantage that different structures that are regarded as the same structure in the ordinary microscopic observation but have different crystal orientations can be distinguished by the difference in color tone.

The metal structure after the processing may be constituted either from only the structures described above (namely, a mixed structure of bainitic ferrite+martensite and residual austenite), or may include other structure such as ferrite (the term ferrite used herein refers to polygonal ferrite, that is a ferrite structure that includes no or very few dislocations) or pearlite to such an extent that the effect of the present invention is not compromised. Such additional components are structures that can inevitably remain in the manufacturing process of the present invention, of which concentration is preferably as low as possible, within 9%, preferably less than 5% and more preferably less than 3% according to the present invention.

In order to maintain high hydrogen embrittlement resisting property after the forming process, for example, high proportion of residual austenite, 5% or more, may be contained in the steel sheet prior to the forming process, or a large amount of fine residual austenite grains may be dispersed in the structure. Alternatively, forming process conditions may be controlled so as to make the residual austenite less likely to transform (for example, form the part by bending operation or control the forming temperature and/or stretching speed). The most desirable means of improving the workability and hydrogen embrittlement resisting property at the same time while maintaining the content of residual austenite before and after the processing substantially constant within an appropriate range and maintaining other properties (high strength, etc.) is to satisfy the requirements (A) and (B) described previously.

While this embodiment is characterized in that the metal structure is controlled after processing, it is necessary to control the other components as described previously, in order to form the metal structure and efficiently improve hydrogen embrittlement resisting property and strength thereby to ensure the level of ductility required for the thin steel sheet.

While the present invention does not specify the manufacturing conditions, it is recommended to apply heat treatment in the following procedure after hot rolling or cold rolling conducted thereafter, in order to form the structure described above that can be easily worked and has high strength and high hydrogen embrittlement resistance after the processing, by using the steel material of the composition described above. The recommended procedure is to keep the steel the composition described above at a temperature (T1) in a range from A3 point to (A3 point+50° C.) for a period of 10 to 1800 seconds (t1), cool down the steel at a mean cooling rate of 3° C./s or higher to a temperature (T2) in a range from Ms point to Bs point and keep the material at this temperature for a period of 60 to 3600 seconds (t2).

It is not desirable that the temperature T1 becomes higher than (A3 point+50° C.) or the period t1 is longer than 1800 seconds, in which case austenite grains grow resulting in poor workability (elongation flanging property). When the temperature T1 is lower than A3 point, on the other hand, desirable bainitic ferrite structure cannot be obtained. When the period t1 is shorter than 10 seconds, austenitization does not proceed sufficiently and therefore cementite and other alloy carbides remain. The period t1 is preferably in a range from 30 to 600 seconds, more preferably from 60 to 400 seconds.

Then the steel sheet is cooled down. The steel is cooled at a mean cooling rate of 3° C./s or higher, for the purpose of preventing pearlite structure from being generated while avoiding the pearlite transformation region. The mean cooling rate should be as high as possible, and is preferably 5° C./s or higher, and more preferably 10° C./s or higher.

After quenching to the temperature between Ms point and Bs point at the rate described above, the steel is subjected to isothermal transformation so as to transform the matrix phase into binary phase structure of bainitic ferrite and martensite. When the heat retaining temperature T2 is higher than Bs, much pearlite that is not desirable for the present invention is formed, thus hampering the formation of the predetermined bainitic ferrite structure. When T2 is below Ms, on the other hand, the amount of residual austenite decreases.

When the temperature holding period t2 is longer than 1800 seconds, density of dislocations in bainitic ferrite becomes low, the amount of trapped hydrogen decreases and the desired residual austenite cannot be obtained. When t2 is less than 60 second, on the other hand, desired bainitic ferrite structure cannot be obtained. The length of t2 is preferably from 90 to 1200 seconds, and more preferably from 120 to 600 seconds. There is no restriction on the method of cooling after maintaining the heating temperature, and air cooling, quenching or air-assisted water cooling may be employed.

In the practical manufacturing process, the annealing process described above can be carried out easily by employing a continuous annealing facility or a batch annealing facility. In case that a cold rolled sheet is plated with zinc by hot dipping, the heat treatment process may be replaced by the plating process by setting the plating conditions so as to satisfy the heat treatment conditions. The plating may also be alloyed.

There is no restriction on the hot rolling process (or cold rolling process as required) that precedes the continuous annealing process described above, and commonly employed process conditions may be used. Specifically, the hot rolling process may be carried out in such a procedure as, after hot rolling at a temperature above Ar3 point, the steel sheet is cooled at a mean cooling rate of about 30° C./s and is wound up at a temperature from about 500 to 600° C. In case that the hot rolled steel sheet has unsatisfactory appearance, cold rolling may be applied in order to rectify the appearance. It is recommended to set the cold rolling ratio in a range from 1 to 70%. Cold rolling beyond 70% leads to excessive rolling load that makes it difficult to carry out the cold rolling.

While the present invention is addressed to thin steel sheet, there is no limitation to the form of product, and may be applied, in addition to steel sheet made by hot rolling or steel sheet made by cold rolling, to those subjected to annealing after hot rolling or cold rolling, followed by chemical conversion treatment, hot-dip coating, electroplating, vapor deposition, painting, priming for painting, organic coating treatment or the like.

The plating process may be either galvanizing or aluminum plating. The method of plating may be either hot-dip coating or electroplating, and the plating process may also be followed by alloying heat treatment or multi-layer plating. A steel sheet, that is plated or not plated, may also be laminated with a film.

When the coating operation described above is carried out, chemical conversion treatment such as phosphating or electrodepositing coating may be applied in accordance to the application. The coating material may be a known resin that can be used in combination with a known hardening agent such as epoxy resin, fluorocarbon resin, silicone acrylic resin, polyurethane resin, acrylic resin, polyester resin, phenol resin, alkyd resin, or melamine resin. Among these, epoxy resin, fluorocarbon resin or silicone acrylic resin is preferably used in consideration of corrosion resistance. Known additives that are added to coating materials such as coloring agent, coupling agent, leveling agent, sensitization agent, antioxidant agent, anti-UV protection agent, flame retarding agent or the like may be used.

There is also no restriction on the coating and solvent-based coating, powder coating, water-based coating, water-dispersed coating, electrodeposition coating or like may be employed. Desired coating layer of the coating material described above can be formed on the steel by a known technique such as dipping, roll coater, spraying, or curtain flow coater. The coating layer may have any proper thickness.

The high strength thin steel sheet of the present invention may be applied to high-strength automotive components such as bumper, door impact beam, pillar and other reinforcement members and interior parts such as seat rail, etc. Automobile components that are manufactured by forming process also have sufficient properties (strength) and high hydrogen embrittlement resisting property.

The present invention will now be described below by way of examples, but the present invention is not limited to the examples. Various modifications may be conceived without departing from the technical scope of the present invention.

Example 1

Sample steels A-1 through Y-1 having the compositions described in Table 1 were melt-refined in vacuum to make test slabs. The slabs were processed in the following procedure (hot rolling→cold rolling→continuous annealing) thereby to obtain hot-rolled steel plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.

<Hot rolling> Starting temperature (SRT): Held at a temperature between 1150 and 1250° C. for 30 minutes.

Finishing temperature (FDT): 850° C.

Cooling rate: 40° C./s

Winding-up temperature: 550° C.

<Cold rolling> Rolling ratio: 50%

<Continuous annealing> Each steel specimen was kept at a temperature of A3 point+30° C. for 120 seconds, then cooled in air at a mean cooling rate of 20° C./s to temperature T0 shown in Table 2, and was kept at T0 for 240 seconds, followed by air-assisted water cooling to the room temperature.

No. 116 shown in Table 2 was made by heating a cold-rolled steel sheet to 830° C., keeping at this temperature for 5 minutes followed by quenching in water and tempering at 300° C. for 10 minutes, thereby to form a martensite steel as a comparative example of the high-strength steel of the prior art. No. 120 was made by heating a cold-rolled steel sheet to 800° C., keeping at this temperature for 120 seconds, cooling down at a mean cooling rate of 20° C./s to 350° C. and keeping at this temperature for 240 seconds.

JIS No. 5 test pieces were prepared from the steel sheets obtained as described above, and were subjected to stretch forming process with elongation of 3% mimicking the actual manufacturing process. Metal structures of the test pieces were observed before and after the processing, tensile strength (TS) and elongation (total elongation E1) before the processing and hydrogen embrittlement resisting property after the processing were measured by the following procedures.

Observation of Metal Structure

Metal structures of the test pieces were observed before and after the processing as follows. A measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness was photographed at measuring intervals of 0.1 μm, and area proportions of bainitic ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by the method described previously. Then similar measurements were made in two fields of view that were arbitrarily selected, and the measured values were averaged. Area proportions of other structures (ferrite, pearlite, etc.) were subtracted from the entire structure.

Mean axis ratio of the residual austenite grains of the steel sheet before and after the processing were measured by the method described previously. Test pieces having mean axis ratio of 5 or higher were regarded to satisfy the requirements of the present invention (∘), and those having mean axis ratio of lower than 5 were regarded to fail to satisfy the requirements of the present invention (x).

Measurement of tensile strength (TS) and elongation (E1) Tensile test was conducted on the JIS No. 5 test piece before processing, so as to measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa as measured by the method described previously, those which showed elongation of 10% or more were evaluated as high in elongation property.

Evaluation of Hydrogen Embrittlement Resisting Property

In order to evaluate hydrogen embrittlement resisting property, the JIS No. 5 test piece was stretched so as to elongate by 3%. Then after bending with a radius of curvature of 15 mm, load of 1000 MPa was applied and the test piece was immersed in 5% solution of hydrochloric acid, and the time before crack occurred was measured.

Hydrogen-charged 4-point bending test was also conducted for some steel species. Specifically, a rectangular test piece measuring 65 mm by 10 mm made of each steel sheet elongated by 3% was immersed in a solution of 0.5 mol of H₂SO₄ and 0.01 mol of KSCN and was subjected to cathode hydrogen charging. Maximum stress endured without breaking for 3 hours was determined as the critical fracture stress (DFL).

Results of these tests are shown in Table 2.

TABLE 1
Steel species	Chemical composition (mass %)*	Ac3	Bs	Ms
Symbol	C	Si	Mn	P	S	Al	Mo	Nb	Others	(° C.)	(° C.)	(° C.)
A-1	0.40	2.01	2.01	0.012	0.002	0.033	0.2	0.05	—	839.1	524.5	300.9
B-1	0.27	2.02	2.54	0.011	0.002	0.031	0.2	0.06	—	845.0	511.9	345.0
C-1	0.55	2.00	2.54	0.011	0.002	0.031	0.2	0.06	—	799.1	436.3	212.3
D-1	0.39	2.55	2.51	0.011	0.002	0.030	0.2	0.05	—	847.9	482.2	289.1
E-1	0.41	2.01	1.32	0.011	0.002	0.031	0.2	0.04	—	856.7	583.9	318.9
F-1	0.40	2.02	3.15	0.011	0.002	0.030	0.2	0.05	—	803.4	421.9	263.3
G-1	0.40	2.02	2.51	0.012	0.002	0.031	0.8	0.05	—	842.6	429.7	271.8
H-1	0.40	1.98	2.48	0.011	0.002	0.033	0.2	0.08	—	822.9	482.2	285.4
I-1	0.40	2.02	2.50	0.011	0.002	0.031	0.2	0.05	B: 0.0005	823.3	480.4	284.7
J-1	0.40	1.98	2.53	0.011	0.002	0.033	0.2	0.05	Ca: 0.004	821.4	477.7	283.7
K-1	0.40	1.49	2.48	0.012	0.002	0.033	0.2	0.05	Mg: 0.005	801.7	482.2	285.4
L-1	0.39	1.51	2.50	0.011	0.002	0.032	0.2	0.06	REM: 0.005	802.5	483.1	289.4
M-1	0.35	1.50	2.50	0.011	0.002	0.033	0.2	0.05	Ca: 0.005, Mg: 0.005,	809.2	493.9	308.4
									B: 0.0005
N-1	0.19	1.51	1.52	0.014	0.005	0.031	0.2	0.05	—	871.9	625.3	416.6
O-1	0.30	1.40	0.90	0.011	0.005	0.032	0.2	0.05	—	861.2	651.4	384.9
P-1	0.35	0.16	1.42	0.014	0.002	0.043	0.2	0.05	—	787.8	591.1	344.0
Q-1	0.7	2.01	2.01	0.013	0.003	0.033	0.2	0.05	—	798.3	443.5	158.7
R-1	0.40	2.02	1.60	0.012	0.002	0.033	1.5	0.05	—	892.8	453.5	287.1
S-1	0.40	2.00	1.20	0.012	0.002	0.031	0.2	0.2	—	862.1	597.4	327.6
T-1	0.27	2.01	2.50	0.011	0.002	0.053	0.2	0.05	—	854.6	515.5	346.3
U-1	0.28	2.02	2.50	0.012	0.002	0.263	0.2	0.05	—	937.8	512.8	341.6
V-1	0.28	2.02	2.51	0.011	0.002	0.418	0.2	0.05	Ca: 0.004	998.8	511.9	341.3
W-1	0.28	2.02	2.51	0.011	0.002	0.681	0.2	0.05	Ca: 0.004	1104.0	511.9	341.3
X-1	0.27	1.98	2.55	0.012	0.002	1.01	0.2	0.05	Mg: 0.005, B: 0.0005	1235.2	511.0	344.7
Y-1	0.28	2.02	2.60	0.011	0.002	1.68	0.2	0.05	—	1500.9	503.8	338.3
*The balance consists of iron and inevitable impurities.

	TABLE 2
	Before processing	After processing
	Steel					Mean axis						Mean axis	Hydrogen
	species	To	Residual γ	BF + M	Others	ratio of	TS	El	Residual γ	BF + M	Others	ratio of	embrittlement	DFL
Test No.	Symbol	° C.	%	%	%	residual γ	MPa	%	%	%	%	residual γ	h	MPa
101	A-1	350	9	91	0	∘	1510	14	5	95	0	∘	Over 24	—
102	B-1	350	7	93	0	∘	1495	12	3	97	0	∘	Over 24	420
103	C-1	350	14	86	0	∘	1521	13	5	95	0	∘	Over 24	—
104	D-1	350	12	88	0	∘	1495	13	4	96	0	∘	Over 24	—
105	E-1	350	11	89	0	∘	1417	14	4	96	0	∘	Over 24	—
106	F-1	350	12	88	0	∘	1376	14	5	95	0	∘	Over 24	—
107	G-1	350	11	88	1	∘	1433	14	4	96	0	∘	Over 24	—
108	H-1	350	11	88	1	∘	1476	14	3	97	0	∘	Over 24	—
109	I-1	350	12	88	0	∘	1486	13	4	96	0	∘	Over 24	—
110	J-1	350	11	89	0	∘	1491	12	3	97	0	∘	Over 24	—
111	K-1	350	12	88	0	∘	1488	13	4	96	0	∘	Over 24	—
112	L-1	350	12	88	0	∘	1487	14	3	97	0	∘	Over 24	—
113	M-1	350	11	89	0	∘	1454	13	4	96	0	∘	Over 24	—
114	N-1	430	7	93	0	∘	1417	8	<1	99	<1	∘	Over 24	—
115	O-1	350	<1	99	<1	x	1436	8	<1	99	<1	x	11	—
116	P-1	—	<1	99	<1	x	1503	4	<1	99	<1	x	7	—
117	Q-1	350	15	85	0	∘	1511	3	3	97	0	∘	3	—
118	R-1	350	11	89	0	∘	1341	1	—	—	—	—	—	—
119	S-1	350	11	89	0	∘	1370	2	—	—	—	—	—	—
120	A-1	350	8	20	72	x	960	13	15	85	0	x	Over 24	—
121	T-1	350	7	93	0	∘	1495	12	3	97	0	∘	Over 24	515
122	U-1	350	8	92	0	∘	1499	12	4	96	0	∘	Over 24	650
123	V-1	350	8	92	0	∘	1503	13	4	96	0	∘	Over 24	680
124	W-1	350	8	92	0	∘	1505	13	5	95	0	∘	Over 24	700
125	X-1	350	8	92	0	∘	1498	13	5	95	0	∘	Over 24	715
126	Y-1	350	12	59	29	x	1290	14	9	91	0	x	8	290

The results shown in Tables 1 and 2 can be interpreted as follows (numbers in the following description are test Nos. in Table 2).

Test pieces Nos. 101 through 113 (inventive steel sheets 2) and test pieces Nos. 121 through 125 (inventive steel sheets 1) that satisfy the requirements of the present invention have high strength of 1180 MPa or higher, and high hydrogen embrittlement resisting property in harsh environment after the forming process. They also have high elongation property required of the TRIP steel sheet, thus providing steel sheets best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere. Test pieces Nos. 121 through 125, in particular, show even better hydrogen embrittlement resisting property.

Test pieces Nos. 114 through 120 and 126 that do not satisfy the requirements of the present invention, in contrast, have the following drawbacks.

No. 114 made of steel species N-1 that includes insufficient C content does not have good workability.

No. 115 made of steel species O-1 that includes insufficient Mn content does not retain sufficient residual austenite and is inferior in hydrogen embrittlement resisting property after the processing.

No. 116, martensite steel that is a conventional high strength steel made of steel species P-1 that includes insufficient Si content, hardly contains residual austenite and is inferior in hydrogen embrittlement resisting property. It also does not show the elongation property required of a thin steel sheet.

No. 117 made of steel species Q-1 that includes excessive C content has precipitation of carbide and is inferior in both forming workability and hydrogen embrittlement resisting property after processing.

No. 118 made of steel species R-1 that includes excessive Mo content and No. 119 made of steel species S-1 that includes excessive Nb content are inferior in forming workability. Nos. 118 and 119 could not undergo the processing, making it impossible to investigate the property after the processing.

No. 120, that was made of a steel that has the composition specified in the present invention but was not manufactured under the recommended conditions, resulted in the conventional TRIP steel. As a result, the residual austenite does not have the mean axis ratio specified in the present invention, while the matrix phase is not formed in binary phase structure of bainitic ferrite and martensite, and therefore sufficient level of hydrogen embrittlement resisting property is not achieved.

No. 126 includes Al content higher than that specified for the inventive steel sheet 1. As a result, although the predetermined amount of residual austenite is retained, the residual austenite does not have the mean axis ratio specified in the present invention, the desired matrix phase is not obtained and inclusions such as AlN are generated thus resulting in poor hydrogen embrittlement resisting property.

Then parts were made by using steel species A-1, J-1 shown in Table 1 and comparative steel sheet (590 MPa class high strength steel sheet of the prior art). Performance (pressure collapse resistance and impact resistance) of the formed test piece were studied by conducting pressure collapse test and impact resistance test as follows.

Pressure Collapse Test

The part 1 (hat channel as test piece) shown in FIG. 1 was made by using steel species A-1, J-1 shown in Table 1 and the comparative steel sheet, and was subjected to pressure collapse test. The part was spot welded at the positions 2 of the part shown in FIG. 1 at 35 mm intervals as shown in FIG. 1 by supplying electric current of a magnitude less than the expulsion generating current by 0.5 kA from an electrode measuring 6 mm in diameter at the distal end. Then a die 3 was pressed against the part 1 from above the mid portion thereof in the longitudinal direction as shown in FIG. 2, and the maximum tolerable load was determined. Absorbed energy was determined from the area under the load-deformation curve. The results are shown in Table 3.

	TABLE 3
	Evaluation
	of test piece
Steel sheet used	Maximum	Energy
	TS	EL	Residual γ	load	absorbed
Steel species	(MPa)	(%)	(Area %)	(kN)	(kJ)
Symbol A-1	1510	15	9	14.1	0.72
Symbol J-1	1491	12	11	14	0.68
Comparative steel	613	22	0	5.7	0.33
sheet

From Table 3, it can be seen that the part (test piece) made from the steel sheet of the present invention has higher load bearing capability and absorbs greater energy than a part made of the conventional steel sheet having lower strength, thus showing high pressure collapse resistance.

Impact Resistance Test

The parts 4 (hat channel as test piece) shown in FIG. 3 were made by using steel species A-1, J-1 shown in Table 1 and the comparative steel sheet, and were subjected to impact resistance test. FIG. 4 is a sectional view along A-A of the part 4 shown in FIG. 3. In the impact resistance test, after the part was spot welded at the positions 5 of the part 4 similarly to the pressure collapse test, the part 4 was placed on a base 7 as schematically shown in FIG. 5. A weight 6 (weighing 10 kg) was dropped onto the part 4 from a height of 11 meters, and the energy absorbed before the part 4 underwent deformation of 40 mm in the direction of height. The results are shown in Table 4.

TABLE 4
	Evaluation of
Steel sheet used	test piece
	TS	EL	Residual γ	Energy absorbed
Steel species	(MPa)	(%)	(Area %)	(kJ)
Symbol A-1	1510	15	9	6.94
Symbol J-1	1491	12	11	6.65
Comparative	613	22	0	3.56
steel sheet

From Table 4, it can be seen that the part (test piece) made from the steel sheet of the present invention absorbs greater energy than a part made of the conventional steel sheet that has lower strength, thus showing higher impact resistance.

TEM photograph of the test piece made in this example is shown as reference. FIG. 6 is a photograph of TEM observation of No. 101 of the present invention. From FIG. 6, it can be seen that the high strength thin steel sheet of the present invention contains lath-shaped residual austenite (black portion of bar shape in FIG. 6) specified in the present invention dispersed therein. FIG. 7 is a photograph of TEM observation of No. 120 of a comparative example. From FIG. 7, it can be seen that the high strength thin steel sheet of No. 120 contains residual austenite (black portion of somewhat round shape in FIG. 7), although the residual austenite has a block shape that does not satisfy the requirements of the present invention.

Example 2

Sample steels A-2 through Y-2 having the compositions described in Table 5 were melt-refined in vacuum to make test slabs. The slabs were processed in the following procedure (hot rolling→cold rolling→continuous annealing) thereby to obtain hot-rolled steel plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.

<Hot Rolling>

Starting temperature (SRT): Held at a temperature between 1150 and 1250° C. for 30 minutes.

Finishing temperature (FDT): 850° C.

Cooling rate: 40° C./s

Winding-up temperature: 550° C.

<Cold Rolling>

Rolling ratio: 50%

<Continuous Annealing>

Each steel specimen was kept at a temperature of A3 point+30° C. for 120 seconds, then rapidly cooled (air cooling) at a mean cooling rate of 20° C./s to temperature T0 shown in Table 6, and was kept at T0 for 240 seconds, followed by air-assisted water cooling to the room temperature.

No. 217 in Table 6 was made by heating a cold-rolled steel sheet to 830° C., keeping at this temperature for 5 minutes followed by quenching in water and tempering at 300° C. for 10 minutes, thereby to form a martensite steel as a comparative example of the high-strength steel of the prior art. No. 220 was made by heating a cold-rolled steel sheet to 800° C., keeping at this temperature for 120 seconds, cooling down at a mean cooling rate of 20° C./s to 350° C. and keeping at this temperature for 240 seconds.

JIS No. 5 test pieces were prepared from the steel sheets obtained as described above, and were subjected to stretch forming process with elongation of 3% mimicking the actual manufacturing process. Metal structures of the test pieces were observed before and after the processing, tensile strength (TS) and elongation (total elongation E1) before the processing and hydrogen embrittlement resisting property after the processing were measured by the following procedures.

Observation of Metal Structure

Metal structures of the test pieces were observed before and after the processing as follows. A measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness was photographed at measuring intervals of 0.1 μm, and area proportions of bainitic ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by the method described previously. Then similar measurements were made in two fields of view that were arbitrarily selected, and the measured values were averaged. Area proportions of other structures (ferrite, pearlite, etc.) were subtracted from the entire structure.

Mean axis ratio of the residual austenite grains of the steel sheet before and after the processing were measured by the method described previously. Test pieces having mean axis ratio of 5 or higher were regarded to satisfy the requirements of the present invention (∘), and those having mean axis ratio of lower than 5 were regarded to fail to satisfy the requirements of the present invention (x).

Measurement of Tensile Strength (TS) and Elongation (E1)

Tensile test was conducted on the JIS No. 5 test piece before processing, so as to measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa as measured by the method described previously, those which showed elongation of 10% or more were evaluated as high in elongation property.

Evaluation of Hydrogen Embrittlement Resisting Property

In order to evaluate the hydrogen embrittlement resisting property, the JIS No. 5 test piece was stretched so as to elongate by 3%. Then after bending with a radius of curvature of 15 mm, load of 1000 MPa was applied and the test piece was immersed in 5% solution of hydrochloric acid, to measure the time before crack occurred.

The bent test pieces prepared as described above were subjected to accelerated exposure test in which 3% solution of NaCl was sprayed once every day for 30 days simulating the actual operating environment, and the number of days before crack occurred was determined.

Hydrogen-charged 4-point bending test was also conducted for some steel species. Specifically, a rectangular test piece measuring 65 mm by 10 mm made of each steel sheet elongated by 3% was immersed in a solution of 0.5 mol of H₂SO₄ and 0.01 mol of KSCN and was subjected to cathode hydrogen charging. Maximum stress endured without breaking for 3 hours was determined as the critical fracture stress (DFL). Then the ratio (DFL ratio) of this value to the value of DFL of test No. 203 (steel species C-2) shown in Table 6 was determined.

Results of these tests are shown in Table 6.

TABLE 5
Steel
species	Chemical composition (mass %)*	Ac3	Bs	Ms
Symbol	C	Si	Mn	P	S	Al	Cu	Ni	Ti	V	Nb	Ma	Others	° C.	° C.	° C.
A-2	0.40	2.03	2.01	0.012	0.002	0.033	0.3	—	—	—	—	—	—	827.7	541.1	305.1
B-2	0.27	2.02	2.48	0.011	0.002	0.031	—	0.3	—	—	—	—	—	836.0	522.8	346.1
C-2	0.40	2.55	2.51	0.011	0.002	0.030	—	0.3	—	—	—	—	—	835.4	485.0	283.5
D-2	0.39	2.01	1.25	0.011	0.002	0.031	—	0.3	—	—	—	—	—	851.1	601.1	329.8
E-2	0.41	2.00	3.15	0.011	0.002	0.030	—	0.3	—	—	—	—	—	790.1	424.7	257.6
F-2	0.40	1.99	2.53	0.011	0.002	0.033	0.3	0.05	—	—	—	—	—	808.8	492.5	287.1
G-2	0.41	2.01	2.50	0.011	0.002	0.031	0.3	0.2	0.05	—	—	—	—	825.9	486.9	280.8
H-2	0.40	1.98	2.48	0.011	0.002	0.033	0.3	0.2	—	0.05	—	—	—	812.8	491.4	286.2
I-2	0.40	2.02	2.50	0.011	0.002	0.031	0.3	0.2	0.05	0.05	—	—	—	833.2	489.6	285.5
J-2	0.40	1.51	2.48	0.011	0.002	0.033	0.3	0.2	0.035	—	—	—	Zr: 0.02	800.6	491.4	286.2
K-2	0.39	2.02	2.50	0.011	0.002	0.032	0.3	0.2	0.05	—	0.05	0.2	—	836.3	475.7	286.0
L-2	0.40	2.01	2.50	0.011	0.002	0.033	0.3	0.2	0.05	—	0.06	0.2	—	834.6	473.0	281.3
M-2	0.35	1.50	2.50	0.012	0.002	0.033	0.3	0.2	0.05	—	0.05	0.2	B: 0.0005	820.8	486.5	305.0
N-2	0.55	1.52	1.53	0.012	0.002	0.033	0.3	0.2	0.05	—	0.05	0.2	Ca: 0.004,	820.4	519.8	242.2
													Mg: 0.005
O-2	0.19	1.51	1.52	0.011	0.005	0.031	—	0.2	—	—	—	—	—	860.5	634.5	417.4
P-2	0.30	1.40	0.90	0.011	0.005	0.032	—	0.2	—	—	—	—	—	851.9	660.6	385.7
Q-2	0.35	0.16	1.42	0.012	0.002	0.043	—	0.2	—	—	—	—	—	777.0	600.3	344.8
R-2	0.7	2.01	2.01	0.012	0.002	0.033	—	0.2	—	—	—	—	—	788.3	452.7	159.5
S-2	0.38	2.02	1.20	0.012	0.002	0.033	—	—	—	—	—	—	—	860.8	619.4	341.3
T-2	0.40	2.02	2.01	0.012	0.002	0.050	0.3	0.2	—	—	—	—	—	831.0	533.7	301.7
U-2	0.39	2.01	2.02	0.011	0.002	0.250	0.3	0.2	0.05	—	—	—	—	931.1	535.5	306.1
V-2	0.41	1.98	2.00	0.012	0.002	0.411	0.3	0.2	0.05	—	0.05	0.2	—	998.6	515.3	293.1
W-2	0.40	2.01	1.99	0.012	0.002	0.761	0.3	0.2	0.05	—	0.05	0.2	—	1141.8	518.9	298.1
X-2	0.39	2.00	2.00	0.012	0.002	1.03	0.3	0.2	0.05	—	0.05	0.2	Ca: 0.004,	1250.3	520.7	302.5
													Mg: 0.005
Y-2	0.41	1.98	2.01	0.011	0.002	1.76	0.3	0.2	0.05	—	0.05	0.2	—	1537.2	514.4	292.7
*The balance consists of iron and inevitable impurities.

	TABLE 6
	Before processing	After processing
						Mean						Mean	Hydrochloric
						axis						axis	acid
	Steel					ratio of						ratio of	immersion	Exposure
Test	species	To	Residual γ	BF + M	Others	residual	TS	El	Residual γ	BF + M	Others	residual	test	test	DFL
No.	Symbol	° C.	%	%	%	γ	MPa	%	%	%	%	γ	h	Days	ratio
201	A-2	400	13	87	0	∘	1512	14	6	94	0	∘	Over 24	Over 30	—
202	B-2	400	11	89	0	∘	1233	16	4	96	0	∘	Over 24	Over 30	—
203	C-2	350	12	88	0	∘	1468	13	6	94	0	∘	Over 24	Over 30	1.00
204	D-2	350	11	88	1	∘	1510	13	4	96	0	∘	Over 24	Over 30	—
205	E-2	350	12	88	0	∘	1467	14	6	94	0	∘	Over 24	Over 30	—
206	F-2	350	12	88	0	∘	1491	13	5	95	0	∘	Over 24	Over 30	—
207	G-2	350	12	88	0	∘	1492	13	4	96	0	∘	Over 24	Over 30	—
208	H-2	350	12	88	0	∘	1505	14	4	96	0	∘	Over 24	Over 30	—
209	I-2	350	12	88	0	∘	1501	13	3	97	0	∘	Over 24	Over 30	—
210	J-2	350	13	87	0	∘	1461	14	5	95	0	∘	Over 24	Over 30	—
211	K-2	350	13	87	0	∘	1485	14	5	95	0	∘	Over 24	Over 30	—
212	L-2	350	13	87	0	∘	1495	13	3	97	0	∘	Over 24	Over 30	—
213	M-2	350	11	88	1	∘	1490	14	4	96	0	∘	Over 24	Over 30	—
214	N-2	350	14	86	0	∘	1478	12	5	95	0	∘	Over 24	Over 30	—
215	O-2	430	6	92	2	∘	1410	11	<1	99	<1	∘	9	7	—
216	P-2	400	<1	99	<1	x	1432	8	<1	99	<1	x	8	6	—
217	Q-2	—	<1	99	<1	x	1487	3	<1	99	<1	x	3	3	—
218	R-2	350	14	86	0	∘	1495	6	5	95	0	∘	17	8	—
219	S-2	350	11	88	1	∘	1448	8	3	97	0	∘	Over 24	14	—
220	A-2	350	9	20	71	x	961	14	15	85	0	x	Over 24	Over 30	—
221	T-2	400	11	89	0	∘	1490	14	7	93	0	∘	Over 24	Over 30	1.19
222	U-2	400	12	88	0	∘	1498	13	8	92	0	∘	Over 24	Over 30	1.49
223	V-2	400	12	88	0	∘	1509	14	8	92	0	∘	Over 24	Over 30	1.58
224	W-2	400	12	88	0	∘	1511	14	9	91	0	∘	Over 24	Over 30	1.58
225	X-2	400	12	88	0	∘	1503	13	8	92	0	∘	Over 24	Over 30	1.63
226	Y-2	400	15	55	30	x	1290	15	11	89	0	x	14	11	0.79

The results shown in Tables 5 and 6 can be interpreted as follows (numbers in the following description are test Nos. in Table 6).

Test pieces Nos. 201 through 214 (inventive steel sheets 2) and test pieces Nos. 221 through 225 (inventive steel sheets 1) that satisfy the requirements of the present invention have high strength of 1180 MPa or higher, and high hydrogen embrittlement resisting property in harsh environment after the forming process. They also have high elongation property required of the TRIP steel sheet, thus providing steel sheets best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere. Test pieces Nos. 221 through 225, in particular, show even better hydrogen embrittlement resisting property.

Test pieces Nos. 215 through 220 and 226 that do not satisfy the requirements of the present invention, in contrast, have the following drawbacks.

No. 215 made of steel species O-2 that includes insufficient C content has the amount of residual austenite significantly decreased after the processing, and fails to show the required level of hydrogen embrittlement resisting property of the present invention.

No. 216 made of steel species P-2 that includes insufficient Mn content does not retain sufficient residual austenite and is inferior in hydrogen embrittlement resisting property after the processing.

No. 217, martensite steel that is a conventional high strength steel made of steel species Q-2 that includes insufficient Si content, hardly contains residual austenite and is inferior in hydrogen embrittlement resisting property. It also does not show the elongation property required of a thin steel sheet.

No. 218 made of steel species R-2 that includes excessive C content has precipitation of carbide and is inferior in both the forming workability and the hydrogen embrittlement resisting property after processing.

No. 219 made of steel species S-2 that does not include Cu and/or Ni shows insufficient corrosion resistance and fails to show the required level of hydrogen embrittlement resisting property of the present invention.

No. 220, that was made of a steel that has the composition specified in the present invention but was not manufactured under the recommended conditions, resulted in the conventional TRIP steel. As a result, the residual austenite does not have the mean axis ratio specified in the present invention, while the matrix phase is not formed in binary phase structure of bainitic ferrite and martensite, and therefore sufficient level of hydrogen embrittlement resisting property is not achieved.

No. 226 includes Al content higher than that specified for the inventive steel sheet 1. As a result, although the predetermined amount of residual austenite is retained, the residual austenite does not have the mean axis ratio specified in the present invention, the desired matrix phase is not obtained and inclusions such as AlN are generated thus resulting in poor hydrogen embrittlement resisting property.

Then parts were made by using steel species A-2, K-2 shown in Table 5 and comparative steel sheet (590 MPa class high strength steel sheet of the prior art). Performance (pressure collapse resistance and impact resistance) of the formed test piece were studied by conducting pressure collapse test and impact resistance test as follows.

Pressure Collapse Test

Maximum tolerable load was determined similarly to Example 1 by using steel species A-2, K-2 shown in Table 5 and the comparative steel sheet. Absorbed energy was determined from the area lying under the load-deformation curve. The results are shown in Table 7.

	TABLE 7
	Evaluation
	of test piece
Steel sheet used	Maximum	Energy
	TS	EL	Residual γ	load	absorbed
Steel species	(MPa)	(%)	(Area %)	(kN)	(kJ)
Symbol A-2	1512	14	13	14.1	0.7
Symbol K-2	1485	14	13	13.9	0.68
Comparative steel	613	22	0	5.7	0.33
sheet

From Table 7, it can be seen that the part (test piece) made from the steel sheet of the present invention has higher load bearing capability and absorbs greater energy than a part made of the conventional steel sheet that has lower strength, thus showing higher pressure collapse resistance.

Impact Resistance Test

The impact resistance test was conducted similarly to Example 1 on the steel sheets made of steel species A-2, K-2 shown in Table 5 and the comparative steel sheet. The results are shown in Table 8.

TABLE 8
	Evaluation of test
Steel sheet used	piece
	TS	EL	Residual γ	Energy absorbed
Steel species	(MPa)	(%)	(Area %)	(kJ)
Symbol A-2	1512	14	13	7.06
Symbol K-2	1485	14	13	6.92
Comparative steel	613	22	0	3.56
sheet

From Table 8, it can be seen that the part (test piece) made from the steel sheet of the present invention absorbs greater energy than a part made of the conventional steel sheet having lower strength, thus showing higher impact resistance.

TEM photograph of the test piece made in this example is shown as reference. FIG. 8 is a photograph of TEM observation of No. 201 of the present invention. From FIG. 8, it can be seen that the high strength thin steel sheet of the present invention contains lath-shaped residual austenite (black portion of bar shape in FIG. 8) specified in the present invention dispersed therein. FIG. 9 is a photograph of TEM observation of No. 220 of a comparative example. From FIG. 9, it can be seen that the high strength thin steel sheet of No. 220 contains residual austenite (black portion of somewhat round shape in FIG. 9), although the residual austenite has a block shape that does not satisfy the requirements of the present invention.

Example 3

Sample steels A-3 through Q-3 having the compositions shown in Table 9 were melt-refined in vacuum to make test slabs. The slabs were processed in the following procedure (hot rolling→cold rolling→continuous annealing) thereby to obtain hot-rolled steel plates measuring 3.2 mm in thickness. The steel plates were pickled to remove scales from the surface and then cold rolled so as to reduce the thickness to 1.2 mm.

<Hot rolling> Starting temperature (SRT): Held at a temperature between 1150 and 1250° C. for 30 minutes.

Finishing temperature (FDT): 850° C.

Cooling rate: 40° C./s

Winding-up temperature: 550° C.

<Cold rolling> Rolling ratio: 50%

<Continuous annealing> Each steel specimen was kept at a temperature of A3 point+30° C. for 120 seconds, then rapidly cooled (air cooling) at a mean cooling rate of 20° C./s to temperature T0 shown in Table 10, and was kept at T0 for 240 seconds, followed by air-assisted water cooling to the room temperature.

No. 311 shown in Table 10 was made by heating a cold-rolled steel sheet to 830° C., keeping at this temperature for 5 minutes followed by quenching in water and tempering at 300° C. for 10 minutes, thereby to form a martensite steel as a comparative example of the high-strength steel of the prior art. No. 312 was made by heating a cold-rolled steel sheet to 800° C., keeping at this temperature for 120 seconds, cooling at a mean cooling rate of 20° C./s down to 350° C. and keeping at this temperature for 240 seconds.

JIS No. 5 test pieces were prepared from the steel sheets obtained as described above, and were subjected to stretch forming process with elongation of 3% mimicking the actual manufacturing process. Metal structures of the test pieces were observed before and after the processing, tensile strength (TS) and elongation (total elongation E1) before the processing and hydrogen embrittlement resisting property after the processing were measured by the following procedures.

Observation of Metal Structure

Metal structures of the test pieces were observed before and after the processing as follows. A measurement area (about 50 by 50 μm) at an arbitrarily chosen position in a surface parallel to the rolled surface at a position of one quarter of the thickness was photographed at measuring intervals of 0.1 μm, and area proportions of bainitic ferrite (BF), martensite (M) and residual austenite (residual γ) were measured by the method described previously. Then similar measurements were made in two fields of view that were arbitrarily selected, and the measured values were averaged. Area proportions of other structures (ferrite, pearlite, etc.) were subtracted from the entire structure.

Mean axis ratio, mean length of minor axes and minimum distance between the residual austenite grains of the steel sheet before and after the processing were measured by the method described previously. Test pieces having mean axis ratio of 5 or higher were regarded to satisfy the requirements of the present invention (∘), and those having mean axis ratio of lower than 5 were regarded to fail to satisfy the requirements of the present invention (x).

Measurement of tensile strength (TS) and elongation (E1) Tensile test was conducted on the JIS No. 5 test piece before processing, so as to measure the tensile strength (TS) and elongation (E1). Stretching speed of the tensile test was set to 1 mm/sec. Among the steel sheets having tensile strength of 1180 MPa as measured by the method described previously, those which showed elongation of 10% or more were evaluated as high in elongation.

Evaluation of Hydrogen Embrittlement Resisting Property

In order to evaluate the hydrogen embrittlement resisting property, flat test piece 1.2 mm in thickness was subjected to slow stretching rate test (SSRT) with a stretching speed of 1×10⁻⁴/sec, to determine hydrogen embrittlement risk index (%) defined by the equation shown below. Hydrogen embrittlement risk index (%)=100×(1−E1/E0)

E0 represents the elongation before rupture of a steel test piece that does not substantially contain hydrogen, E1 represents the elongation before rupture of a steel test piece that has been charged with hydrogen electrochemically in sulfuric acid. Hydrogen charging was carried out by immersing the steel test piece in a mixed solution of H₂SO₄ (0.5 mol/L) and KSCN (0.01 mol/L) and supplying constant current (100 A/m²) at room temperature.

A steel sheet having hydrogen embrittlement risk index higher than 50% is likely to undergo hydrogen embrittlement during use. In the present invention, steel sheets having hydrogen embrittlement risk index not higher than 50% were evaluated to have high hydrogen embrittlement resisting property.

Results of the test are shown in Table 10.

TABLE 9
Steel
species	Chemical composition (mass %)*	Ac3	Bs	Ms
Symbol	C	Si	Mn	P	S	Al	Cu	Ni	Ti	V	Nb	Ma	Others	(° C.)	(° C.)	(° C.)
A-3	0.29	1.45	1.50	0.014	0.002	0.043	—	—	—	—	—	—	—	847.5	616.7	374.0
B-3	0.40	2.03	2.01	0.012	0.002	0.033	0.3	0.2	—	—	—	—	—	826.4	533.7	301.7
C-3	0.42	1.99	2.53	0.011	0.002	0.033	—	0.2	—	—	—	—	—	809.4	481.5	275.0
D-3	0.39	2.02	2.51	0.011	0.002	0.030	—	—	0.05	0.05	—	—	—	843.1	498.8	293.3
E-3	0.41	1.98	2.48	0.011	0.002	0.031	—	—	—	—	0.05	0.2	—	820.5	479.5	280.6
F-3	0.39	2.00	2.53	0.011	0.002	0.030	—	—	—	—	—	—	B: 0.0005	816.4	497.0	292.7
G-3	0.35	1.50	2.50	0.012	0.002	0.033	0.3	0.3	0.05	0.05	0.05	0.2	B: 0.0005	824.5	482.8	303.3
H-3	0.55	1.52	1.53	0.012	0.002	0.033	0.3	0.3	0.05	—	0.05	0.2	Ca: 0.004,	818.8	516.1	240.5
													Mg: 0.005
I-3	0.40	1.98	2.48	0.011	0.002	0.033	—	—	—	—	—	—	REM: 0.005	816.6	498.8	289.6
J-3	0.7	1.51	1.52	0.011	0.005	0.031	—	—	—	—	—	—	—	782.2	504.2	179.0
K-3	0.39	0.14	1.42	0.012	0.002	0.043	—	—	—	—	—	—	—	772.5	596.9	329.3
L-3	0.29	1.49	1.51	0.011	0.002	0.052	—	—	—	—	—	—	—	850.5	615.8	373.7
M-3	0.30	1.49	1.50	0.012	0.002	0.241	0.3	0.2	—	—	—	—	—	916.2	606.6	365.9
N-3	0.31	1.51	1.52	0.012	0.002	0.387	0.3	0.3	0.05	—	—	—	—	991.5	598.4	358.8
O-3	0.31	1.51	1.49	0.012	0.002	0.740	0.3	0.3	0.05	—	—	—	—	1133.6	601.1	359.8
P-3	0.30	1.49	1.50	0.012	0.002	1.02	0.3	0.3	0.05	—	0.05	0.2	Ca: 0.004,	1252.6	586.3	360.0
													Mg: 0.005
Q-3	0.29	1.50	1.51	0.11	0.002	1.65	—	—	—	—	—	—	—	1490.1	615.8	373.7
*The balance consists of iron and inevitable impurities.

	TABLE 10
	Before processing
					Minimum
				Mean	distance
				axis	between	Mean
	Steel			ratio of	residual	axis
Test	species	To	Residual γ	residual γ	γ grains	ratio of	BF + M	Others	TS	El
No.	Symbol	° C.	%	nm	nm	residual γ	%	%	MPa	%
301	A-3	400	8.5	180	490	∘	91.5	0	1221	14
302	B-3	370	11	160	760	∘	89	0	1497	14
303	C-3	370	12	150	710	∘	88	0	1495	13
304	D-3	370	11	140	630	∘	89	0	1460	13
305	E-3	370	12	160	710	∘	88	0	1489	16
306	F-3	370	11	140	630	∘	89	0	1488	12
307	G-3	370	10	130	590	∘	90	0	1480	14
308	H-3	350	12	110	690	∘	87	1	1475	14
309	I-3	350	10	100	610	∘	89	1	1504	12
310	J-3	350	6	1200	700	∘	92	2	1422	7
311	K-3	350	<1	—	—	—	99	<1	1410	5
312	A-3	350	15	1500	610	x	20	65	967	15
313	L-3	400	9	190	750	∘	90	1	1230	14
314	M-3	370	10	160	700	∘	90	0	1410	13
315	N-3	370	10	150	620	∘	90	0	1471	12
316	O-3	370	10	150	620	∘	90	0	1471	12
317	P-3	350	11	140	610	∘	89	0	1480	11
318	Q-3	370	17	1100	1300	x	55	28	1320	14
	After processing
				Minimum
			Mean	distance
			axis	between	Mean			Hydrogen
			ratio of	residual	axis			embrittlement
	Test	Residual γ	residual γ	γ grains	ratio of	BF + M	Others	risk index
	No.	%	nm	nm	residual γ	%	%	%
	301	4	170	510	∘	96	0	30
	302	5	140	690	∘	95	0	25
	303	6	130	650	∘	94	0	23
	304	4	120	600	∘	96	0	20
	305	5	150	750	∘	95	0	25
	306	4	120	600	∘	96	0	23
	307	4	120	600	∘	96	0	13
	308	6	100	690	∘	94	0	15
	309	5	90	630	∘	95	0	18
	310	4	1180	800	∘	96	0	60
	311	<1	—	—	—	99	<1	90
	312	<1	1480	710	x	99	<1	85
	313	5	170	680	∘	95	0	24
	314	6	140	680	∘	94	0	22
	315	6	140	600	∘	94	0	18
	316	6	140	600	∘	94	0	18
	317	7	120	600	∘	96	0	15
	318	9	1000	1200	x	91	0	70

The results shown in Tables 9 and 10 can be interpreted as follows (numbers in the following description are test Nos. in Table 10).

Test pieces Nos. 301 through 309 (inventive steel sheets 2) and test pieces Nos. 313 through 317 (inventive steel sheets 1) that satisfy the requirements of the present invention have high strength of 1180 MPa or higher, and show high hydrogen embrittlement resisting property in harsh environment after the forming process. They also have high elongation property required of the TRIP steel sheet, thus providing steel sheets best suited for reinforcement parts of automobiles that are exposed to corrosive atmosphere.

Test pieces Nos. 310 through 312 and 318 that do not satisfy the requirements of the present invention, in contrast, have the following drawbacks.

No. 310 made of steel species J-3 that includes excessive C content has carbide precipitated and residual austenite of longer mean length of minor axis, thus resulting poor performance in both workability and hydrogen embrittlement resisting property after processing.

No. 311, martensite steel that is a conventional high strength steel made of steel species K-3 that includes insufficient Si content, hardly contains residual austenite and is inferior in hydrogen embrittlement resisting property. It also does not show the elongation property required of a thin steel sheet.

No. 312, that was made of a steel that has the composition specified in the present invention but was not manufactured under the recommended conditions, resulted in the conventional TRIP steel. As a result, the residual austenite does not have the mean axis ratio and the mean length of minor axis specified in the present invention, while the matrix phase is not formed in binary phase structure of bainitic ferrite and martensite, thus resulting in low strength and poor hydrogen embrittlement resisting property.

No. 318 includes Al content higher than that specified for the inventive steel sheet 1. As a result, although the predetermined amount of residual austenite is retained, the residual austenite does not have the mean axis ratio specified in the present invention, the desired matrix phase is not obtained and inclusions such as AlN are generated thus resulting in poor hydrogen embrittlement resisting property.

Then parts were made by using steel species A-3, G-3 shown in Table 9 and comparative steel sheet (590 MPa class high strength steel sheet of the prior art). Performance (pressure collapse resistance and impact resistance) of the formed test piece were studied by conducting pressure collapse test and impact resistance test as follows.

Pressure Collapse Test

Maximum tolerable load was determined similarly to Example 1 by using steel species A-3, G-3 shown in Table 9 and the comparative steel sheet. Absorbed energy was determined from the area under the load-deformation curve. The results are shown in Table 11.

	TABLE 11
	Evaluation
	of test piece
Steel sheet used	Maximum	Energy
	TS	EL	Residual γ	load	absorbed
Steel species	(MPa)	(%)	(Area %)	(kN)	(kJ)
Symbol A-3	1221	14	8.5	11.3	0.58
Symbol G-3	1480	14	10	13.8	0.69
Comparative steel	613	22	0	5.7	0.33
sheet

From Table 11, it can be seen that the part (test piece) made from the steel sheet of the present invention has higher load bearing capability and absorbs greater energy than a part made of the conventional steel sheet having lower strength, thus showing high pressure collapse resistance.

Impact Resistance Test

The impact resistance test was conducted similarly to Example 1 on the steel sheets made of steel species A-3, G-3 shown in Table 9 and the comparative steel sheet. The results are shown in Table 12.

TABLE 12
	Evaluation of test
Steel sheet used	piece
	TS	EL	Residual γ	Energy absorbed
Steel species	(MPa)	(%)	(Area %)	(kJ)
Symbol A-3	1221	14	8.5	5.72
Symbol G-3	1480	14	10	6.88
Comparative steel	613	22	0	3.56
sheet

From Table 12, it can be seen that the part (test piece) made from the steel sheet of the present invention absorbs greater energy than a part made of the conventional steel sheet that has lower strength, thus showing high impact resistance.

TEM photographs of the test pieces made in this example are shown as reference. FIG. 12 is a photograph of TEM observation (magnification factor 15000) of No. 301 of the present invention. FIG. 13 is a photograph of TEM observation (magnification factor 60,000) of a portion shown in the photograph of FIG. 12. From FIGS. 12, 13, it can be seen that the high strength thin steel sheet of the present invention contains fine residual austenite grains (black portion of bar shape in FIGS. 12, 13) specified in the present invention dispersed therein, and that the residual austenite has the lath shape specified in the present invention. FIG. 14 is a photograph of TEM observation of No. 313 of a comparative example. From FIG. 14, it can be seen that the high strength thin steel sheet of No. 313 contains residual austenite (black portion of somewhat round shape in FIG. 14), although the residual austenite has a block shape that does not satisfy the requirements of the present invention.

Claims

1. A high strength thin steel sheet having high hydrogen embrittlement resisting property and high workability, which comprises: C: higher than 0.25 up to 0.60%; Si: 1.0 to 3.0%; Mn: 1.0 to 3.5%; P: 0.15% or less; S: 0.02% or less; and Al: 1.5% or less (higher than 0%) in terms of percentage by weight, with balance of iron and inevitable impurities; wherein the metal structure after stretch forming operation with elongation of 3% comprises: residual austenite; 1% by area or more in proportion to the entire structure; bainitic ferrite and martensite: 80% or more in total; and ferrite and pearlite: 9% or less (may be 0%) in total, while the mean axis ratio (major axis/minor axis) of said residual austenite grains is 5 or higher, and the steel has tensile strength of 1180 MPa or higher.

2. The high strength thin steel sheet according to claim 1, wherein the metal structure after said stretch forming operation with elongation of 3% further satisfies the requirements that: mean length of minor axes of said residual austenite grains is 1 μm or less; and minimum distance between said residual austenite grains is 1 μm or less.

3. The high strength thin steel sheet according to claim 1, wherein 0.5% or less (higher than 0%) by weight of Al is contained.

4. The high strength thin steel sheet according to claim 1, wherein 0.003 to 0.5% of Cu and/or 0.003 to 1.0% of Ni in terms of percentage by weight are further contained.

5. The high strength thin steel sheet according to claim 1, wherein 0.003 to 1.0% of Ti and/or V in terms of percentage by weight are further contained.

6. The high strength thin steel sheet according to claim 1, wherein 1.0% or less (higher than 0%) of Mo and 0.1% or less (higher than 0%) of Nb in terms of percentage by weight are further contained.

7. The high strength thin steel sheet according to claim 1, wherein 0.0002 to 0.01% of B in terms of percentage by weight is further contained.

8. The high strength thin steel sheet according to claim 1, wherein at least one element selected from the group consisting of: 0.0005 to 0.005% of Ca; 0.0005 to 0.01% of Mg; and 0.0005 to 0.01% of REM in terms of percentage by weight is further contained.

9. A high strength thin steel sheet having high hydrogen embrittlement resisting property and high workability, which comprises: C: higher than 0.25 up to 0.60%; Si: 1.0 to 3.0%; Mn: 1.0 to 3.5%; P: 0.15% or less; S: 0.02% or less; and Al: 1.5% or less (higher than 0%) in terms of percentage by weight, with balance of iron and inevitable impurities, wherein the metal structure after stretch forming operation with elongation of 3% comprises: residual austenite; 1% by area or more in proportion to the entire structure; while the mean axis ratio (major axis/minor axis) of said residual austenite grains is 5 or higher; mean length of minor axes of said residual austenite grains is 1 μm or less; and minimum distance between said residual austenite grains is 1 μm or less; and the steel has tensile strength of 1180 MPa or higher.

10. The high strength thin steel sheet according to claim 9, wherein 0.5% or less (higher than 0%) by weight of Al is contained.

11. The high strength thin steel sheet according to claim 9, wherein 0.003 to 0.5% of Cu and/or 0.003 to 1.0% of Ni in terms of percentage by weight are further contained.

12. The high strength thin steel sheet according to claim 9, wherein 0.003 to 1.0% of Ti and/or V in terms of percentage by weight are further contained.

13. The high strength thin steel sheet according to claim 9, wherein 1.0% or less (higher than 0%) of Mo and 0.1% or less (higher than 0%) of Nb in terms of percentage by weight are further contained.

14. The high strength thin steel sheet according to claim 9, wherein 0.0002 to 0.01% of B in terms of percentage by weight is further contained.

15. The high strength thin steel sheet according to claim 9, wherein at least one element selected from the group consisting of: 0.0005 to 0.005% of Ca; 0.0005 to 0.01% of Mg; and 0.0005 to 0.01% of REM in terms of percentage by weight is further contained.