JP4268079B2 - Ultra-high strength steel sheet having excellent elongation and hydrogen embrittlement resistance, method for producing the same, and method for producing ultra-high strength press-formed parts using the ultra-high strength steel sheet - Google Patents

Description

  The present invention is an ultra-high strength steel plate excellent in elongation (that is, total elongation) and hydrogen embrittlement resistance in an ultra-high strength region of 1180 MPa class or higher, a method capable of efficiently producing the ultra-high strength steel plate, In addition, the present invention relates to a method for manufacturing an ultra-high-strength press-formed part using the ultra-high-strength steel sheet.

  Steel plates used by press forming in automobiles and industrial machines are required to have both excellent strength and ductility, and in recent years, the need for ultra-high strength steel plates of 1180 MPa class or higher has been increasing. .

  Until now, such ultra-high strength steel plates have often been used for reinforcing members obtained by simple bending, such as bumpers, door impact bars and beams. However, recently, from the viewpoint of reducing the weight of the vehicle body in order to achieve both collision safety and environmental issues, it is applied to members that require complex press molding, such as front side members, rear side members, lockers, and pillars. Attempts to do so are increasing, and improvement in ductility in the ultra-high strength steel sheet is desired.

  However, it is known that a new adverse effect such as delayed fracture (cracking, etc.) due to hydrogen embrittlement occurs in the ultrahigh strength region of 1180 MPa class or higher. Delayed fracture is a high-strength steel in which hydrogen generated from a corrosive environment or atmosphere diffuses into defects such as dislocations, vacancies, and grain boundaries, embrittles the material, and breaks when stress is applied. This is a phenomenon, and as a result, it has adverse effects such as a decrease in the ductility and toughness of the metal material.

  Therefore, even in an ultra-high strength steel plate of 1180 MPa class or higher, it is a matter of great concern that the steel plate suddenly breaks due to hydrogen generated by a corrosion reaction in the atmospheric environment.

  Therefore, in order to improve the hydrogen embrittlement resistance in the ultra-high strength steel sheet, Patent Document 1 discloses that Ca is added to increase the bonding strength of austenite crystal grains, or Ti is added to refine crystal grains and crystal grains. A method for producing an ultra-high strength steel sheet that suppresses the growth of steel is disclosed. Although the hydrogen embrittlement resistance is improved by this method, the crack generation time by the cathode charge test is 1000 seconds or less, and further improvement is required. Further, since the steel sheet obtained by the above method is mainly composed of tempered martensite, some examples have an elongation of only about 10%. There is an urgent need to improve the crystallization characteristics and achieve excellent ductility.

On the other hand, TRIP (Transformation Induced Plasticity) steel sheets have attracted attention as high-strength steel sheets having excellent elongation and delayed fracture resistance. The TRIP steel sheet has a retained austenite structure, and when deformed at a temperature equal to or higher than the martensite transformation start temperature (Ms point), the retained austenite (γ _R ) is induced and transformed into martensite by stress, resulting in a large elongation. For example, Patent Document 2 discloses a high-tensile cold-rolled steel sheet having fine ferrite (polygonal ferrite), fine martensite, and 3-20% retained austenite (γ _R ). Yes. This is what with improved resistance to delayed fracture by refinement of ferrite and martensite, aimed elongation improving effect due to induced transformation by the formation of gamma _R. However, in the above publication, since the upper limit is set to 20% based on the reason that γ _R has a high hydrogen absorption capacity and significantly deteriorates the delayed fracture resistance, the effect of improving the elongation by γ _R is effectively exhibited. First, the elongation property of the 1180 MPa class ultra-high strength steel sheet is only 17%. As described above, in TRIP steel sheet (PF steel sheet) having polygonal ferrite as a matrix, γ _R deteriorates the delayed fracture resistance. The reason is that γ _R exists at the triple point of the grain boundary, so the stability is poor. It can be considered that the martensitic transformation is easily caused by the strain-induced transformation described above.

It is also described in Non-Patent Document 1 that γ _R adversely affects delayed fracture characteristics. According to this, it is reported that the ultra high strength cold-rolled steel sheet is more likely to cause delayed fracture as the γ _R is larger and the strain is larger, and the influence of the processing is also pointed out. γ _R transforms into martensite by processing, but the solid solution amount of hydrogen in the martensite phase is smaller than the austenite phase, so hydrogen accumulates and molecularizes at the interface between the austenite phase and the martensite phase, increasing the internal pressure, Alternatively, it is presumed that delayed fracture occurs due to, for example, the accumulation of hydrogen promoting processing-induced transformation.

Recently, however, it has been pointed out that γ _R is not necessarily a factor that degrades delayed fracture characteristics. For example, in Non-Patent Document 2, when a fracture surface of a TRIP steel sheet (TAM steel sheet) having tempered martensite as a matrix is observed, pseudo-cleavage fracture due to hydrogen storage is suppressed; It has been reported that the reduction in total elongation after hydrogen storage is suppressed to a small level compared to martensitic single-phase steel sheets not containing _R. The above TAM steel sheet has fine γ _R generated between the laths, and as compared to γ _R present in the PF steel sheet, the stability against strain-induced transformation is increased and the toughness is improved. It is thought that it can be prevented. Furthermore, due to the generation of the fine γ _R described above, the ductility is remarkably improved as compared with the conventional PF steel sheet.

However, while the present invention aims to improve delayed fracture resistance in an ultrahigh strength region of 1180 MPa class or higher, the strength of the TAM steel sheet described in Non-Patent Document 2 is as low as 1030 MPa, and the TAM Even if it is a steel plate, it still does not satisfy the required characteristics of the present invention.
JP-A-8-134549 (Claims, [0011], [0013], etc.) JP 2001-81533 A (Claims, [0020] to [0024], Table 3 etc.) Kazumasa Yamazaki, Yaichiro Mizuyama, “Effects of retained austenite and strain on delayed fracture properties of ultra-high-strength cold-rolled steel sheets”, Iron and Steel, 1997, Vol. 83, No. 11, p.66-71 Tomohiko Kitajo, 5 others, "Hydrogen embrittlement of ultra-high strength low alloy TRIP steel (1st report: Hydrogen storage properties and ductility)", Proceedings of the 51st Annual Conference of the Society of Materials Science, 2002, Vol. 8, p.17-18

  The present invention has been made paying attention to the above circumstances, and its purpose is to provide an ultra-high strength steel sheet having excellent elongation and hydrogen embrittlement resistance in an ultra-high strength region of 1180 MPa class or higher. An object of the present invention is to provide a method capable of efficiently producing a high-strength steel sheet and a method for producing an ultra-high-strength press-formed part using the ultra-high-strength steel sheet.

The ultra-high strength steel sheet of the present invention excellent in elongation and hydrogen embrittlement resistance in an ultra-high strength region with a tensile strength of 1180 MPa or more that can solve the above-mentioned problems is expressed by mass%.
C: 0.06 to 0.6%,
Si + Al: 0.5 to 3%
Mn: 0.5-3%,
P: 0.15% or less (excluding 0%),
S: 0.02% or less (including 0%)
And containing
The organization is the space factor for the whole organization,
15-60% tempered martensite,
5-50% of ferrite,
It contains 15 to 45% of massive martensite with a retained austenite of 5% or more and an aspect ratio of 3 or less, and the space factor occupied by fine martensite with an average particle size of 5 μm or less in the massive martensite is 30 It has a gist where it is at least%.

Furthermore, in the present invention, by mass%,
(I) Mo: 1% or less (not including 0%), Ni: 0.5% or less (not including 0%), Cu: 0.5% or less (not including 0%), Cr: 1% Containing at least one of the following (not including 0%);
(Ii) At least one of Ti: 0.1% or less (not including 0%), Nb: 0.1% or less (not including 0%), V: 0.1% or less (not including 0%) Containing:
(Iii) Ca: 0.003% or less (excluding 0%) and / or REM: 0.003% or less (excluding 0%):
(Iv) A case in which the amount of dissolved carbon in the retained austenite is 0.85% or more is a preferable embodiment.

Furthermore, the manufacturing method of the ultra-high strength steel sheet according to the present invention that can solve the above-mentioned problems is as follows.
(I) at least a step of heating and holding a steel satisfying the above components to a temperature of A ₃ point or more and 1100 ° C. or less for 10 seconds or more and then cooling to a temperature of Ms point or less at an average cooling rate of 30 ° C./s or more. And (ii) (A ₃ point-25 ° C.) to A ₃ point at a temperature of 120 to 600 seconds after heating and holding, at an average cooling rate of 3 ° C./s or higher, Ms point or higher and Bs point It has a gist in that it includes a step of cooling to the following temperature and holding in the temperature range for 1 second or longer.

  In the above ultra-high strength steel sheet, particularly when producing a steel sheet in which the amount of dissolved carbon in the retained austenite satisfies 0.85% or more, in the step (ii), the Ms point or more and the Bs point or less are used. It is recommended to cool to a temperature and keep the holding time in the temperature range at least 200 seconds.

  The present invention also includes a method for producing an ultra-strength press-molded part by warm press-molding an ultra-high strength steel sheet having a solid solution carbon content of 0.85% or more in retained austenite.

  Since the present invention is configured as described above, in an ultrahigh strength region of 1180 MPa class or higher, an ultrahigh strength steel plate having excellent elongation and hydrogen embrittlement resistance, and the ultrahigh strength steel plate are press-molded. An ultra-high-strength press-formed product could be produced efficiently. In particular, ultra-high-strength press-molded parts made by warm press-forming steel sheets that have a dissolved carbon content of 0.85% or more in retained austenite ensure excellent delayed fracture resistance because the retained austenite is stabilized. can do.

  In view of the fact that the present inventors have not yet obtained an ultra-high-strength steel sheet having excellent elongation and hydrogen embrittlement resistance in an ultra-high strength region of 1180 MPa class or higher, A method capable of efficiently producing an ultra-high strength steel sheet has been intensively studied based on the TAM steel sheet described above. This TAM steel plate is one of the new TRIP steel plates developed by the present inventors, and has already been filed as a low alloy TRIP steel plate having a large total elongation while maintaining high stretch flangeability. (Japanese Patent Application No. 2002-54595, hereinafter referred to as the prior application invention).

The details of the TAM steel sheet are as described in the specification of the invention of the prior application.
(I) A mixed structure of tempered martensite and ferrite composed of a soft lath structure with low dislocation density is used as a parent phase to improve workability (stretch flangeability and total elongation);
(Ii) As the second phase, by controlling to a retained austenite (γ _R ) structure having a high C concentration (Cγ _R ), preferably lath shape, the total elongation is particularly improved. From this research, the steel sheet was found to be excellent in delayed fracture resistance, and was published in Non-Patent Document 2.

  Thus, the TAM steel sheet is very useful in that it has not only excellent elongation and stretch flangeability, but also excellent delayed fracture resistance in a high strength region of about 1030 MPa. However, the TAM steel sheet disclosed in the invention of the prior application is not intended to provide an ultra-high-strength steel sheet that is the subject of the present invention. In the ultra-high-strength region of 1180 MPa class or higher, elongation characteristics and resistance There is an urgent need to provide ultra-high strength with excellent delayed fracture properties.

Therefore, the present inventors, in providing an ultra-high-strength steel sheet having all the desired characteristics, the “elongation improving effect by the combination of tempered martensite, ferrite and γ _R ” disclosed in the prior application invention is In order to ensure delayed fracture resistance in the desired ultra-high strength region while following it as it is, paying attention to martensite, which was not considered in the prior invention, particularly by changing the aspect ratio and particle size. I came. When a basic experiment was conducted focusing on the martensitic structure of the prior invention, it was difficult to ensure excellent delayed fracture resistance in the ultra-high strength region when coarse martensite having an average particle size exceeding 5 μm was generated. This is because it was proved to be.

As a result, it has been clarified that it is necessary to generate a large amount of fine massive martensite in order to ensure delayed fracture resistance in the ultrahigh strength region of 1180 MPa class or higher. In particular,
(I) In order to ensure a desired ultra-high strength level, annealing in the two-phase region is carried out at a high temperature for a long time to generate massive martensite (hard martensite);
(Ii) In order to ensure good delayed fracture resistance in the ultra-high strength region, a specific quenching process is repeated at least twice to generate fine martensite. By
The present invention was completed by finding out that the intended purpose was achieved.

In other words, the ultra-high strength steel sheet of the present invention is the structure disclosed in the prior invention,
(I) “A tempered martensite composed of a soft lath structure having a low dislocation density, ferrite, and a predetermined amount of γ _R having a high C concentration (Cγ _R ) and preferably having a lath shape” are generated. While ensuring excellent elongation characteristics in the above ultra-high strength region,
(Ii) To ensure the “ultra-high strength level of 1180 MPa class or higher” that could not be achieved in the prior invention, and to ensure “excellent delayed fracture resistance in the ultra-high strength region” The most important point exists when a large number of “fine massive martensite” that could not be obtained by the method of the invention of the prior application was generated.

  As described above, the present invention has not yet been achieved in the field of steel sheets. “Elongation characteristics (over 18% in total elongation) and hydrogen embrittlement resistance (measured by the method described later) in an ultrahigh strength region of 1180 MPa class or higher. In this case, the cathode CH life is 1000 seconds or more) and the present invention is based on the unique solution of “providing an ultra-high-strength steel sheet.” By adopting the above-described unique heat treatment method, It has the greatest characteristics when a tissue useful for achieving the above object is generated.

  Hereinafter, the ultra high strength steel sheet of the present invention will be described.

  First, the organization that most characterizes the present invention will be described.

15-60% space factor for tempered martensite
The “tempered martensite” in the present invention has the following characteristics.

First, “tempered martensite” in the present invention means a soft material having a low dislocation density and having a lath-like structure. On the other hand, martensite is different from the tempered martensite in that it is a hard structure having a high dislocation density, and both can be distinguished by observation with a transmission electron microscope (TEM), for example. The conventional γ _R steel sheet is also different from the steel sheet of the present invention having the tempered martensite as a parent phase in that it has a soft block ferrite structure with a low dislocation density.

  Second, the tempered martensite tends to have a generally higher Vickers hardness (Hv) than polygonal ferrite in the same component system (system in which the basic components C, Si, and Mn are the same). . FIG. 1 shows tempered martensite in the same component steel grade (C: 0.1 to 0.3%, Mn: 1.0 to 2.0%, Si: 1.0 to 2.0%). It is the graph which contrasted the hardness (vertical axis) of tempered bainite, and the hardness (horizontal axis) of polygonal ferrite. Incidentally, the Vickers hardness is obtained by measuring the Vickers hardness (Hv) of the parent phase (gray) part by observation with an optical microscope by Repeller corrosion (load 1 g). For reference, the straight line y = x is shown as a dotted line in the figure, and as a result, the hardness of tempered martensite is higher than that of polygonal ferrite; such a tendency increases as the hardness increases. , You can see that it is noticeable.

  In addition, FIG. 2 is an arrangement of the data of FIG. 1 divided into each case of C amount: 0.1%, 0.2%, and 0.3%, tempered martensite, tempered bainite, And the influence of the amount of C on the hardness of polygonal ferrite. FIG. 2 shows that when the C content is the same, the hardness of tempered martensite tends to be higher than that of polygonal ferrite: such a tendency can be seen more prominently as the C content increases. I understand.

  Based on these results, when the hardness in tempered martensite, tempered bainite, and polygonal ferrite is expressed in relation to the basic components of C, Mn, and Si, the following relational expression is generally obtained.

Hardness of tempered martensite and tempered bainite (Hv)
≧ 500 [C] +30 [Si] +3 [Mn] +50
Polygonal ferrite hardness (Hv) ≒ 200 [C] +30 [Si] +3 [Mn] +50
In the formula, [] means the content (% by mass) of each element.

  Incidentally, it has been confirmed that the hardness (calculated value) obtained by the above relational expression substantially reflects the actually measured value.

  Further, the hardness obtained by the above relational expression is not only when the C amount is 0.1 to 0.3%, but also when it is 0.3 to 0.6%, further 0.06 to 0.1%. In some cases, it is confirmed that the measured values are reflected in the same manner.

  The upper limit of the tempered martensite hardness may vary depending on the component composition, etc., but is generally 500 [C] +30 [Si] +3 [Mn] +200, preferably 500 [C] +30 [Si ] +3 [Mn] +150 is recommended.

In order to effectively exhibit the effect of improving the elongation by the tempered martensite, it is necessary to have 15% or more of tempered martensite in a space factor with respect to the entire structure. The specific space factor of tempered martensite is determined by the balance with ferrite and γ _R described later, and it is recommended that the space ratio be appropriately controlled so that desired characteristics can be exhibited. However, if the space factor of tempered martensite exceeds 60%, the desired strength cannot be obtained, so the upper limit is made 60%. Preferably it is 55% or less, More preferably, it is 50% or less.

5-50% space factor for ferrite in all structures
The “ferrite” in the present invention means polygonal ferrite, that is, ferrite having a low dislocation density.

  In order to effectively exhibit the action according to the present invention, it is recommended that ferrite is present in an amount of 5% or more. Preferably it is 10% or more. In particular, from the viewpoint of improving the elongation characteristics, the amount of ferrite is preferably large, and it is recommended that the ferrite content be 30% or more, more preferably 40% or more. However, if it exceeds 50%, it becomes difficult to ensure the required ultra-high strength, so the upper limit is made 50%. Preferably it is 40% or less, More preferably, it is 30% or less.

Residual austenite (γ _R ) has a space factor of 5% or more with respect to the entire structure. Γ _R is useful for improving the total elongation. It is necessary that the volume ratio be 5% or more (preferably 7% or more). On the other hand, if it is present in a large amount, the desired ultra-high strength cannot be ensured, so it is recommended that the upper limit be 30%. More preferably, it is 25%.

In the present invention, the form of γ _R is preferably a lath shape. Here, “the shape is lath” means that the average axial ratio (major axis / minor axis) is 2 or more (preferably 4 or more, more preferably 6 or more). Such a lath-like γ _R is extremely useful in that it exhibits not only the TRIP effect similar to that of the conventional γ _R but also a significant stretch flangeability improvement effect. Although the upper limit of the average axial ratio is not particularly defined, the preferable upper limit is 30 in view of the fact that a certain thickness of γ _R is necessary in order to effectively exert the effect of TRIP. Preferably it is 20.

In addition, in order to effectively exhibit the effect of the above-mentioned lath γ _R, the space factor of the lath-shaped γ _R occupied in γ _R is the more the better. Specifically, it is determined by the balance with the above-mentioned tempered martensite and ferrite, and it is recommended to control appropriately so that the desired characteristics can be exhibited, but from the viewpoint of improving the strength. For example, the space factor of the lath-like γ _R may be 50% or more, more preferably 60% or more, still more preferably 70% or more, still more preferably 80% or more, and still more preferably 85% or more. Recommended. Incidentally, all the gamma _R is may be composed of a lath gamma _R, considering the limitations or the like of the heating equipment and cooling equipment, at a practical level, it is recommended that the upper limit of about 95%.

Further solute carbon concentration in the γ _R (Cγ _R) is recommended to be 0.8% or more. This C gamma _R is largely affects the characteristics of the TRIP (strain-induced transformation process), controlling over 0.8%, in particular, is effective in improving the elongation and the like. Preferably it is 1% or more, More preferably, it is 1.2% or more. Although preferred as the content of the C gamma _R is large, the actual operation, adjustable upper limit is believed to roughly 1.6%.

When the ultra-high-strength steel sheet of the present invention is press-molded into an ultra-high-strength press-molded part, the Cγ _R range is 0.85% or more (preferably 0.87% or more, more preferably 0). controls over .89%) suppressed as much as possible the transformation to martensite stabilized thereby (gamma _R a gamma _R in) it is recommended. Thereby, even if it is a case where it is an ultra-high-strength press-molded part, it is possible to ensure excellent delayed fracture resistance. Although preferred as the content of the C gamma _R is large, the actual operation, adjustable upper limit is believed to roughly 1.6%.

This point will be described in more detail. As reported in the above-mentioned Non-Patent Document 1 and the like, it is known that γ _R is likely to cause delayed fracture, and the greater the strain, the delayed fracture occurs. Since the steel sheet of the present invention contains a large amount of “fine massive martensite” effective in improving delayed fracture resistance, the above-mentioned problems due to γ _R are avoided at the stage of the steel sheet, When a molded part (molded member) is formed by press working, the above-mentioned problems become obvious, and γ _R may be transformed into martensite (work-induced transformation) by processing, and the hydrogen embrittlement resistance may deteriorate. There is. Therefore, in the present invention, it was to stabilize the gamma _R by the amount of C gamma _R as high as 0.85% or more. Furthermore, at the stage of processing, it is recommended to press-form (warm processing) the steel plate with stabilized γ _R in particular at a warm temperature (about 50 to 400 ° C.). In general, warm processing is said to be easier to suppress transformation of γ _R than room temperature processing, so it is useful to provide a means to further stabilize γ _R not only in the steel plate stage but also in the processing stage. As a result of the combined use of these, the fact that the martensitic transformation of γ _R is minimized, and as a result, an ultra-high-strength press-molded part with further improved hydrogen embrittlement resistance can be obtained is confirmed by Examples described later. ing.

It contains 15 to 45% of bulk martensite (aspect ratio is 3 or less) with respect to the entire structure, and the space ratio occupied by fine martensite (average particle size of 5 μm or less) in the bulk martensite is The above requirement is 30% or more . In short, a large amount of fine massive martensite is generated, and thereby, excellent delayed fracture resistance in an ultrahigh strength region of 1180 MPa class or higher is ensured. Since the action of the martensite that characterizes the present invention most differs depending on whether it is “bulky” or “fine”, it will be described separately below.

  First, “bulky martensite” in the present invention means an aspect ratio (major axis / minor axis) of 3 or less, and a structure useful for securing a desired ultrahigh strength level (1180 MPa class or more). It is. As the aspect ratio becomes smaller and approaches 1, the shape of martensite becomes a block shape, and the strengthening effect is more effectively exhibited. Preferably it is 2.5 or less, More preferably, it is 2 or less.

  In order to effectively exhibit the ultra-high strength action by the massive martensite, it is necessary to contain 15% in space factor with respect to the whole structure. Preferably it is 20% or more, More preferably, it is 25% or more. However, if the mass ratio of the massive martensite exceeds 45%, the desired elongation characteristics and further the desired cathode CH life cannot be obtained, so the upper limit is made 45%. Preferably it is 40% or less, More preferably, it is 38% or less.

  Furthermore, in this invention, the space factor which fine martensite (average particle diameter is 5 micrometers or less) accounts for 30% or more in the said massive martensite. This is because generation of fine martensite can prevent hydrogen embrittlement in the ultra-high strength region and improve delayed fracture resistance. In order to effectively exhibit the delayed fracture resistance improvement effect by such fine martensite, the space factor with respect to massive martensite is preferably 35% or more, more preferably 40% or more. The upper limit is not particularly limited, and the higher the space factor of fine martensite, the better.

Other: Bainite (including 0%)
The steel sheet of the present invention may be composed of only the above structure (that is, a mixed structure of tempered martensite, ferrite, γ _R , and massive martensite), but in a range not impairing the function of the present invention, You may have bainite. Bainite can inevitably remain in the production process of the present invention, but the smaller the bainite, the better.

  Next, basic components constituting the steel plate of the present invention will be described. Hereinafter, all the units of chemical components are mass%.

C: 0.06 to 0.6%
C is an essential element for securing high strength and securing γ _R. More specifically, it is an important element for containing a sufficient amount of C in the γ phase and leaving the desired γ phase even at room temperature, and is useful for increasing the strength-elongation balance. In particular, when the amount of C is added by 0.25% or more, the amount of γ _R increases and the concentration of C to γ _R increases, so that an extremely high strength-elongation can be obtained.

  However, if added over 0.6%, not only the effect is saturated, but also defects due to center segregation during casting and the like are observed. Moreover, when it adds 0.25% or more, weldability will deteriorate.

  Therefore, considering mainly weldability, it is preferable to control C: 0.06 to 0.25% (more preferably 0.2% or less, still more preferably 0.15% or less), When high elongation or the like is required without requiring spot welding, it is recommended to control C: 0.25 to 0.6% (more preferably 0.3% or more).

Si + Al: 0.5 to 3%
Si and Al are elements that effectively suppress the decomposition of γ _R and the formation of carbides. In particular, Si is useful as a solid solution strengthening element. In order to effectively exhibit such an effect, it is necessary to add 0.5% or more in total of Si and Al. Preferably it is 0.7% or more, More preferably, it is 1% or more. However, even if the above elements are added in total exceeding 3%, the above effects are saturated, and it is economically wasteful. If added in a large amount, hot brittleness is caused, so the upper limit is 3%. And Preferably it is 2.5% or less, More preferably, it is 2% or less.

Mn: 0.5 to 3%
Mn is an element necessary for stabilizing γ and obtaining a desired γ _R. In order to exhibit such an action effectively, it is necessary to add 0.5% or more. Preferably it is 0.7% or more, More preferably, it is 1% or more. However, if it is added in excess of 3%, adverse effects such as slab cracking are observed. Preferably it is 2.5% or less, More preferably, it is 2% or less.

P: 0.15% or less (excluding 0%)
P is an element effective for securing desired γ _R. In order to effectively exhibit such an action, it is recommended to add 0.03% or more (more preferably 0.05% or more). However, when it exceeds 0.1%, secondary workability deteriorates. More preferably, it is 0.1% or less.

S: 0.02% or less (including 0%)
S is an element that forms sulfide inclusions such as MnS and degrades workability as a starting point of cracking. Preferably it is 0.02% or less, More preferably, it is 0.015% or less. Note that the workability deterioration suppression effect due to the reduction of S is saturated when S is reduced to 0.003% or less, and conversely, considering that the cost for reducing S is high, the lower limit is 0. It is recommended to exceed 0.003%, more preferably 0.005% or more.

  The steel of the present invention basically contains the above components, and the balance: substantially iron and impurities, but the following allowable components can be added as long as the effects of the present invention are not impaired.

Mo: 1% or less (not including 0%), Ni: 0.5% or less (not including 0%), Cu: 0.5% or less (not including 0%), Cr: 1% or less (0 These elements are not only useful as steel strengthening elements, but also effective for stabilizing γ _R and securing a predetermined amount. In order to effectively exhibit such an action, Mo: 0.05% or more (more preferably 0.1% or more), Ni: 0.05% or more (more preferably 0.1% or more), Cu : 0.05% or more (more preferably 0.1% or more) and Cr: 0.05% or more (more preferably 0.1% or more) are recommended to be added. However, even if Mo and Cr are added in excess of 1% and Ni and Cu are added in excess of 0.5%, the above effect is saturated, which is economically wasteful. More preferably, Mo is 0.8% or less, Ni is 0.4% or less, Cu is 0.4% or less, and Cr is 0.8% or less.

At least one of these elements : Ti: 0.1% or less (not including 0%), Nb: 0.1% or less (not including 0%), V: 0.1% or less (not including 0%) Has an effect of precipitation strengthening and refinement of structure, and is an element useful for increasing the strength. In order to effectively exhibit such an action, Ti: 0.01% or more (more preferably 0.02% or more), Nb: 0.01% or more (more preferably 0.02% or more), V : 0.01% or more (more preferably 0.02% or more) is recommended to be added respectively. However, if any element is added in excess of 0.1%, the above effect is saturated, which is economically useless. More preferably, Ti is 0.08% or less, Nb is 0.08% or less, and V is 0.08% or less.

Ca: 0.003% or less and / or REM: 0.003% or less
(Excluding 0%)
Ca and REM (rare earth elements) are elements that control the form of sulfide in steel and are effective in improving workability. Here, examples of rare earth elements used in the present invention include Sc, Y, and lanthanoids. In order to effectively exhibit the above action, it is recommended to add 0.0003% or more (more preferably 0.0005% or more). However, even if added over 0.003%, the above effect is saturated, which is economically useless. More preferably, it is 0.0025% or less.

  Next, a method for producing the steel sheet of the present invention will be described.

In the method for producing a steel sheet of the present invention, steel satisfying the above components is heated and held at a temperature of A ₃ point or higher and 1100 ° C. or lower for 10 seconds or more, and then at an average cooling rate of 30 ° C./s or higher to a temperature of Ms point or lower. A step including at least two cooling steps, and after heating and holding at a temperature of (A ₃ point −25 ° C.) to A ₃ point for 120 to 600 seconds, at an average cooling rate of 3 ° C./s or higher, an Ms point or higher It is characterized in that it includes a step of cooling to a temperature below the Bs point and holding at that temperature range for 1 second or longer (particularly preferably 200 seconds or longer).

As described above, the ultra-high strength steel sheet of the present invention is the structure disclosed in the prior application invention,
(I) “A tempered martensite composed of a soft lath structure having a low dislocation density, ferrite, and a predetermined amount of γ _R having a high C concentration (Cγ _R ) and preferably having a lath shape” are generated. While ensuring excellent elongation characteristics in the above ultra-high strength region,
(Ii) To ensure the “ultra-high strength level of 1180 MPa class or higher” that could not be achieved in the prior invention, and to ensure “excellent delayed fracture resistance in the ultra-high strength region” The most important point exists where a large number of "fine massive martensite" that could not be obtained by the method of the invention of the prior application exists, and in order to obtain such an ultra-high strength steel sheet, There are two typical methods described in the invention of the prior application;
(1) [Hot rolling process] → [Cold rolling process] → [First continuous annealing process]
→ [Second continuous annealing process or plating process]
(2) [Hot rolling process] and [Continuous annealing process or Plating process]
Based on the above, the following modifications (a) and (b) are further made.

(A) In order to generate a predetermined amount of massive martensite (hard martensite) that could not be obtained by the method of the prior invention, the [second continuous annealing step] of (1) above, In the continuous annealing step or plating step], annealing in a two-phase region is performed at a substantially higher temperature for a longer time than in the invention of the prior application, and (b) by refining the massive martensite to obtain fine martensite. In order to generate a predetermined amount, the [first continuous annealing step] in (1) is a step in which the rapid heating and quenching treatment is repeated at least twice; or in the [hot rolling step] in (2) above After the rapid heating and quenching process is performed once, cold rolling is performed as necessary, and then the rapid heating and quenching process is performed at least once in the [continuous annealing process or plating process] of (2).

  In the following, description will be given separately for each case.

(1) [Hot rolling process] → [Cold rolling process] → [First continuous annealing process] → [Second continuous annealing process or plating process]
The method (1) is a method for producing a desired steel sheet through a hot rolling step, a cold rolling step, a first continuous annealing step, and a second continuous annealing step or a plating step.

  First, a hot rolling process and a cold rolling process are performed, but these processes are not particularly limited, and usually, conditions to be performed can be appropriately selected and employed. In the method (1), the desired structure is not secured by the hot rolling process and the cold rolling process, but the first continuous annealing process and the second continuous annealing process or plating process to be performed thereafter are performed. This is because the desired structure is obtained by controlling.

Specifically, as the hot rolling step, conditions such as cooling at an average cooling rate of about 30 ° C./s and winding at a temperature of about 500 to 600 ° C. are adopted after the hot rolling is completed at _{Ar 3} or higher. be able to. In the cold rolling process, it is recommended to perform cold rolling at a cold rolling rate of about 30 to 70%. Of course, this is not intended to be limited to this.

  Next, (1-1) the first continuous annealing step and (1-2) the second continuous annealing step or plating step characterizing the method (1) will be described with reference to FIG. FIG. 3 shows a process chart in which the predetermined rapid heating and quenching process is performed twice in the first continuous annealing process, but it is of course not intended to be limited to this.

(1-1) First Continuous Annealing Step The above-described step is: “A ₃ point or more and 1100 ° C. or less (T1 in FIG. 3) heated and held for 10 seconds or more (t1 in FIG. 3), then 30 ° C. The process of “cooling to a temperature below the Ms point (T2 in FIG. 3) at an average cooling rate (CR1 in FIG. 3) of at least / s” (hereinafter, these series of methods are referred to as “rapid thermal quenching treatment”). In some cases) at least twice. By the above process, martensite can be secured after cooling while avoiding pearlite transformation, and it becomes possible to refine the old γ grain size to 20 μm or less, which is finally useful for improving delayed fracture resistance. This is because fine martensite is obtained.

Here, the heating temperature of T1 is a A ₃ point or higher 1100 ° C. or less. This is because the austenite grain growth can be prevented and the prior γ grain size can be refined by adjusting to the above range. The lower limit of the preferred heating temperature may vary depending on the components in the steel, but is generally preferably 880 ° C. or higher. This is because, even if the temperature is controlled to ₃ points or more, if the heating temperature is low, undissolved carbides are generated, and it is somewhat difficult to refine the old γ grains. On the other hand, if the heating temperature exceeds 1100 ° C., γ grains grow and the old γ grains cannot be refined. In addition, it is recommended to control the temperature to 980 ° C. or less in consideration of the loss of economic efficiency, for example, as the heating temperature increases, the equipment needs to be enlarged accordingly.

  The soaking time t1 at the heating temperature is 10 to 300 seconds. This is because if the soaking time is short, austenite is not sufficiently formed and cementite and other alloy carbides remain, so it is difficult to make the old γ grains fine. The soaking time t1 is preferably as long as possible, and is preferably 60 seconds or longer, more preferably 120 seconds or longer. However, if the length is too long, growth of old γ grains is caused and the production efficiency is also deteriorated. Therefore, it is recommended that the upper limit be 300 seconds.

The average cooling rate CR1 after the heating is particularly important for generating a predetermined amount of martensite and refining the old γ grain size, and in the present invention, the rapid cooling process is 30 ° C./s or more. Further, the average cooling rate CR1 also affects the last γ _R form, and if the average cooling rate is high, it exhibits a lath shape. Preferably it is 50 degrees C / s or more. The upper limit of the average cooling rate is not particularly limited, and a larger value is better for obtaining a desired fine martensite. However, it is recommended that the average cooling rate be appropriately controlled in relation to the actual operation level.

  It cools to temperature T2 below Ms point with said average cooling rate. This is because desired fine martensite can be obtained. In addition, although T2 is so preferable that it is small, considering the relationship with an actual operation level, it is generally recommended to set it as about room temperature.

The present invention is characterized in that the series of rapid heating and quenching processes described above are performed at least twice (multiple times). This is because a desired fine old γ can be obtained. Incidentally, in FIG. 3, as described above, the rapid heating and quenching process is performed twice, and the first rapid heating and quenching process is as described above, but the second rapid heating and quenching process is “A ₃ points or more and 1100 ° C. After heating and holding at the following temperature (T3 in FIG. 3) for 10 seconds or longer (t3 in FIG. 3), at an average cooling rate of 30 ° C./s or higher (CR2 in FIG. 3) In FIG. 3, the process includes the step of “cooling to T4)”.

  In addition, when performing rapid heating and quenching treatment a plurality of times, each rapid heating and quenching treatment may be performed in exactly the same process as long as the above-described conditions are satisfied, or different processes may be performed within the range of the conditions. Also good. In addition, the number of treatments of the “rapid heating / cooling treatment” is not particularly limited as long as it is at least 2 times, and as the number of treatments is increased, the effect of miniaturization increases. More preferably, it is recommended that it be 3 times or less, most preferably 2 times.

  The reason why the grain size of the former γ is refined by multiple “rapid quenching treatments” in this way is unknown in detail, but when the previous structure is martensite, the number of γ nucleation sites increases. For this reason, it is considered that the martensite is refined by repeating the above-mentioned “rapid heating / cooling treatment”.

Incidentally, the method itself for refining the grain size of the former γ by repeating rapid heating and quenching immediately above the A _c3 point has already been reported, for example, by Grange et al. (“Steel Heat Treatment Revision 5”, pages 80-81). (Issued on August 20, 1981, edited by the Japan Iron and Steel Institute). However, Grange et al. Only disclosed the refinement technology of the old γ grains in general, and it is concrete until the refinement technology of the former γ grains can be applied to the TRIP steel plate targeted by the present invention. Is not disclosed. Of course, even if the above-mentioned gazette is scrutinized, it is useful to improve delayed fracture resistance in providing an ultra-high strength steel sheet having both excellent elongation and delayed fracture resistance in an ultrahigh strength range of 1180 MPa class or higher. It is not even suggested about the technical idea of the present invention that it is effective to apply the above method to obtain fine martensite. Therefore, the present invention cannot be immediately derived based on the above-mentioned publication, which only discloses a general refinement technique of old γ grains.

(1-2) Second continuous annealing step (later continuous annealing step) or plating step The above steps are performed at a temperature of (A ₃ point-25 ° C.) to A ₃ point (T5 in FIG. 3) 120-600. After heating and holding for 2 seconds (t5 in FIG. 3), it is cooled to a temperature (T6 in FIG. 3) at an average cooling rate (CR3 in FIG. 3) of 3 ° C./s or more and Ms point to Bs point, A step of holding for 1 second or longer (t6 in FIG. 3) in the temperature range is included. These conditions are set to obtain the desired amount of ferrite, γ _R, and massive martensite by tempering the martensite mainly generated in the first continuous annealing step described above to obtain the desired tempered martensite. It has been done. In particular, in the present invention, the annealing process in the two-phase region is performed at a substantially high temperature for a long time as compared with the method of the prior application invention (that is, T5 and t5 are substantially higher than those of the prior application invention). This is characterized in that a predetermined amount of massive martensite (hard martensite) that could not be obtained by the invention of the prior application is generated.

First, desired tempered martensite and ferrite are generated by soaking at a temperature T5 of (A ₃ point-25 ° C.) to A ₃ point for 120 to 600 seconds (two-phase region annealing). This is because if the temperature is exceeded, all become γ, while if it is below the temperature, desired bulk martensite cannot be obtained. Control of the heating and holding time t5 in addition to the heating temperature T5 is extremely important for obtaining a desired structure. This is because if it is less than 120 seconds, the amount of γ produced is small and the desired massive martensite and γ _R cannot be obtained. Preferably it is 130 seconds or more, More preferably, it is 140 seconds or more. In addition, when 600 seconds are exceeded, the lath-like structure which is the characteristic of tempered martensite cannot be maintained, and the mechanical characteristics deteriorate. Preferably it is 500 seconds or less, More preferably, it is 400 seconds or less.

Next, the average cooling rate CR3 is controlled to 3 ° C./s or more (preferably 5 ° C./s or more), and avoiding pearlite transformation, the Ms point or more (preferably 350 ° C. or more) and the Bs point or less (preferably 450 ° C.). The temperature is cooled to a temperature T6 of the following) and further maintained in this temperature range for 1 second or longer (preferably 5 seconds or longer) (austempering). Thereby, the C concentration to γ _R can be obtained in a large amount and in a very short time. Especially to the C gamma _R than 0.85% is the retention time more than 200 seconds (preferably 220 seconds or more, more preferably more than 240 seconds) is recommended that a.

  Here, when the average cooling rate CR3 falls below the above range, desired tempered martensite cannot be obtained, and pearlite or the like is generated. The upper limit is not particularly defined, and the larger the better, the better. However, it is recommended to control appropriately in relation to the actual operation level.

As described above, the austempering treatment is performed after cooling, and the austempering temperature T6 is particularly important for securing the desired structure and exerting the action of the present invention. If the temperature is controlled within the above temperature range, a stable and large amount of γ _R can be obtained, whereby the TRIP effect by γ _R is exhibited. If the temperature is lower than the Ms point, the desired γ _R cannot be obtained, and if it exceeds the Bs point, a large amount of pearlite is generated.

  The upper limit of the holding time t6 at the temperature T6 is not particularly limited, but it is recommended to control it to 3000 seconds or less, preferably 2000 seconds or less in consideration of the time for austenite to transform into bainite.

Moreover, in the said process, the bainite structure may be produced | generated in the range which does not impair the effect | action of this invention other than a desired mixed structure | tissue and (gamma) _R. Further, plating or further alloying treatment may be performed without significantly degrading the desired structure and within the range not impairing the action of the present invention.

(2) [Hot rolling process] → [Continuous annealing process or plating process]
This method is a method of manufacturing a desired steel sheet through (2-1) hot rolling step, cold rolling if necessary, and (2-2) continuous annealing step or plating step. Of these, the (2-1) hot rolling process includes the first rapid quenching process in the (1-1) first continuous annealing process of (1) described above, and (2-2) The continuous annealing step or the plating step includes the (1-1) first rapid annealing process (first continuous annealing step) in (1) described above, the rapid heating and quenching treatment after the second time, and (1) ( 1-2) Includes both the second continuous annealing step (the last continuous annealing step). That is, in the method (2), the step of “repeating a predetermined rapid heating and quenching process at least twice”, which is a characteristic part of the method of the present invention, is combined with the hot rolling step (2-1) and the (2- In each step of the continuous annealing step or the plating step of 2), it is performed at least once, and the specific treatment method is as described in detail in the above (1).

  Following the hot rolling in (2-1) above, the continuous annealing or plating in (2-2) is performed. If the shape after hot rolling is poor, the heat in (2-1) above is used for the purpose of shape correction. After performing the rolling, it may be cold-rolled before performing the continuous annealing or plating of (2-2). Here, the cold rolling rate is recommended to be 1 to 30%. This is because cold rolling exceeding 30% increases the rolling load and makes cold rolling difficult.

The method for producing the steel sheet of the present invention has been described in detail above. Steel, especially C gamma _R is 0.85% of the steel sheets obtained in this manner is especially suitable as press forming a steel sheet for ultra Thus, elongation and resistance to hydrogen embrittlement are both significantly elevated A high-strength press-formed part is obtained.

In particular, in producing an ultra-high strength press-molded part with extremely excellent hydrogen embrittlement resistance, as described above, the austempering time is increased to 200 seconds or more and the Cγ _R is 0.85% at the steel plate production stage. As described above, it is useful to stabilize γ _R , and in the processing stage, the steel plate with γ _R stabilized in this way is press-formed particularly at a warm temperature (about 50 to 400 ° C.) Processing) is recommended. Therefore, according to the present invention, it is very useful in that it can provide an ultra-high strength steel sheet suitable for warm working performed when press molding automotive parts and the like. Here, the warm working in the present invention means forming at 50 to 400 ° C. (preferably 100 ° C. or more and 300 ° C. or less; most preferably around 200 ° C.), and the entire steel sheet falls within the temperature range. The temperature may be soaked appropriately.

  Hereinafter, the present invention will be described in detail based on examples. However, the following examples are not intended to limit the present invention, and all modifications made without departing from the spirit of the preceding and following descriptions are included in the technical scope of the present invention.

Example 1: Properties as a steel sheet (examination of component composition)
In this example, the test steels No. A to O having the component compositions shown in Table 1 (the balance is iron and impurities, and the unit in the table is mass%) are vacuum-melted to obtain an experimental slab. In accordance with the method of the present invention comprising the following steps (hot rolling → cold rolling → first continuous annealing → second continuous annealing), a hot rolled steel sheet having a thickness of 3.2 mm is obtained, and then the surface scale is removed by pickling. And cold rolled to a thickness of 1.2 mm.

Hot rolling process : Start temperature (SRT) 1150 ℃, Finishing temperature (FDT) 850 ℃, Cooling rate 40 ℃ / s, Winding temperature 550 ℃
Cold rolling process : 50% cold rolling rate
First continuous annealing step : The process of “holding at 930 ° C. for 120 seconds and then cooling to room temperature at an average cooling rate of 50 ° C./s (water cooling)” was repeated twice.

Second continuous annealing step : After holding at 850 ° C. for 180 seconds, after cooling to 420 ° C. at an average cooling rate of 25 ° C./s (water cooling), holding at 420 ° C. for 120 seconds.

  Each steel plate thus obtained was measured for tensile strength (TS), elongation [total elongation (EI)], and hydrogen embrittlement resistance (cathode CH life) in the following manner.

[Measurement of tensile strength (TS) and elongation]
The tensile test used a JIS No. 5 test piece and measured tensile strength (TS) and elongation (EI). The strain rate in the tensile test was 1 mm / sec.

[Measurement of hydrogen embrittlement resistance]
In measuring the hydrogen embrittlement resistance, a strip test piece in which each steel plate was adjusted to a size of 15 mm × 65 mm was used. A stress of 980 MPa was applied to the strip test piece by four-point bending, and in a mixed solution of (0.5 mol sulfuric acid + 0.01 mol KSCN), a potential of −80 mV, which was lower than natural potential using a potentiostat. The hydrogen embrittlement resistance [cathode charge (CH) life] was evaluated by measuring the time during which cracking occurred when an electric potential was applied. In the present invention, a steel sheet having a measurement time of 1000 seconds or more according to the above method is evaluated as a “steel plate excellent in hydrogen embrittlement resistance”.

[Measurement of space factor of organization]
The space factor (area ratio) of the structure (excluding γ _R ) in the steel sheet identifies the structure by repeller corrosion of the steel sheet and observing the sheet thickness center position of the cross section in the rolling direction with an optical microscope (1000 times magnification). And then measured. Incidentally, gamma space factor of _R (volume fraction) and gamma C concentration of _R (C gamma _R), after grinding to 1/4 the thickness of the steel sheet was measured by X-ray diffraction method from the chemical polishing (ISIJ Int. Vol. 33. (1933), No. 7, P. 776).

  In addition, the “space factor of fine massive martensite” which is an object of the present invention was measured in the following manner.

  First, as described above, two steel plate structure photographs are prepared in which the steel plate is subjected to repeller corrosion and the center position of the thickness in the rolling direction section is observed with an optical microscope (× 1000). An area of 50 μm × 50 μm is arbitrarily selected from each photograph and cut out. After obtaining the area ratio of massive martensite (aspect ratio is 3 or less) occupying the total area (50 μm × 50 μm × 2) and the average particle size of the martensite for the two cut out photographs, The area ratio of fine martensite having an average particle diameter of 5 μm or less is calculated and is defined as “the space factor of fine massive martensite”.

  These results are shown in Table 2.

  From these results, it can be considered as follows (the following No. means the experiment No. in Table 2).

  First, No. 2 to 4 and 8 to 15 are steel sheets having a predetermined structure according to the present invention method using steel types (No. B to D and H to N in Table 1) that satisfy the scope of the present invention. Although it is an example which manufactured this, it turns out that it is excellent in both the elongation and the hydrogen embrittlement resistance in the ultra high strength region of 1180 MPa or more.

  On the other hand, the following examples using steel types (No. A and E to G in Table 1) that do not satisfy the scope of the present invention have the following problems.

  First, No. 1 is an example using a steel type A with a small amount of C, and a predetermined amount of massive martensite (hard) cannot be obtained, and the ferrite is excessive, so the strength is reduced. The hydrogen embrittlement resistance was not measured because the strength was so low that a four-point bending stress could not be applied.

No. 5 is an example using steel type E with a small total amount of (Si + Al), and the desired γ _R cannot be obtained, so the elongation is low.

No. 6 is an example using steel type F with a small amount of Mn, and the desired γ _R cannot be obtained, so the elongation is low.

  No. 7 is an example using steel type G with a large amount of Mn, and because the strength was too high, hot rolling cracks occurred during hot rolling, and the subsequent annealing treatment could not be performed.

Example 2: Characteristics as a steel sheet (examination of manufacturing conditions)
In this example, an experimental slab using the steel type B shown in Table 1 (steel satisfying the composition of the present invention) was used, and after hot rolling and cold rolling under the same conditions as in Example 1, various types shown in Table 3 were used. The cold-rolled steel sheets shown in No. 1 to 22 of Table 4 were obtained by performing the first continuous annealing (implementing the predetermined rapid heating and quenching treatment twice) and the second continuous annealing under the conditions described above. The plate thickness is all 1.2 mm.

  Next, the structure and various characteristics of the steel sheet were examined in the same manner as in Example 1. These results are shown in Table 4.

  First, No. 1 is an example of the present invention in which a cold rolled steel sheet having a desired structure is manufactured by performing hot rolling → cold rolling → first continuous annealing → second continuous annealing under the conditions specified in the present invention; and No. 20 is an example of the present invention that was subjected to the above-described steps and then further alloyed (heated at 500 ° C. for the purpose of alloying after being immersed in hot dip galvanizing). It has ultra-high strength and is excellent in both elongation and hydrogen embrittlement resistance.

  On the other hand, Nos. 2 to 22 that do not satisfy any of the conditions defined in the present invention have the following problems.

  Of these, Nos. 2 to 6 are examples in which any of the conditions of the first rapid quenching process in the first continuous smelting process does not satisfy the scope of the method of the present invention.

First, since No. 2 has a heating temperature T1 of 850 ° C. and less than A ₃ point, the old γ grains cannot be made fine, and a large amount of massive martensite is generated, resulting in hydrogen embrittlement resistance. The characteristics deteriorated.

  In No. 3, since the heating time t1 was too short, austenitization did not proceed sufficiently, and cementite remained, so the old γ grains could not be refined. As a result, massive martensite was produced in a large amount, and the hydrogen embrittlement resistance deteriorated.

  No. 4 is an example in which the heating temperature T1 was too high, leading to the growth of γ grains, and the old γ grains could not be made finer. As a result, a large amount of massive martensite was formed, resulting in hydrogen embrittlement resistance. The aging characteristics deteriorated.

  No. 5 is an example in which the cooling rate CR1 is small; No. 6 is an example in which the heating temperature T2 after cooling is 400 ° C. and exceeds the Ms point, both of which have a martensite structure after the first rapid heating and quenching. The old γ grains could not be made fine. As a result, a large amount of massive martensite was generated, and the hydrogen embrittlement resistance was deteriorated.

  Next, Nos. 7 to 12 are examples in which any of the conditions in the second rapid quenching process in the first continuous smelting process does not satisfy the scope of the method of the present invention.

First, No. 7 is an example in which the heating temperature T3 is 930 ° C., which is lower than the A ₃ point, and the old γ grains cannot be made fine, so that massive martensite is generated in large amounts, The hydrogen embrittlement characteristics decreased.

  No. 8 is an example in which the heating time t3 is too short, and austenitization did not proceed sufficiently, and martensite was not sufficiently obtained after cooling. As a result, a predetermined amount of tempered martensite was not obtained, and the elongation decreased. Moreover, the amount of massive fine martensite became small, and the hydrogen embrittlement resistance decreased.

  No. 9 is an example in which the heating temperature T3 is too high, leading to the growth of γ grains and making the old γ grains fine, resulting in the formation of large amounts of massive martensite and resistance to hydrogen embrittlement. Decreased.

  No. 10 is an example in which the heating time t3 is too long, and since the old γ grains could not be made fine by causing the grain growth of γ, a large amount of massive martensite was generated, and the hydrogen embrittlement resistance was low. Declined.

  No. 11 is an example in which the cooling rate CR2 is small; No. 12 is an example in which the temperature T4 after cooling exceeds 400 ° C. and exceeds the Ms point, both of which have a martensite structure after the second rapid heating and quenching. I couldn't. As a result, a predetermined amount of tempered martensite was not obtained, and the elongation decreased. Moreover, the amount of massive fine martensite became small, and the hydrogen embrittlement resistance decreased.

  Next, Nos. 13 to 19 are examples in which any of the conditions in the second continuous smelting step does not satisfy the scope of the method of the present invention.

  First, No. 13 is an example in which the heating temperature T5 is 930 ° C., which exceeds the upper limit of the heating temperature specified in the present invention, and a predetermined lath-like structure cannot be obtained by austenitization, and the desired tempered martensite, etc. Was not obtained at all, so the elongation was low.

  No. 14 is an example in which the heating temperature T5 is too low at 800 ° C., and a predetermined amount of massive martensite cannot be obtained, resulting in a low TS.

  No. 15 is an example in which the heating time t5 is too long, and the desired tempered martensite cannot be obtained, so the elongation becomes low.

No. 16 is an example in which the heating time t5 is too short. Since the amount of γ produced is small, a predetermined amount of γ _R cannot be obtained, the elongation becomes low, and the hydrogen embrittlement resistance is also deteriorated.

  No. 17 is an example in which the cooling rate CR3 is small, and a large amount of pearlite is generated, and a predetermined amount of massive martensite cannot be obtained, and a desired strength cannot be ensured.

No. 18 is an example in which the heating temperature T6 after cooling is too high. Since pearlite is produced in a large amount, finally obtained γ _R is reduced, elongation is lowered, and hydrogen embrittlement resistance is lowered. did.

No. 19 is an example in which the heating temperature T6 after cooling is too low, and the final obtained γ _R is as very low as 1%, so the EL becomes low. Moreover, since γ _R is reduced and a large amount of massive martensite is generated, the hydrogen embrittlement resistance is deteriorated.

  Further, No. 21 is an example in which the first rapid heating and quenching process in the first continuous tempering process is omitted, and corresponds to the prior invention. Since the predetermined rapid heating / cooling treatment is not performed twice as in the method of the present invention, the refinement of the old γ grains is not performed, and a predetermined amount of fine martensite is not generated. Although the elongation could be secured, the hydrogen embrittlement resistance decreased.

  No. 22 is an example in which the first continuous tempering step is omitted, and corresponds to a conventional polygonal ferrite-containing TRIP steel sheet. Since the predetermined first continuous annealing process, which is a feature of the present invention method, is not performed, the refinement of the old γ grains is not performed, and no fine martensite is generated at all. Further, no tempered martensite contributing to the improvement of ductility was obtained, and the ferrite was excessive, so that the elongation was lowered.

  For reference, FIG. 4 shows optical micrographs (1000 × magnification) of the inventive example (No. 1) and the comparative example (No. 21), respectively. From these figures, No. 1 produced by the method of the present invention produced the desired fine massive martensite, whereas No. 21 produced by the method of the prior invention had coarse massive martensite. It can be seen that

Example 3 (Processing strain simulation experiment)
In this example, the resistance to hydrogen embrittlement when a steel sheet obtained by variously changing the austempering time was applied with various working strains to form a press-formed part (working strain simulation experiment) was examined.

  Specifically, steel type B shown in Table 1 (steel satisfying the composition of the present invention) was used, and in Example 1, the second continuous smelting step was held at 420 ° C. for 120 seconds or 360 seconds. A steel plate was obtained in the same manner as in Example 1.

  The following experiments were conducted for the purpose of investigating the effects of processing on each steel sheet thus obtained.

First, a tensile test piece (JIS No. 5) is collected from the steel sheet, and after applying various pre-strains (up to 10%) at room temperature or warm (100 ° C.), the width from the parallel part of the tensile test piece is increased. A test piece having a length of 15 mm and a length of 65 mm was collected. This specimen, JIS Z Vickers hardness test based on 2244; as well as implement (Hv test load 9.807N), γ space factor of _R and martensite, gamma C concentration (C gamma _R) of _R, and resistance to hydrogen The embrittlement characteristics (cathode CH life) were measured in the same manner as in Example 1.

  These results are shown in Table 5 (results when processed at room temperature) and Table 6 (results when processed warm). In this example, unlike Example 1 described above, the cathode CH lifetime measurement time of 900 seconds or more by the above method was evaluated as “a press-formed part excellent in hydrogen embrittlement resistance”. In this example, unlike Example 1 and the like, processing strain is imparted, and from the viewpoint that deterioration of hydrogen embrittlement resistance due to processing is unavoidable to some extent, the evaluation standard is 900 seconds or more. It is up to setting a little lower than 1.

Further, these tables, although the space factor does not describe other tissues other than martensite and gamma _R (tempered martensite and ferrite), any space factor of the other tissues, the It is confirmed that the scope of the invention is satisfied.

  From these tables, it can be considered as follows.

  Of these, Table 5 summarizes the results of the structure and various characteristics when the austempering time is 120 seconds or 360 seconds and the strain amount is changed in the range of 0 to 10% at room temperature.

First, when the austempering time was 120 seconds, a decrease in hydrogen embrittlement resistance was observed with an increase in processing strain (high processing). That is, when the strain amount is up to 6%, all of the massive martensite, γ _R , and Cγ _R satisfy the scope of the present invention, so that a desired cathode CH life (900 seconds or more) is obtained. On the other hand, when the amount of strain is higher than 8%, a large amount of massive martensite is generated and the retained austenite is reduced, so that the cathode CH life is remarkably lowered and the hydrogen embrittlement resistance is inferior. Incidentally, Cγ _{R in} the steel sheet before processing (that is, strain amount 0%) is 0.80%.

On the other hand, when the austempering time is set to 360 seconds, even if the strain amount is increased to 10%, all of the massive martensite, γ _R , and Cγ _R satisfy the scope of the present invention. Excellent hydrogen embrittlement resistance. Incidentally, Cγ _{R in} the steel plate before processing (ie, strain amount 0%) is 0.88%.

From these results, when producing the ultra-high strength press-molded part by subjecting the steel sheet of the present invention to room temperature, a method of stabilizing the γ _R by increasing the austempering time so that the Cγ _R is 0.85% or more. This makes it possible to obtain a press-molded part having excellent hydrogen embrittlement resistance even after high processing.

  Next, Table 6 summarizes the results of the structure and various characteristics when the austempering time is 120 seconds or 360 seconds and the strain amount is changed in the range of 0 to 10% in the warm (100 ° C.). Is.

When the results of each austempering time (120 seconds / 360 seconds) are compared with the results of Table 5 described above, the γ _R transformation rate (γ _R) associated with machining is higher in warm machining than in room temperature machining. It is recognized that there is a strong tendency to suppress the degree of transformation to the martensite, and as a result, the hydrogen embrittlement resistance is also improved.

That is, when the amount of decrease in γ _{R and} the increase rate of massive martensite when the strain amount was increased to 2%, 4%, and 6% were compared, respectively, when the austempering time was 120 seconds, As the amount increases to 2%, 4%, and 6%, γ _R decreases to 13%, 8%, and 6%, and conversely, massive martensite increases to 36%, 42%, and 44%. On the other hand, when the austempering time is 360 seconds, the decrease rate of γ _{R and} the increase rate of massive martensite with increasing strain are slightly suppressed as compared with the case where the austempering time is 120 seconds. As the amount increases to 2%, 4%, and 6%, γ _R decreases to 14%, 11%, and 8%, and conversely, massive martensite increases to 35%, 39%, and 42%. These rates of change are slightly higher than when the austempering time is 120 seconds. It tends to slow down.

From the above results, in producing an ultra-high strength press-molded part having excellent hydrogen embrittlement resistance, in the steel plate production stage, the austempering time is lengthened so that Cγ _R in the steel plate is 0.85% or more and γ It is useful to stabilize _R (not so much compared with the cathode CH life of 0% strain in Table 5); however, in the processing stage, γ _R transformation compared to room temperature processing It is useful to further stabilize γ _R by applying warm working, which is said to be easy to suppress: The combined use of these materials may further improve the resistance to hydrogen embrittlement. I understand.

It is the graph which contrasted the hardness of the tempered martensite in the same component system, and the hardness of polygonal ferrite. It is a graph which shows the influence of the amount of C exerted on the hardness of tempered martensite and polygonal ferrite. It is the schematic explaining the typical process of this invention method (what performed the rapid heating rapid cooling process twice in the 1st continuous annealing process). It is an optical microscope photograph (1000-times multiplication factor) of the example (No. 1) of this invention in Example 2, and a comparative example (No. 21).

Claims (8)

% By mass
C: 0.06 to 0.6%,
Si + Al: 0.5 to 3%
Mn: 0.5-3%,
P: 0.15% or less (excluding 0%),
S: 0.02% or less (including 0%)
Containing
The rest: iron and inevitable impurities
and,
The organization is the space factor for the whole organization,
15-60% tempered martensite,
5-50% of ferrite,
It contains 15 to 45% of massive martensite with a retained austenite of 5% or more and an aspect ratio of 3 or less, and the space factor occupied by fine martensite with an average particle size of 5 μm or less in the massive martensite is 30 % Ultra-high strength steel sheet excellent in elongation in an ultra-high strength region having a tensile strength of 1180 MPa or more and resistance to hydrogen embrittlement. Furthermore, in mass%,
Mo: 1% or less (excluding 0%),
Ni: 0.5% or less (excluding 0%),
Cu: 0.5% or less (excluding 0%),
Cr: 1% or less (excluding 0%)
The ultra-high-strength steel sheet according to claim 1, which contains at least one of the following. Furthermore, in mass%,
Ti: 0.1% or less (excluding 0%),
Nb: 0.1% or less (excluding 0%),
V: 0.1% or less (excluding 0%)
The ultra high strength steel sheet according to claim 1 or 2, which contains at least one of the following. Furthermore, in mass%,
Ca: 0.003% or less (not including 0%) and / or REM: 0.003% or less (not including 0%)
The ultra-high strength steel sheet according to any one of claims 1 to 3.   The ultra high strength steel sheet according to any one of claims 1 to 4, wherein a solid solution carbon amount in the retained austenite is 0.85% or more. A method for producing the ultra high strength steel sheet according to any one of claims 1 to 4,
The steel satisfying the component according to any one of claims 1 to 4 is heated and held at a temperature of A ₃ point or more and 1100 ° C. or less for 10 seconds or more, and then at an average cooling rate of 30 ° C./s or more and Ms point or less. Including at least twice a step of cooling to temperature, and after heating and holding at a temperature of (A ₃ point −25 ° C.) to A ₃ point for 120 to 600 seconds, at an average cooling rate of 3 ° C./s or more, Ms The manufacturing method of the ultra high strength steel plate characterized by including the process cooled to the temperature below a point and below a Bs point, and hold | maintaining at this temperature range for 1 second or more. A method for producing the ultra high strength steel sheet according to claim 5,
The steel satisfying the component according to any one of claims 1 to 4 is heated and held at a temperature of A ₃ point or more and 1100 ° C. or less for 10 seconds or more, and then at an average cooling rate of 30 ° C./s or more and Ms point or less. Including at least twice a step of cooling to temperature, and after heating and holding at a temperature of (A ₃ point −25 ° C.) to A ₃ point for 120 to 600 seconds, at an average cooling rate of 3 ° C./s or more, Ms A method for producing an ultra-high strength steel sheet, comprising a step of cooling to a temperature not lower than the point and not higher than the Bs point and holding the temperature in the temperature range for 200 seconds or longer.   A method for producing an ultra-high-strength press-molded part, wherein the ultra-high-strength steel sheet according to claim 5 is warm press-molded.

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