JP5110970B2 - High strength steel plate with excellent stretch flangeability - Google Patents

Description

  The present invention relates to a high-strength steel sheet excellent in stretch flangeability. Specifically, the tensile strength (TS) is 980 MPa or more, and the product of TS (MPa) and stretch flangeability [λ (%)] [TS] × λ (MPa ·%)] is 50000 or more, and relates to a high-strength steel sheet having an excellent balance between tensile strength and stretch flangeability. The high-strength steel sheet of the present invention is widely used in industrial fields such as automobiles, electric machines, and machines, and is particularly suitable for automobile parts in which the application of high-strength steel is expanding. Specifically, for example, collision parts such as front and rear side members and crash boxes, pillars such as center pillar reinforcements, body components such as roof rail reinforcements, side sills, floor members, and kick parts. used.

  Steel plates used in the fields of automobiles, electrical machinery, machinery, etc. are required to have both excellent strength and ductility. As a steel plate that meets such needs, TRIP (Transformation Induced Plasticity) ) Steel sheets are attracting attention. The TRIP steel sheet contains retained austenite (residual γ) in which the austenite (γ) structure remains. When the TRIP steel sheet is deformed by processing, the retained austenite is induced and transformed into martensite by stress, and the excellent elongation and martensite due to γ are obtained. High strength can be obtained. Depending on the type of matrix, the TRIP steel sheet is a TRIP type composite structure steel (hereinafter referred to as TDP steel) having polygonal ferrite as a matrix, and a TRIP type tempered martensite steel having a tempered martensite as a matrix ( Hereinafter, it is referred to as TAM steel), and TRIP type bainitic steel (hereinafter referred to as TBF steel) having bainitic ferrite as a matrix.

  Among these, TBF steel has the characteristics that high strength is easily obtained by a hard bainite structure and has high elongation characteristics, but there is a problem that it is inferior in stretch flangeability (hole expandability, local ductility). .

In order to solve the above problem, for example, Patent Documents 1 and 2 disclose a technique for improving the hole expandability (synonymous with stretch flangeability) and hydrogen embrittlement resistance of TBF steel by suppressing the ratio of residual γ as much as possible. Proposed. This is made in view of the conventional recognition that “when martensite or retained austenite is used in the second phase, the hole expandability is significantly reduced”.
JP 2004-323951 A (claims, paragraph [0010], etc.) JP 2004-332100 A (Claims, paragraph [0007], etc.)

  The present invention has been made in view of the above circumstances, and the purpose thereof is a TRIP steel sheet having bainitic ferrite as a matrix, and has a tensile strength (TS) of 980 MPa or more and excellent stretch flangeability. It is to provide a high strength steel sheet.

  The high-strength steel sheet excellent in stretch flangeability according to the present invention that can solve the above-mentioned problems has components in the steel of C: 0.10 to 0.20% (meaning mass%, hereinafter the same), Si: 0.8 to 2.5%, Mn: 1.5 to 2.5%, Al: 0.01 to 0.10%, P: less than 0.1% (excluding 0%), S: 0.0. Less than 002% (excluding 0%), the balance: iron and inevitable impurities are satisfied, and the structure contains at least bainitic ferrite and retained austenite, and the area ratio to the entire structure is Tick ferrite: 70% or more, retained austenite: 2 to 20%, polygonal ferrite and / or quasi-polygonal ferrite: 15% or less (excluding 0%), and average in the retained austenite Residual austicule with particle size of 5μm or less Proportion of Knight exist gist that 60% or more.

  In a preferred embodiment, Ca is further contained in an amount of 0.002% or less (excluding 0%).

  Since the present invention is configured as described above, a high-strength steel sheet having a high strength of 980 MPa or more and an excellent stretch flangeability can be obtained. If the steel plate of the present invention is used, it is possible to satisfactorily form automobile parts and other industrial machine parts that require high strength.

  The present inventor has intensively studied a TRIP steel sheet having a bainitic ferrite (BF) structure as a parent phase mainly for improving stretch flangeability. As a result, if the retained austenite (residual γ) is made into a fine shape with an average particle size of 5 μm or less, the stretch flangeability (λ) is dramatically increased. For that purpose, in particular, tempering in the continuous annealing process. It has been determined that the hot-dip Zn plating process may be performed after the temperature is controlled to be lower than before, and the present invention has been conceived.

  With reference to FIG. 1, the effect of increasing λ by miniaturizing the residual γ will be described. In FIG. 1, the horizontal axis represents the ratio (%) of residual γ having an average particle diameter of 5 μm or less (hereinafter sometimes referred to as “fine residual γ”) to the total residual γ, and the vertical axis represents It is a hole expansion rate (λ).

  As shown in FIG. 1, in the TBF steel (in the figure, ●) having BF as a matrix, λ is remarkably increased when the fine residual γ ratio is 60% or more. For reference, Fig. 1 also shows the results of TRIP steel (TDP steel) with polygonal ferrite (PF) as the parent phase (■ in the figure), but contrasts the results of TBF steel and TDP steel. Then, it can be seen that the effect of increasing λ due to fine residual γ is recognized significantly more in TBF steel than in TDP steel.

  In addition, the result of TBF steel shown in FIG. 16 to 21 are plotted, and the results for TDP steel are shown in No. 5 of Table 5. The data of 1-4 is plotted.

  The high-strength steel sheet of the present invention has a tensile strength (TS) of 980 MPa or more, a product of TS and stretch flangeability (λ) of 50000 or more, and preferably a tensile strength (TS) of 1000 MPa or more, TS and stretch flange. The product with the property (λ) satisfies 60000 or more. Further, the elongation (total elongation) characteristic satisfies 18% or more, preferably 20% or more.

  In addition to the hot-dip Zn-plated steel sheet, the present invention also includes an alloyed hot-dip Zn-plated steel sheet that has been further alloyed.

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

(Organization)
The high-strength steel sheet of the present invention includes at least bainitic ferrite and residual γ, and the area ratio to the entire structure is bainitic ferrite: 70% or more, residual γ: 2 to 20%, polygonal ferrite And / or quasi-polygonal ferrite: 15% or less (excluding 0%), and the proportion of residual γ having an average particle size of 5 μm or less in the residual γ has a structure satisfying 60% or more. ing.

Bainitic ferrite: 70% or more The steel sheet of the present invention contains bainitic ferrite (BF) as a main component (matrix).

Here, BF will be described in detail. BF means a substructure having a high dislocation density (initial dislocation density) (a lath-like structure may or may not be included), and is a bainite in that there is very little carbide in the structure. It is clearly different from the organization. BF is also different from polygonal ferrite (PF) having a substructure with a low dislocation density and quasi-polygonal ferrite (quasi-PF) having a substructure such as fine subgrains. These structures are shown, for example, in p4 Table 1 of Bainite Photobook-1 (published by the Japan Iron and Steel Institute Basic Research Group). In Table 1, BF is α _B and α _B °, PF is α _p , quasi PF is α _q . Furthermore, BF and PF / quasi-PF are clearly distinguished as follows by observation with a transmission electron microscope (TEM: magnification of about 10,000 times).
PF: White in the TEM photograph, polygonal shape, and hardly contains residual γ and martensite inside.
-Quasi-PF: White in a TEM photograph, almost spherical, and hardly contains residual γ and martensite inside.
-In a BF: TEM photograph, gray is shown by dislocations present inside. In many cases, BF and martensite cannot be separated and distinguished by TEM observation.

  Since BF has a higher dislocation density than PF and quasi-PF, it has characteristics that it can easily achieve high strength and has high stretch characteristics and stretch flangeability. In order to effectively exhibit such an action, the area ratio of BF is set to 70% or more. BF is preferably 80% or more, and more preferably 90% or more. In order to obtain a high-strength steel sheet having even more excellent stretch flangeability, it is recommended to control so as to have a substantially two-phase structure of BF and residual γ described later.

Residual γ: 2 to 20%
Residual γ is a particularly useful structure for improving elongation. As described in Patent Documents 1 and 2 described above, if residual γ is present in TBF steel, the stretch flangeability usually deteriorates, but as in the present invention, the fine residual γ occupies the residual γ. Stretch flangeability can be improved by increasing the ratio of.

  In order to effectively exhibit such an action, the area ratio of the total residual γ is set to 2% or more. However, if the residual γ is too large, the stretch flangeability tends to be reduced, so the upper limit is made 20%. The area ratio of the residual γ is preferably 5% or more and 18% or less, and more preferably 7% or more and 16% or less.

  Further, the C concentration in the residual γ is preferably 0.8% or more. The C concentration in the residual γ greatly affects the characteristics of TRIP (strain-induced transformation processing), and when controlled to 0.8% or more, the elongation characteristics are particularly enhanced. The elongation property is improved as the C concentration in the residual γ is increased, more preferably 1% or more, and further preferably 1.2% or more. The upper limit of the C concentration in the residual γ is not particularly limited, but the upper limit that can be adjusted in actual operation is considered to be approximately 1.6%.

Ratio of residual γ with an average particle size of 5 μm or less in the residual γ ≧ 60%
In order to obtain the desired stretch flangeability and elongation characteristics, it is necessary that the residual γ having an average particle size of 5 μm or less be present as much as possible. In the present invention, the fine residual γ occupies the total residual γ. The ratio is 60% or more. The larger the ratio of the fine residual γ, the better. It is preferably 70% or more, more preferably 80% or more, and most preferably 100%.

  The average particle size is measured as follows. First, after corroding a steel plate with a repeller and identifying residual γ by observation with an optical microscope (magnification: 1000 times), an average value of the particle size (maximum diameter) of residual γ existing in a visual field of 60 mm × 80 mm is calculated. Similarly, the average particle diameter in a total of 5 fields of view was calculated, and the average value thereof was defined as “average particle diameter of residual γ”.

  The smaller the average particle size of the fine residual γ, the better, and it is preferably 4 μm or less, more preferably 3 μm or less. Note that the lower limit is not particularly limited, and the lower limit can be a limit that allows the residual γ to be identified by the observation method described above and from which the average particle diameter can be calculated.

Polygonal ferrite and / or quasi-polygonal ferrite: 15% or less (excluding 0%)
The high-strength steel sheet of the present invention may be composed of a structure composed of the above-described bainitic ferrite (BF) and residual γ, and thereby the stretch flangeability is maximized. Polygonal ferrite (PF) and / or quasi-polygonal ferrite (quasi-PF) may be contained within a total range of 15% or less within a range not impairing the action. These area ratios are better as they are smaller, preferably 10% or less, and most preferably 0%.

In addition, the steel sheet of the present invention may be composed of only the above-described structure (that is, a mixed structure of BF and residual γ, or a mixed structure of BF, residual γ and PF / quasi-PF). Other structures (pearlite, bainite, martensite, cementite, etc.) that can remain in the manufacturing process may be included within a range that does not impair the action of the present invention. The smaller the area ratio, the better. For example, the total area ratio is preferably 10% or less at most.

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

C: 0.10 to 0.20%
C is an element necessary for ensuring high strength and obtaining a predetermined residual γ. More specifically, it is an important element for containing a sufficient amount of C in the austenite (γ) phase and leaving the desired γ phase at room temperature. In order to effectively exhibit such an action, 0.10% or more of C is added. On the other hand, if the amount of C is excessively added, the weldability is lowered, so the upper limit is made 0.20% or less. The C content is preferably 0.12% or more and 0.18% or less.

Si: 0.8-2.5%
Si is an element that effectively suppresses the formation of carbides, and is also useful as a solid solution strengthening element. In order to effectively exhibit such an action, Si is added by 0.8% or more. However, if the amount of Si is excessive, the above action is saturated and problems such as hot brittleness occur, so the upper limit is set to 2.5%. The amount of Si is preferably 1.2% or more and 2.3% or less.

Mn: 1.5 to 2.5%
Mn is an element necessary for stabilizing γ and obtaining a desired residual γ. In order to exhibit such an action effectively, 1.5% or more of Mn is added. However, if Mn is added excessively, adverse effects such as slab cracking appear, so the upper limit of Mn is set to 2.5%. The amount of Mn is preferably 1.8% or more and 2.3% or less.

Al: 0.01-0.10%
Al, like Si, is an element that effectively suppresses the formation of carbides. In order to effectively exhibit such an effect, 0.01% or more of Al is added. However, if Al is added in excess of 0.10%, polygonal ferrite is likely to be formed, and the stretch flangeability cannot be sufficiently improved. The amount of Al is preferably 0.03% or more and 0.06% or less.

P: less than 0.1% (excluding 0%)
If P is added excessively, the workability deteriorates, so the upper limit is suppressed to less than 0.1%. The smaller the amount of P, the better.

S: Less than 0.002% (excluding 0%)
S is a harmful element which forms sulfide inclusions such as MnS and degrades workability as a starting point of cracking, so the upper limit is made less than 0.002%. The smaller the amount of S, the better.

  The steel sheet of the present invention contains the above components, with the balance being iron and inevitable impurities. Examples of unavoidable impurities include N (nitrogen), which is an element brought in depending on the situation of raw materials, materials, manufacturing facilities, and O (oxygen) of 0.01% or less.

  Further, for the purpose of improving workability, it is preferable to contain Ca in the following range.

Ca: 0.002% or less (excluding 0%)
Ca is an element that controls the form of sulfide in steel and is effective in improving workability. In order to effectively exhibit such an action, it is preferable to add 0.0003% or more of Ca. However, even if Ca is added excessively, the above action is saturated and is economically wasteful, so the upper limit is preferably made 0.002% or less. The Ca content is more preferably 0.0005% or more and 0.0015% or less.

(Production method)
Next, the preferable method for manufacturing the steel plate of this invention is demonstrated.

  The manufacturing method of the steel sheet of the present invention includes (1) a hot rolling step (further, if necessary, a cold rolling step), (2) a continuous annealing step, and (3) a hot-dip Zn plating step (further, if necessary, an alloy Process). In order to efficiently obtain the steel sheet of the present invention, it is particularly important to appropriately control the tempering temperature in the continuous annealing process. By this process and the subsequent hot dip Zn plating process, bainitic A structure having a matrix structure of ferrite and a fine residual γ can be secured.

(1) Hot-rolling step The hot- rolling step is not particularly limited, and a generally used method can be adopted. Specifically, for example, after performing hot rolling at a temperature of about 1000 to 1300 ° C. for about 20 to 60 minutes and performing finish rolling at a temperature of about 850 to 890 ° C., about 40 to 60 ° C. It is preferable to cool at an average cooling rate of / s and wind at a temperature of about 500 to 600 ° C. Thus, it is preferable that the thickness of the hot-rolled steel sheet obtained is generally 2.0 to 3.5 mm.

  The hot-rolled steel sheet obtained as described in (1) is pickled to remove scale.

  Further, for the purpose of improving workability, cold rolling may be performed at a cold rolling rate of about 50 to 80%. The thickness of the cold-rolled steel sheet thus obtained is preferably about 0.7 to 1.4 mm.

(2) Continuous annealing process A continuous annealing process is demonstrated, referring FIG. This step is mainly set for obtaining a desired quenched structure (quenched martensite). 2 (a) and FIG. 2 (b) differ only in the cooling conditions after the soaking step [T1 (° C.) × t1 (seconds)], and FIG. 2 (a) shows the roll quench (RQ). examples of subjecting a, FIG. 2 (b) is an example subjected to c _o Takuenchi (WQ).

First, hold at (t1) for 10 seconds or more at a temperature (T1) of _Ar3 point or higher (soaking step).

  By setting it as said soaking temperature (T1), a carbide | carbonized_material melt | dissolves completely and desired residual (gamma) is obtained efficiently. It is also effective in forming bainitic ferrite in the cooling step after soaking. Although the upper limit of soaking temperature (T1) is not specifically limited, For example, from a viewpoint of cost reduction etc., it is generally preferable to set it as 1000 degrees C or less.

  Further, if the holding time (t1) at the soaking temperature (T1) is less than 10 seconds, the soaking effect of the whole steel sheet is insufficient, and finally, the ratio of fine residual γ and bainitic ferrite Generation is reduced (see examples below). In order to effectively exhibit such an action, the longer the holding time (t1) is, the better. However, when t1 exceeds 300 seconds, the above action is saturated and only the cost is increased. The holding time (t1) is preferably 60 seconds or more and 200 seconds or less.

Next, cooling is performed from the soaking temperature (T1) to a predetermined temperature (T2) while avoiding ferrite transformation and pearlite transformation at an average cooling rate (CR) of 10 ° C./s or more.

  T2 (tempering temperature) is 100 ° C. or higher and 300 ° C. or lower. In order to generate a large amount of fine residual γ, it is extremely important to set T2 in the above range. When the temperature of T2 exceeds 300 ° C., the hardened structure decreases, so finally, As a result of the study by the present inventors, it became clear that a lot of lath-like and coarse residual γ was generated and the desired fine residual γ could not be obtained (see Examples described later). On the other hand, when the temperature of T2 is less than 100 ° C., there is a problem that it is not tempered. The temperature of T2 is preferably 100 ° C. or higher and 280 ° C. or lower.

  In addition, when the average cooling rate (CR) from the soaking temperature (T1) to the temperature of T2 is less than 10 ° C / s, the generation of bainitic ferrite in the parent phase is reduced, and polygonal ferrite or quasi-polygonal ferrite is reduced. (See examples below). The average cooling rate is preferably 20 ° C./s or more. The upper limit is not particularly limited, and the faster it is, the better. However, appropriate control is recommended in relation to the actual operation level. Examples of the cooling method include air cooling, mist cooling or water cooling of a roll used during cooling.

  Specifically, as shown in FIG. 2A, the range from the soaking temperature (T1) to the temperature of T2 may be cooled at a predetermined average cooling rate (CR), or FIG. As shown in b), at a predetermined average cooling rate (CR), the temperature may be cooled from the soaking temperature (T1) to room temperature at once, and then heated to the temperature of T2. The heating rate from room temperature to the temperature of T2 is not particularly limited, but in consideration of the heating performance of the equipment, it is preferably in the range of about 1 to 10 ° C./s. As shown in Examples described later, a desired tissue can be obtained by applying any method. However, in consideration of productivity and economy, it is preferable to adopt the method of FIG.

  Next, after maintaining (t2) for 100 to 600 seconds at the above temperature (T2) (tempering treatment), it is cooled to room temperature. As a result, C concentration to residual γ can be efficiently performed in a short time, and finally, a large amount of fine residual γ can be generated, and a predetermined amount of bainitic ferrite can be secured. You can also. The holding time (t2) is preferably 120 seconds or longer and 300 seconds or shorter. The method for cooling from the temperature (T2) to room temperature is not particularly limited, and for example, water cooling, gas cooling, air cooling, or the like can be employed.

  In the series of steps described above, for example, a CAL (actual machine) may be used, or a CAL simulator may be used.

(3) Hot Zn plating step The hot Zn plating step will be described with reference to FIG. This process is mainly set to refine the residual γ by using the structure (quenched martensite) generated by the above-described continuous annealing process. Can be secured.

Specifically, first, the steel sheet treated as described above is held for 30 to 300 seconds (t3) at a two-phase region temperature (T3) between A _r1 point and A _r3 point. When the soaking temperature (T3) exceeds the _Ar3 point, all the structures become austenite (γ). On the other hand, when the temperature falls below the _Ar1 point, γ is not generated and the desired residual γ cannot be obtained. . Similarly, even if the holding time (t3) is out of the above range, a desired tissue cannot be obtained. t3 is preferably within a range of 150 seconds or longer and 250 seconds or shorter.

Here, the A _r1 point and the A _r3 point are calculated based on the following equations, respectively.
A _r1 point = 723-29 × [Si] -11 × [Mn] + 31 × [Al]
A _r3 point = 910−203 × √ [C] + 45 × [Si] −23 × [Mn] + 95 × [Al]
In the formula, [] is the content (% by mass) of each element.

  Then, it is immersed in a galvanizing bath. The immersion conditions are not particularly limited. For example, the immersion may be performed at a temperature (T4) of about 330 to 380 ° C. for about 1 to 10 seconds (t4).

  Further, an alloying treatment may be performed as necessary. The alloying conditions are preferably maintained at a temperature (T5) of about 400 to 650 ° C. for 30 to 300 seconds (t5) (T5 × t5 is approximately 20,000 or more). When the alloying temperature (T5) is less than 350 ° C. or the alloying time (t5) is less than 30 seconds, the C concentration in the residual γ does not proceed sufficiently and the desired fine residual γ cannot be obtained. On the other hand, when T5 exceeds 650 ° C. or t5 exceeds 300 seconds, tissue including residual γ is decomposed.

  In addition, when obtaining the hot-dip Zn plating steel plate which is not alloyed, after being immersed in a galvanization bath as mentioned above, the heat pattern (within the range of about 30-100 second at the temperature of about 350-450 degreeC ( The temperature x time is generally less than 20,000).

  It is preferable to perform said process using the hot dip Zn plating equipment of a reducing atmosphere.

  In the series of steps described above, for example, a CGL (actual machine) may be used, or a CGL simulator or the like may be used.

  Hereinafter, the present invention will be described more specifically with reference to examples. However, the present invention is not limited by the following examples, and can be carried out by making appropriate modifications within a range that can be adapted to the gist of the preceding and following descriptions, all of which are within the technical scope of the present invention. include.

Example 1: Examination of component composition and production conditions using RQ (1)
A steel material having the chemical composition shown in Table 1 (remainder: iron and inevitable impurities, unit: mass%) was melted to obtain a slab, and then hot rolled. In hot rolling, heating is performed at 1250 ° C. for 30 minutes and then rolling (finish rolling temperature: 880 ° C.), cooling is performed at an average cooling rate of 50 ° C./s, holding at 650 ° C. for 30 minutes, and then air cooling. (Take-up equivalent treatment) A hot rolled steel sheet having a thickness of about 3.2 mm was obtained. Furthermore, after pickling the obtained hot-rolled steel sheet, cold rolling (rolling rate: 50%) was performed to obtain a cold-rolled steel sheet having a thickness of about 1.6 mm.

  Next, a heat treatment corresponding to the continuous annealing step (CAL) and a heat treatment corresponding to the alloying hot-dip Zn plating step (CGL) were performed under the conditions shown in Tables 2 to 3. In these tables, T1, t1, T2, and t2 during the CAL process are the same as those shown in FIG. 2A, and T3, t3, T5, and t5 during the heat treatment process equivalent to CGL are This coincides with the reference shown in FIG. In this example, the plating immersion conditions were all set at 460 ° C. for 10 seconds (in FIG. 3, T4 = 460 ° C., t4 = 10 seconds).

  The structure and mechanical properties of each steel plate thus obtained were measured as follows. The method for measuring fine residual γ is as described above.

(Organization)
The steel sheet was repeller-corroded, and the area ratio (area%) of the structure was measured by optical microscope observation (magnification 1000 times). The area ratio of residual γ and the C concentration (mass%) of residual γ were measured by an X-ray diffraction method after being ground to ¼ thickness of the steel plate, then chemically polished. Detailed measurement methods are described in ISIJ Int. Vol. 33. (19933), No. 7, p776.

(Tensile strength and elongation)
A tensile test was performed using a JIS No. 5 test piece, and tensile strength (TS) and elongation [total elongation (El)] were measured. The strain rate in the tensile test was 1 mm / sec.

(Stretch flangeability)
The stretch flangeability was evaluated based on the Steel Federation Standard JFST 1001. Specifically, after using a disk-shaped test piece having a diameter of 100 mm and a plate thickness of 2.0 mm, punching out a hole of φ10 mm with a punch, expanding the hole with a burr up using a 60 ° conical punch, and cracking The hole expansion rate (λ) at the time of penetration was measured.

  In the present invention, TS (MPa) × λ (%) ≧ 50000 and El ≧ 18% were evaluated as acceptable (examples of the present invention).

  These results are also shown in Tables 2-3. In these tables, the symbols of the steel materials used (No. in Table 1) are also shown. In Tables 2 to 3 and Tables 4 to 5 to be described later, “other” structures are other structures (pearlite, bainite, martensite, cementite, etc.) that can remain in the manufacturing process of this example.

  From Tables 2-3, it can be considered as follows.

  No. in Table 2 Nos. 1 to 14 are for examining the influence of the components in the steel, and the production conditions all satisfy the preferable requirements of the present invention.

  Of these, No. 4, 7, 10 and 14 are all examples of the present invention using a steel material in which the components in the steel satisfy the requirements of the present invention, a desired structure is obtained, and a balance between tensile strength and stretch flangeability, And excellent elongation properties.

  In contrast, no. No. 1 and No. 2 in Table 1 where the amount of C is small. This is a comparative example using 1-2, a predetermined structure (BF and residual γ) could not be obtained, and a target level of mechanical characteristics could not be ensured.

  No. No. 6 in Table 1 with a small amount of Si. This is a comparative example using No. 6, and a predetermined matrix structure was not obtained. On the other hand, no. Nos. 11-12 are Nos. This is a comparative example using 11 to 12, and a predetermined amount of matrix structure and residual γ were not obtained. Therefore, in all of these comparative examples, the mechanical properties deteriorated.

  No. in Table 3 Nos. 15 to 35 are Nos. 1 in Table 1 that satisfy the requirements of the present invention. 4 was used to examine the influence of manufacturing conditions.

  First, in Table 3, No. 15-20 is the example which changed the tempering temperature (T2) of a CAL process in the range of 100-600 degreeC.

  Of these, No. Nos. 15 to 17 are examples of the present invention produced at the tempering temperature defined in the present invention. A desired structure is obtained, and the balance between tensile strength and stretch flangeability, and elongation characteristics are excellent.

  In contrast, no. Nos. 18 to 20 are comparative examples having a high T2, and the ratio of fine residual γ is reduced. In No. 20, the generation of BF structure as a parent phase was reduced, and the mechanical properties were deteriorated.

  For reference, FIG. 16 (example of the present invention), an optical micrograph (magnification: 1000 times) is shown in FIG. 19 (comparative example) optical micrographs (magnification: 1000 times) are shown respectively. As is clear from the comparison of these photographs, in the example of the present invention shown in FIG. 4, a lot of fine residual γ is generated, whereas in the comparative example shown in FIG. It can be seen that a lot of lath-like residual γ is generated.

  No. in Table 3 21-23 are examples in which the soaking temperature (T1) of the CAL process is changed in the range of 800-950 ° C.

  No. 21 to 22 are examples of the present invention produced at a soaking temperature defined in the present invention, a desired structure is obtained, and the balance between tensile strength and stretch flangeability and the stretch properties are excellent.

In contrast, no. No. 23 is a comparative example in which T1 is lower than the temperature defined in the present invention ( _Ar3 point = 874 ° C. or higher), the ratio of fine residual γ is small, and a predetermined matrix structure cannot be obtained. The characteristics deteriorated.

  No. in Table 3 24 is an example in which the average cooling rate (CR) of the CAL process is 8 ° C./s.

  No. No. 24 is a comparative example with a slow CR, and a predetermined matrix structure was not obtained, so that the mechanical properties were deteriorated.

  No. in Table 3 26 to 27 are examples in which the soaking time (t1) of the CAL process is changed in the range of 5 to 120 seconds.

  No. No. 26 is an example of the present invention produced with a soaking time (t1) defined in the present invention, and a desired structure can be obtained and excellent in mechanical properties.

  In contrast, no. No. 27 is a comparative example with a short t1, and the ratio of fine residual γ is small and a large amount of PF or quasi-PF structure is formed, so that the mechanical properties are deteriorated.

  No. in Table 3 28 to 29 are examples in which the tempering time (t2) of the CAL process was changed in the range of 50 to 180 seconds.

  No. Reference numeral 29 is an example of the present invention manufactured with the austempering time (t2) defined in the present invention, and a desired structure is obtained and excellent in mechanical properties.

  In contrast, no. No. 28 is a comparative example having a short t2, and tempering was insufficient, and the PF structure after the CGL process was increased, so that the mechanical properties were lowered.

  No. in Table 3 30 to 34 are comparative examples in which the conditions of the CGL process are changed, and each has the following problems.

  No. 30 is a comparative example having a high soaking temperature (T3), in which the ratio of fine residual γ is small and a large amount of PF or quasi-PF structure is formed, so that the mechanical properties are deteriorated.

  No. No. 31 is a comparative example with a short soaking time (t3), a predetermined matrix structure was not obtained, and the mechanical properties were deteriorated.

  No. No. 32 is a comparative example having a long soaking time (t3). The ratio of fine residual γ is small, and a large amount of PF or quasi-PF structure is generated, so that the mechanical properties are deteriorated.

  No. No. 33 is a comparative example with a short alloying time (t5), No. 33. No. 34 is a comparative example with a long alloying time (t5), and in all cases, a large amount of PF or quasi-PF structure was formed, so that the mechanical properties were deteriorated.

Example 2: Examination of manufacturing conditions using WQ (2)
In this example, the CAL process was performed with the heat pattern shown in FIG. 2B instead of the heat pattern shown in FIG. 2A as in Example 1 described above.

  For details, see No. 1 in Table 1. After producing a cold-rolled steel sheet in the same manner as in Example 1 using the steel material No. 4, heat treatment equivalent to the continuous annealing step (CAL) and heat treatment equivalent to the alloying hot-dip Zn plating step (CGL) were performed under the conditions shown in Table 4. went. In Table 4, T1, t1, T2, and t2 during the CAL process are the same as those shown in FIG. 2B, and T3, t3, T5, and t5 during the CGL process are the same as those shown in FIG. Match. The plating immersion conditions are the same as in Example 1.

  The structure and mechanical properties of each steel plate thus obtained were measured in the same manner as in Example 1, and these results are also shown in Table 4.

  From Table 4, it can be considered as follows.

  No. in Table 4 1-3 are the examples of this invention manufactured on the manufacturing conditions prescribed | regulated by this invention, a desired structure | tissue is obtained, and it is excellent in the balance of tensile strength and stretch flangeability, and an elongation characteristic.

  In contrast, no. Nos. 4 to 6 are comparative examples in which the tempering temperature (T2) in the CAL process is high, and the ratio of fine residual γ is reduced. In No. 6, the generation of the BF structure as a parent phase was reduced, and the mechanical characteristics were deteriorated.

Example 3: Examination of TRIP steel sheet having polygonal ferrite as matrix (reference)
In this example, in TRIP steel plate (TDP steel) having polygonal ferrite as a parent phase, the influence of fine residual γ formation on stretch flangeability was examined for reference. Here, in order to generate a PF structure, the continuous annealing step shown in FIG. 6 was performed. As shown in FIG. 6, the range from the soaking temperature (T1) to the tempering temperature (T2) was cooled by changing the cooling rate.

  Specifically, No. 1 in Table 1 is used. After producing a cold-rolled steel sheet using the steel material No. 4 in the same manner as in Example 1, a continuous annealing step (CAL) was performed under the conditions shown in Table 5. In Table 5, T1, t1, T2, and t2 during the CAL process coincide with the symbols shown in FIG.

  The structure and mechanical properties of each steel plate thus obtained were measured in the same manner as in Example 1. These results are also shown in Table 5.

  As shown in Table 5, in the TBF steel using PF as the matrix, even if the ratio of the fine residual γ was increased, almost no increase in stretch flangeability was observed.

It is a graph which shows the relationship between the ratio of fine residual (gamma), and stretch flangeability ((lambda)). And the heat pattern of the continuous annealing step (CAL) a process diagram schematically showing FIG. 2 (a), FIG subjected to roll quenching (RQ), facilities to FIG. 2 (b) U _o Takuenchi (WQ) FIG. It is process drawing which shows typically the heat pattern of an alloying hot-dip Zn plating process (CGL). No. 1 in Table 3 of Example 1. 16 is an optical micrograph of 16 (Example of the present invention) (magnification 1000 times). No. 1 in Table 3 of Example 1. 19 is an optical micrograph (magnification 1000 times) of 19 (comparative example). It is process drawing which shows typically the heat pattern of the continuous annealing process (CAL) in Example 3. FIG.

Claims (2)

Components in steel are
C: 0.10 to 0.20% (meaning mass%, hereinafter the same),
Si: 0.8 to 2.5%
Mn: 1.5 to 2.5%
Al: 0.01 to 0.10%,
P: less than 0.1% (excluding 0%),
S: Less than 0.002% (excluding 0%)
Containing
The rest: satisfying iron and inevitable impurities,
The structure includes at least bainitic ferrite and retained austenite,
Area ratio for all tissues, bainitic ferrite: 70% or more,
Retained austenite: 2-20%,
Polygonal ferrite and / or quasi-polygonal ferrite: satisfying 15% or less (excluding 0%), and
A high-strength steel sheet excellent in stretch flangeability, wherein the proportion of retained austenite having an average particle size of 5 μm or less in the retained austenite is 60% or more.   Furthermore, the high-strength steel plate according to claim 1 containing 0.002% or less (not including 0%) of Ca.