CA2283924C - Dual-phase type high-strength steel sheets having high impact energy absorption properties and a method of producing the same - Google Patents
Dual-phase type high-strength steel sheets having high impact energy absorption properties and a method of producing the same Download PDFInfo
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- CA2283924C CA2283924C CA002283924A CA2283924A CA2283924C CA 2283924 C CA2283924 C CA 2283924C CA 002283924 A CA002283924 A CA 002283924A CA 2283924 A CA2283924 A CA 2283924A CA 2283924 C CA2283924 C CA 2283924C
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 138
- 239000010959 steel Substances 0.000 title claims abstract description 138
- 238000010521 absorption reaction Methods 0.000 title claims abstract description 74
- 238000000034 method Methods 0.000 title claims abstract description 27
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 56
- 238000005482 strain hardening Methods 0.000 claims abstract description 37
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 30
- 239000002131 composite material Substances 0.000 claims abstract description 8
- 241001279983 Dendroacalles sigma Species 0.000 claims abstract 2
- 238000001816 cooling Methods 0.000 claims description 43
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- 238000005096 rolling process Methods 0.000 claims description 10
- 239000000047 product Substances 0.000 claims description 9
- 238000005496 tempering Methods 0.000 claims description 9
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- 238000005272 metallurgy Methods 0.000 claims description 8
- 229910052804 chromium Inorganic materials 0.000 claims description 7
- 238000005266 casting Methods 0.000 claims description 5
- 229910052802 copper Inorganic materials 0.000 claims description 5
- 229910052750 molybdenum Inorganic materials 0.000 claims description 5
- 229910052758 niobium Inorganic materials 0.000 claims description 5
- 229910052719 titanium Inorganic materials 0.000 claims description 5
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- 239000012779 reinforcing material Substances 0.000 abstract description 2
- 230000009102 absorption Effects 0.000 description 68
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Classifications
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
- C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
- C23C2/0224—Two or more thermal pretreatments
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
- C23C2/024—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by cleaning or etching
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/008—Martensite
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- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Organic Chemistry (AREA)
- Chemical Kinetics & Catalysis (AREA)
- Physics & Mathematics (AREA)
- Thermal Sciences (AREA)
- Crystallography & Structural Chemistry (AREA)
- Heat Treatment Of Sheet Steel (AREA)
Abstract
The invention relates t.o dual-phase type high-strength steel sheets, for automobiles, which have excellent dynamic deformation properties and exhibit impact absorption properties, and are intended to be used as structural members and reinforcing materials primarily for automobiles, as well as to a method of producing them, which dual-phase type high-strength steel sheets with excellent dynamic deformation properties are characterized in that the final microstructure of the steel sheets is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction between 3% and 50% after 5% deformation of the steel sheet, wherein the difference between the quasi-static deformation strength .sigma.s when deformed in a strain rate range of 5 .times. 10 -4 - 5 .times. 10 -3 (s-1) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength ad when deformed in a strain rate range of 5 .times. 10 2 - 5 .times. 10 3 (s-1) after the aforementioned pre-deformation, i.e. (.sigma.d - .sigma.s), is at least 60 MPa, and the work hardening coefficient at 5~10% strain is at least 0.13.
Description
DESCRIPTION
DUAL-PHASE TYPE HIGH-STRENGTH STEEL SHEETS HAVING
HIGH IMPACT ENERGY ABSORPTION PROPERTIES AND
A METHOD OF PRODUCING THE SAME
Technical FiE~ld The preaent invention relates to dual-phase type high-strength steel sheets, for automobiles use, which have excellent dynamic deformation properties and exhibit excellent impact absorption properties, and are intended to be used as structural members and reinforcing materials primarily for automobiles, as well as to a method of producing them.
Background Art The app:Lications of high-strength steels have been increasing for the purpose of achieving lighter weight vehicle bodies in consideration of fuel consumption restrictions on automobiles and even more applications for high-strength steel are expected as domestic and foreign restrictions, relating to estimated impact absorption properties in automobile accidents, become rapidly more broad and strict. For example, for frontal collisions of passenger cars, the use of materials with high impact absorption properties for members known as "front side :members" can allow impact energy to be absorbed through collapse of the member, thus lessening the impact experienced by passengers.
However, conventional high-strength steels have been developed with a main view toward improving press formability, and doubts exist as to their application in terms of impact absorption properties. Prior art techniques relating to automobile steel with excellent impact absorption properties and methods of producing it have been developed which result in increased yield strength of steel sheets under high deformation speeds as an indicator of impact absorption properties, as disclosed in Japanese Unexamined Patent Publication No.
7-18372, but because the members undergo deformation
DUAL-PHASE TYPE HIGH-STRENGTH STEEL SHEETS HAVING
HIGH IMPACT ENERGY ABSORPTION PROPERTIES AND
A METHOD OF PRODUCING THE SAME
Technical FiE~ld The preaent invention relates to dual-phase type high-strength steel sheets, for automobiles use, which have excellent dynamic deformation properties and exhibit excellent impact absorption properties, and are intended to be used as structural members and reinforcing materials primarily for automobiles, as well as to a method of producing them.
Background Art The app:Lications of high-strength steels have been increasing for the purpose of achieving lighter weight vehicle bodies in consideration of fuel consumption restrictions on automobiles and even more applications for high-strength steel are expected as domestic and foreign restrictions, relating to estimated impact absorption properties in automobile accidents, become rapidly more broad and strict. For example, for frontal collisions of passenger cars, the use of materials with high impact absorption properties for members known as "front side :members" can allow impact energy to be absorbed through collapse of the member, thus lessening the impact experienced by passengers.
However, conventional high-strength steels have been developed with a main view toward improving press formability, and doubts exist as to their application in terms of impact absorption properties. Prior art techniques relating to automobile steel with excellent impact absorption properties and methods of producing it have been developed which result in increased yield strength of steel sheets under high deformation speeds as an indicator of impact absorption properties, as disclosed in Japanese Unexamined Patent Publication No.
7-18372, but because the members undergo deformation
- 2 -during the shaping process or during collision deformation, it is necessary to include a work-hardening aspect to them yield strength as an indicator of impact resistance, and this is inadequate in terms of anti-s collision sa:ety in the prior art described above.
In addil:ion, since the strain rate undergone by each location upon automobile collision reaches about 103 (s 1), consid~aration of the impact absorption properties of the mater_~als requires an understanding of the dynamic deformation properties in such a high strain rate range.
Also, high-si~rength steel sheets with excellent dynamic deformation properties are understood to be important for achieving both lighter weight and improved impact absorption p:coperties for automobiles, and recent reports have highlighted this fact. For example, the present inventors have reported on the high strain rate properties and impact energy absorption properties of high-strength thin steel sheets in CAMP-ISIJ Vol.9 (1966), pp.1.112-1115, wherein they explain that the dynamic strength at a high strain rate of 103 (s 1) increases dramatically compared to the static strength at a low strain rate speed of 10-3 (s-1), that absorption energy durin~~ crashes is increased by greater steel material strengths, that the strain rate dependency of materials depends on the structure of the steel, and that TRIP type steel (Transformation induced plasticity type steel) and dual-phase (hereunder, "DP") type steel exhibit both excellent press formability and high impact absorption properties. Also, the present inventors have already filed Japanese Patent Applications No.8-98000 and No.8-109224 relating to such a DP-type steel, among which there are proposed high-strength steel sheets with higher dynamic strength than static strength, which are suitable for achieving both lighter weights and improved impact absorption properties for automobiles, and a process for their production.
In addil:ion, since the strain rate undergone by each location upon automobile collision reaches about 103 (s 1), consid~aration of the impact absorption properties of the mater_~als requires an understanding of the dynamic deformation properties in such a high strain rate range.
Also, high-si~rength steel sheets with excellent dynamic deformation properties are understood to be important for achieving both lighter weight and improved impact absorption p:coperties for automobiles, and recent reports have highlighted this fact. For example, the present inventors have reported on the high strain rate properties and impact energy absorption properties of high-strength thin steel sheets in CAMP-ISIJ Vol.9 (1966), pp.1.112-1115, wherein they explain that the dynamic strength at a high strain rate of 103 (s 1) increases dramatically compared to the static strength at a low strain rate speed of 10-3 (s-1), that absorption energy durin~~ crashes is increased by greater steel material strengths, that the strain rate dependency of materials depends on the structure of the steel, and that TRIP type steel (Transformation induced plasticity type steel) and dual-phase (hereunder, "DP") type steel exhibit both excellent press formability and high impact absorption properties. Also, the present inventors have already filed Japanese Patent Applications No.8-98000 and No.8-109224 relating to such a DP-type steel, among which there are proposed high-strength steel sheets with higher dynamic strength than static strength, which are suitable for achieving both lighter weights and improved impact absorption properties for automobiles, and a process for their production.
- 3 -As mentioned above, although the dynamic deformation properties oi: high-strength steel sheets are understood at the high strain rates of automobile collisions, it is still unclear what properties should be maximized for automobile members with impact energy absorption properties, and on what criteria the selection of materials should be based. In addition, the automobile members are produced by press forming of steel sheets, and collision impacts are applied to these press formed members. However, high-strength steel sheets with excellent dynamic deformation properties as actual members, bas<~d on an understanding of the impact energy absorption p=roperties after such press forming, are still unknown.
For press forming of members for collision safety, a combination of excellent shape fixability, excellent stretchability (tensile strength x total elongation >_ 18,000) and ~axcellent flangeability (hole expansion ratio <_ 1.2) is desirable, but at the current time no material has provided both excellent impact absorption properties and excellent press formability.
Disclosure of the Invention The present invention has been proposed as a means of overcoming the problems described above, and provides dual-phase type high-strength steel sheets for automobiles use, which have excellent impact absorption properties and excellent dynamic deformation properties, as well as a method of producing them.
The invention further provides dual-phase type high-strength steel sheets, for automobiles, with excellent dynamic deformation properties, which are high-strength steel sheets used for automotive parts, such as front side members, and which are selected based on exact properties and standards for impact energy absorption during collisions and can reliably provide guaranteed safety, as well as a method of producing them.
The invention still further provides dual-phase type
For press forming of members for collision safety, a combination of excellent shape fixability, excellent stretchability (tensile strength x total elongation >_ 18,000) and ~axcellent flangeability (hole expansion ratio <_ 1.2) is desirable, but at the current time no material has provided both excellent impact absorption properties and excellent press formability.
Disclosure of the Invention The present invention has been proposed as a means of overcoming the problems described above, and provides dual-phase type high-strength steel sheets for automobiles use, which have excellent impact absorption properties and excellent dynamic deformation properties, as well as a method of producing them.
The invention further provides dual-phase type high-strength steel sheets, for automobiles, with excellent dynamic deformation properties, which are high-strength steel sheets used for automotive parts, such as front side members, and which are selected based on exact properties and standards for impact energy absorption during collisions and can reliably provide guaranteed safety, as well as a method of producing them.
The invention still further provides dual-phase type
- 4 -high-strength steel sheets for automobiles with excellent dynamic deformation properties, which exhibit all the properties suitable for press forming of members, including exc:ellent shape fixability, excellent stretchabilii~y and excellent flangeability, as well as a method of producing them.
The invE~ntion was devised to achieve the objects stated above by the following concrete means.
(1) A dual-phase type high-strength steel sheets having high impact energy absorption properties, characterized in that the final microstructure of the steel sheet _i.s a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite a1= a volume fraction between 3% and 50% after deformation ~3t 5% equivalent strain of the steel sheet, wherein the difference between the quasi-static deformation strength 6s when deformed in a strain rate range of 5 x 104 - 5 x 103 (s 1) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength od when deformed in ~3 strain rate range of 5 x 10z - 5 x 103 ( s 1 ) after the aforementioned pre-deformation, i.e. (od - 6s), is at least GO MPa, and the work hardening coefficient at 510% strain is at least 0.13.
(2) A dual-phase type high-strength steel sheet having high .impact energy absorption properties, characterized in that the final microstructure of the steel sheet is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, wherein the average value odyn (MPa) of the deformation stress in th~~ range of 310% of equivalent strain when deformed in .a strain rate range of 5 x 102 - 5 x 103
The invE~ntion was devised to achieve the objects stated above by the following concrete means.
(1) A dual-phase type high-strength steel sheets having high impact energy absorption properties, characterized in that the final microstructure of the steel sheet _i.s a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite a1= a volume fraction between 3% and 50% after deformation ~3t 5% equivalent strain of the steel sheet, wherein the difference between the quasi-static deformation strength 6s when deformed in a strain rate range of 5 x 104 - 5 x 103 (s 1) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength od when deformed in ~3 strain rate range of 5 x 10z - 5 x 103 ( s 1 ) after the aforementioned pre-deformation, i.e. (od - 6s), is at least GO MPa, and the work hardening coefficient at 510% strain is at least 0.13.
(2) A dual-phase type high-strength steel sheet having high .impact energy absorption properties, characterized in that the final microstructure of the steel sheet is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, wherein the average value odyn (MPa) of the deformation stress in th~~ range of 310% of equivalent strain when deformed in .a strain rate range of 5 x 102 - 5 x 103
- 5 -(s-1), after pre-deformation of more than 0~ and less than or equa7_ to 10~ of equivalent strain, satisfies the inequality: cfdyn ? 0.766 x TS + 250 as expressed in terms of the tensi7_e strength TS (MPa) in the quasi-static tensile test as measured in a strain rate range of 5 x 104 - 5 x 103 (s-1) prior to pre-deformation, and the work hardening coefficient at 5~10~ strain is at least 0.13.
(3) A dual-phase type high-strength steel sheet having high impact energy absorption properties according to (1) or (2) above, characterized in that the ratio between the yield strength YS(0) and the tensile strength TS'(5) in them tensile test after pre-deformation at 50 of equivalent si~rain or after further bake hardening treatment (BH treatment) satisfies the inequality YS(0)/TS'(5) <- 0.7, and also satisfies the inequality:
yield strength YS(0) x work hardening coefficient ? 70.
(4) A dual-phase type high-strength steel sheet having high :impact energy absorption properties according to any of (1), (2) or (3) above, characterized in that the average grain size of the martensite is 5 ~m or less, and the average grain size of the ferrite is 10 ~.m or less.
(5) A dual-phase type high-strength steel sheet having high .impact energy absorption properties according to any of (1), (2), (3) or (4) above, characterized by satisfying t:he inequality: tensile strength (MPa) x total elongation (~) ? 18,000, and by satisfying the inequality: hole expansion ratio (d/do) >_ 1.2.
(3) A dual-phase type high-strength steel sheet having high impact energy absorption properties according to (1) or (2) above, characterized in that the ratio between the yield strength YS(0) and the tensile strength TS'(5) in them tensile test after pre-deformation at 50 of equivalent si~rain or after further bake hardening treatment (BH treatment) satisfies the inequality YS(0)/TS'(5) <- 0.7, and also satisfies the inequality:
yield strength YS(0) x work hardening coefficient ? 70.
(4) A dual-phase type high-strength steel sheet having high :impact energy absorption properties according to any of (1), (2) or (3) above, characterized in that the average grain size of the martensite is 5 ~m or less, and the average grain size of the ferrite is 10 ~.m or less.
(5) A dual-phase type high-strength steel sheet having high .impact energy absorption properties according to any of (1), (2), (3) or (4) above, characterized by satisfying t:he inequality: tensile strength (MPa) x total elongation (~) ? 18,000, and by satisfying the inequality: hole expansion ratio (d/do) >_ 1.2.
(6) A dual-phase type high-strength steel sheet having high impact energy absorption properties according to any of (1), (2), (3), (4) or (5) above, characterized in that the plastic deformation (T) by either or both a tempering rolling and a tension leveller satisfies the following inequality.
2.5 fYS(0)/TS'(5) - 0.5} + 15 >_ T ? 2.5 ~YS(0)/~'S'(5) - 0.5} + 0.5
2.5 fYS(0)/TS'(5) - 0.5} + 15 >_ T ? 2.5 ~YS(0)/~'S'(5) - 0.5} + 0.5
(7) The dual-phase type high-strength steel sheet having high impact energy absorption~properties according to the invention is also a dual-phase type high-strength steel sheet with excellent dynamic deformation properties according to (1) to (6) above, characterized in that the chemical compositions, in terms of weight percentage, C
at 0.02~0.25~;, either or both Mn and Cr at a total of 0.15~3.5~, one or more from among Si, A1 and P at a total of 0.02~4.0~, if necessary one or more from among Ni, Cu and Mo at a total of no more than 3.5~, one or more from among Nb, Ti and V at no more than 0.30, and either or both Ca and REM at 0.0005~0.01~ for Ca and 0.005~0.05~
for REM, with the remainder Fe as the primary component.
at 0.02~0.25~;, either or both Mn and Cr at a total of 0.15~3.5~, one or more from among Si, A1 and P at a total of 0.02~4.0~, if necessary one or more from among Ni, Cu and Mo at a total of no more than 3.5~, one or more from among Nb, Ti and V at no more than 0.30, and either or both Ca and REM at 0.0005~0.01~ for Ca and 0.005~0.05~
for REM, with the remainder Fe as the primary component.
(8) The dual-phase type high-strength steel sheet having high impact energy absorption properties according to the invention is also a dual-phase type high-strength steel sheet with excellent dynamic deformation properties according to (1) to (7) above, characterized in that one or more from among B (__<0.01), S (__<0.01%) and N (<_0.02~) are further added if necessary to the steel.
(9) The method of producing a dual-phase type high-strength hot--rolled steel sheet having high impact energy absorption properties according to the invention is a method of producing a dual-phase type high strength hot-rolled steel sheet with excellent dynamic deformation properties according to (1) to (8) above, characterized in that after a continuous casting slab is fed directly from casting to a hot rolling step, or is hot rolled upon repeating afi~er momentary cooling, it is subjected to hot rolling at a finishing temperature of Ar3 - 50°C to Ar3 +
120°C, cooled at an average cooling rate of more than 5°C/sec in a run-out table, and then coiled at a temperature of no greater than 350°C; and
120°C, cooled at an average cooling rate of more than 5°C/sec in a run-out table, and then coiled at a temperature of no greater than 350°C; and
(10) a method of producing a dual-phase high-strength hot-rolled steel sheet having high impact energy -absorption properties according to (9) above, characterized in that at the finishing temperature for hot rolling in a range of Ar3 - 50°C~to Ar3 + 120°C, the hot rolling _~s carried out so that the metallurgy parameter A satisfies inequalities (1) and (2) below, the subsequent average cooling rate in the run-out table is at least 5°C/sec, and the coiling is accomplished so that the relationship between the above-mentioned metallurgy parameter A and the coiling temperature (CT) satisfies inequality (:3) below.
9 <_ logA <_ 18 (1) oT <_ 21 x logA - 61 (2) CT <- 6 :.c logA + 242 ( 3 )
9 <_ logA <_ 18 (1) oT <_ 21 x logA - 61 (2) CT <- 6 :.c logA + 242 ( 3 )
(11) The method of producing a dual-phase type high-strength cold rolled steel sheet having high impact energy abosorption properties according to the invention is a method of producing a dual-phase type high-strength cold rolled steel sheet with excellent dynamic deformation properties according to (1) to (8) above, characterized in that after a continuous cast slab is fed directly from casting to a hot rolling step, or is hot rolled upon :reheating after momentary cooling, it is hot rolled, the :hot-rolled and subsequently coiled steel sheet is cold-rolled after acid pickling, and during annealing in a continuous annealing step for preparation of the final product, it is heated to a temperature between Aci and Acz and subjected to the annealing while held in this temperature range for at least 10 seconds, and then cooled at a cooling rate of more than 5°C/sec;
and
and
(12) a :method according to (11) above for producing a dual-phase type high-strength cold rolled steel sheet having high impact energy absorption properties according to (1) to (8) above, characterized in that in the continuous annealing step, the cold rolled steel sheet is heated to a temperature (To) between Acl and Ac3 and _ g _ subjected to the annealing while held in this temperature range for at least 10 seconds, and for subsequent cooling, it is cooled to a secondary~cooling start' temperature (Tq) in the range of 550°C-To at a primary cooling rate of 110°C/sec and then cooled to a secondary cooling end temperature (Te) which is no higher than Tem determined by the chemical compositions and annealing temperature (To), at a secondary cooling rate of 10200°C/sec.
Brief Descrix~tion of the Drawinas Fig. 1 is a graph showing the relationship between the absorption energy (Eab) of a shaped member during collision and the material strength (S), according to the invention.
Fig. 2 _~s a perspective view of a shaped member for measurement of impact absorption energy for Fig. 1.
Fig. 3 is a graph showing the relationship between the work hardening coefficient and dynamic energy absorption for a steel sheet.
Fig. 4 .is a graph showing the relationship between the yield strength x work hardening coefficient and the dynamic energy absorption for a steel sheet.
Fig. 5 .is a general view of a "hat model" used in the impact crush test method relating to Figs. 3 and 4.
Fig. 6 is a cross-sectional view of the shape of the test piece of Fig. 5.
Fig. 7 is a schematic view of the impact crush test method relating to Figs. 3-6.
Fig. 8 is a graph showing the relationship between TS and the difference between the average value 6dyn of the deformation stress in the range of 310% of equivalent strain when deformed in a strain rate range of 5 x lOZ - 5 x 103 (1/S) and TS, as an index of the impact energy absorption property upon collision, according to the invention.
Fig. 9 is a graph showing the change in the static/dynam.ic ratio with tempered rolling for an example of the invention and a comparative example.
Fig. 10 is a graph showing the relationship between oT and the m~atallurgy parameter A for a hot-rolling step according to the invention.
Fig. 11 is a graph showing the relationship between the coiling temperature and the metallurgy parameter A
for a hot-rolling step according to the invention.
Fig. 12 is a graph showing the annealing cycle for continuous annealing according to the invention.
Best Mode fo:r Carrying Out the Invention Impact .absorbing members such as front side members of automobiles are produced by bending and press forming of steel sheets. Because impacts during automobile collisions are absorbed by such members which have undergone press forming, they must have high impact absorption properties even after having undergone the pre-deformation corresponding to the press forming. At the current time, however, no attempt has been made to obtain high-strength steel sheets with excellent impact absorption properties as actual members, with consideration of both the increase in the deformation stress by press forming and the increase in deformation stress due to a higher strain rate, as was mentioned above.
As a result of much experimentation and research with the aim of achieving this purpose, the present inventors have found that steel sheets with a dual-phase (DP) structure are ideal as high-strength steel sheets with excellent impact absorption properties for actual members which are press formed as described above. It was demonstrated that such steel sheets with a dual-phase microstructure, which is a composite microstructure wherein the dominating phase is a ferrite phase responsible for the increase in deformation resistance by an increased. strain rate, and the second phase includes a hard martensite phase, have excellent dynamic deformation properties. That is, it was found that high dynamic deformation properties are exhibited when the microstructure of the final steel sheets is a composite structure wherein the dominating phase is ferrite and another low temperature product phase includes a hard martensite phase at a volume fraction of 350% after deformation at 5% equivalent strain of the steel sheet.
Concerning the volume fraction of 350% for the hard martensite phase, since high-strength steel sheets and even steel sheets with high dynamic deformation properties cannot be obtained if the martensite phase is less than 3%,. the volume fraction of the martensite phase must be at least 3%. Also, if the martensite phase exceeds 50%, this results in a smaller volume fraction of the ferrite phase responsible for greater deformation resistance due to increased deformation speed, making it impossible to obtain steel sheets with excellent dynamic deformation properties compared to static deformation strength while also hindering press formability, and therefore it was found that the volume fraction of the martensite phase must be 350%.
The present inventors then pursued experimentation and research based on these findings and, as a result, found that a:Lthough the degree of pre-deformation corresponding to press forming of impact absorbing members such as front side members sametimes reaches a maximum of over 20%, depending on the location, the majority are locations with 0%~10% of equivalent strain, and that by understanding the effect of pre-deformation in this range, it is possible to estimate the behavior of the member a5 a whole after pre-deformation.
Consequently, according to the invention, a deformation of from 0% to 10% of equivalent strain was selected as the amount of pre-deformation applied to members during press forming.
Fig. 1 is a graph showing the relationship between the absorption energy (Eab) of a press formed member during collision and the material strength (S), for the different st~ael types shown in Table 5, according to an example to bra described later. The material strength S
is the tensile strength (TS) according to the common tensile test. The member absorption energy (Eab) is the absorption energy in the lengthwise direction (direction of the arrow) along a press formed member such as shown in Fig. 2, upon collision with a 400 kg mass weight at a speed of 15 m/sec, to a crushing degree of 100 mm. The shaped member in Fig. 2 consists of a 2.0 mm-thick steel sheet formed into a hat-shaped section 1 with a steel sheet 2 of the same thickness and the same type of steel, joined together by spot welding, the hat-shaped section 1 having a corner radius of 2 mm, and with spot welding points indicated by 3.
From Fig. 1 it is seen that the member absorption energy (Eab) tends to increase with the strength of materials under normal tensile testing, though with considerable variation. Here, the materials in Fig. 1 were subjected to pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and then the static deformation strength 6s when deformed in a strain rate range of 5 x 10 4 - 5 x 10 3 ( s 1 ) and the dynamic deformation strength ad when deformed in a strain rate range of 5 x lOZ - 5 x 103 ( s 1 ) after the pre-deformation, were measured. As a result, a classification was possible based on (~d - mss). The symbols plotted in Fig. 1 were as follows:
o: (ad - 6s) < 60 MPa with any pre-deformation of more than 0% and less than or equal to 10%;
~: 60 M.Pa <_ (6d - as) with any :pre-deformation in the above range, and 60 MPa _< (6d - as) < 80 MPa with pre-deformation of 5%;
1: 60 M.Pa <_ (ad - cs) with any pre-deformation in the above range, and 80 MPa <_ (~d - os) < 100 MPa with. pre-deformation of 5%;
~: 60 MF~a -< (ad - os) with any pre-deformation in the above range, and 100 MPa <- (od - 6s) with pre-deformation of 5%.
Also, when 60 MPa <_ (ad - 6s) with any pre-y deformation i.n the range of more than 0% and less than or equal to 10% of equivalent strain, the values for member absorption energy (Eab) during collision was equal to or greater than the values predicted from the material strength S, thus indicating steel sheets with excellent dynamic deformation properties as impact absorbing members for collision. These predicted values are those shown in the curve in Fig. 1, represented by Eab =
0.062S°'8. Consequently, (od - os) must be at least 60 MPa.
For improved impact absorption properties, it is basically important to increase the work hardening coefficient, specifically to at least 0.13, and preferably at: least 0.16; by controlling the yield strength and the work hardening coefficient to specified ranges it is possible to achieve excellent impact absorption pi_-operties, and for improved press formability it is effective to design the volume percentage and particle size of the martensite to within a specified range.
Fig. 3 shows the relationship between the work hardening coefficient of a steel sheet and the dynamic energy absorption which indicates the member impact absorption properties, for a class of: materials with the same yield strength. Here it is shown that increased work hardening coefficients of the steel sheets result in improved member impact absorption properties (dynamic energy absorption), and that the work hardening coefficient of a steel sheet can properly indicate the member impact absorption properties so long as the yield strength class is the same. Also, when the yield strengths differ, as shown in Fig. 4, the yield strength
Brief Descrix~tion of the Drawinas Fig. 1 is a graph showing the relationship between the absorption energy (Eab) of a shaped member during collision and the material strength (S), according to the invention.
Fig. 2 _~s a perspective view of a shaped member for measurement of impact absorption energy for Fig. 1.
Fig. 3 is a graph showing the relationship between the work hardening coefficient and dynamic energy absorption for a steel sheet.
Fig. 4 .is a graph showing the relationship between the yield strength x work hardening coefficient and the dynamic energy absorption for a steel sheet.
Fig. 5 .is a general view of a "hat model" used in the impact crush test method relating to Figs. 3 and 4.
Fig. 6 is a cross-sectional view of the shape of the test piece of Fig. 5.
Fig. 7 is a schematic view of the impact crush test method relating to Figs. 3-6.
Fig. 8 is a graph showing the relationship between TS and the difference between the average value 6dyn of the deformation stress in the range of 310% of equivalent strain when deformed in a strain rate range of 5 x lOZ - 5 x 103 (1/S) and TS, as an index of the impact energy absorption property upon collision, according to the invention.
Fig. 9 is a graph showing the change in the static/dynam.ic ratio with tempered rolling for an example of the invention and a comparative example.
Fig. 10 is a graph showing the relationship between oT and the m~atallurgy parameter A for a hot-rolling step according to the invention.
Fig. 11 is a graph showing the relationship between the coiling temperature and the metallurgy parameter A
for a hot-rolling step according to the invention.
Fig. 12 is a graph showing the annealing cycle for continuous annealing according to the invention.
Best Mode fo:r Carrying Out the Invention Impact .absorbing members such as front side members of automobiles are produced by bending and press forming of steel sheets. Because impacts during automobile collisions are absorbed by such members which have undergone press forming, they must have high impact absorption properties even after having undergone the pre-deformation corresponding to the press forming. At the current time, however, no attempt has been made to obtain high-strength steel sheets with excellent impact absorption properties as actual members, with consideration of both the increase in the deformation stress by press forming and the increase in deformation stress due to a higher strain rate, as was mentioned above.
As a result of much experimentation and research with the aim of achieving this purpose, the present inventors have found that steel sheets with a dual-phase (DP) structure are ideal as high-strength steel sheets with excellent impact absorption properties for actual members which are press formed as described above. It was demonstrated that such steel sheets with a dual-phase microstructure, which is a composite microstructure wherein the dominating phase is a ferrite phase responsible for the increase in deformation resistance by an increased. strain rate, and the second phase includes a hard martensite phase, have excellent dynamic deformation properties. That is, it was found that high dynamic deformation properties are exhibited when the microstructure of the final steel sheets is a composite structure wherein the dominating phase is ferrite and another low temperature product phase includes a hard martensite phase at a volume fraction of 350% after deformation at 5% equivalent strain of the steel sheet.
Concerning the volume fraction of 350% for the hard martensite phase, since high-strength steel sheets and even steel sheets with high dynamic deformation properties cannot be obtained if the martensite phase is less than 3%,. the volume fraction of the martensite phase must be at least 3%. Also, if the martensite phase exceeds 50%, this results in a smaller volume fraction of the ferrite phase responsible for greater deformation resistance due to increased deformation speed, making it impossible to obtain steel sheets with excellent dynamic deformation properties compared to static deformation strength while also hindering press formability, and therefore it was found that the volume fraction of the martensite phase must be 350%.
The present inventors then pursued experimentation and research based on these findings and, as a result, found that a:Lthough the degree of pre-deformation corresponding to press forming of impact absorbing members such as front side members sametimes reaches a maximum of over 20%, depending on the location, the majority are locations with 0%~10% of equivalent strain, and that by understanding the effect of pre-deformation in this range, it is possible to estimate the behavior of the member a5 a whole after pre-deformation.
Consequently, according to the invention, a deformation of from 0% to 10% of equivalent strain was selected as the amount of pre-deformation applied to members during press forming.
Fig. 1 is a graph showing the relationship between the absorption energy (Eab) of a press formed member during collision and the material strength (S), for the different st~ael types shown in Table 5, according to an example to bra described later. The material strength S
is the tensile strength (TS) according to the common tensile test. The member absorption energy (Eab) is the absorption energy in the lengthwise direction (direction of the arrow) along a press formed member such as shown in Fig. 2, upon collision with a 400 kg mass weight at a speed of 15 m/sec, to a crushing degree of 100 mm. The shaped member in Fig. 2 consists of a 2.0 mm-thick steel sheet formed into a hat-shaped section 1 with a steel sheet 2 of the same thickness and the same type of steel, joined together by spot welding, the hat-shaped section 1 having a corner radius of 2 mm, and with spot welding points indicated by 3.
From Fig. 1 it is seen that the member absorption energy (Eab) tends to increase with the strength of materials under normal tensile testing, though with considerable variation. Here, the materials in Fig. 1 were subjected to pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and then the static deformation strength 6s when deformed in a strain rate range of 5 x 10 4 - 5 x 10 3 ( s 1 ) and the dynamic deformation strength ad when deformed in a strain rate range of 5 x lOZ - 5 x 103 ( s 1 ) after the pre-deformation, were measured. As a result, a classification was possible based on (~d - mss). The symbols plotted in Fig. 1 were as follows:
o: (ad - 6s) < 60 MPa with any pre-deformation of more than 0% and less than or equal to 10%;
~: 60 M.Pa <_ (6d - as) with any :pre-deformation in the above range, and 60 MPa _< (6d - as) < 80 MPa with pre-deformation of 5%;
1: 60 M.Pa <_ (ad - cs) with any pre-deformation in the above range, and 80 MPa <_ (~d - os) < 100 MPa with. pre-deformation of 5%;
~: 60 MF~a -< (ad - os) with any pre-deformation in the above range, and 100 MPa <- (od - 6s) with pre-deformation of 5%.
Also, when 60 MPa <_ (ad - 6s) with any pre-y deformation i.n the range of more than 0% and less than or equal to 10% of equivalent strain, the values for member absorption energy (Eab) during collision was equal to or greater than the values predicted from the material strength S, thus indicating steel sheets with excellent dynamic deformation properties as impact absorbing members for collision. These predicted values are those shown in the curve in Fig. 1, represented by Eab =
0.062S°'8. Consequently, (od - os) must be at least 60 MPa.
For improved impact absorption properties, it is basically important to increase the work hardening coefficient, specifically to at least 0.13, and preferably at: least 0.16; by controlling the yield strength and the work hardening coefficient to specified ranges it is possible to achieve excellent impact absorption pi_-operties, and for improved press formability it is effective to design the volume percentage and particle size of the martensite to within a specified range.
Fig. 3 shows the relationship between the work hardening coefficient of a steel sheet and the dynamic energy absorption which indicates the member impact absorption properties, for a class of: materials with the same yield strength. Here it is shown that increased work hardening coefficients of the steel sheets result in improved member impact absorption properties (dynamic energy absorption), and that the work hardening coefficient of a steel sheet can properly indicate the member impact absorption properties so long as the yield strength class is the same. Also, when the yield strengths differ, as shown in Fig. 4, the yield strength
- 13 -x work hardening coefficient can be an indicator of the member impact absorption properties. While the work hardening coefficient was expressed in terms of an n value of 5~~:10~ strain in consideration of the strain undergone by members during press forming, from the viewpoint of improving the dynamic energy absorption, work hardening coefficients of under 5~ strain or work hardening coE~fficients of even more than 10~ strain may be preferred.
The dynamic energy absorptions for members shown in Fig. 3 and F:ig. 4 were determined in the following manner. Specifically, the steel sheet was shaped into the member shape shown in Fig. 6 (corner R = 5 mm) and spot welded at 35 mm pitch using an electrode with a tip radius of 5..'i mm at a current of 0.9 times the expulsion current, and then after baking and painting treatment at 170°C x 20 minutes, an approximately 150 Kg falling weight was dropped from a height of about 10 m to crush the member in its lengthwise direction, and the displacement work where displacement = 0-150 mm is calculated from the area of the corresponding load displacement diagram to determine the dynamic energy absorption. A schematic illustration of this test method is shown in :Eig. 7. In Fig. 5, 4 is a worktop, 5 is a test piece a:nd 6 is a spot welding section.
In Fig. 6, 7 is a hat-shaped test piece and 8 is a spot welding section. In Fig. 7, 9 is a worktop, 10 is a test piece, 11 is a falling weight (150 kg), 12 is a frame, and 13 is a shock absorber. The work hardening coefficient and yield strength of each steel sheet was determined in the following manner. The steel sheet was shaped into a JIS-#5 test piece (gauge length: 50 mm, parallel width: 25 mm), subjected to tensile test at a strain rate of 0.001 (s1) to determine the yield strength and work hardening coefficient (n value at 5%~10% strain). The steel sheet used had a sheet
The dynamic energy absorptions for members shown in Fig. 3 and F:ig. 4 were determined in the following manner. Specifically, the steel sheet was shaped into the member shape shown in Fig. 6 (corner R = 5 mm) and spot welded at 35 mm pitch using an electrode with a tip radius of 5..'i mm at a current of 0.9 times the expulsion current, and then after baking and painting treatment at 170°C x 20 minutes, an approximately 150 Kg falling weight was dropped from a height of about 10 m to crush the member in its lengthwise direction, and the displacement work where displacement = 0-150 mm is calculated from the area of the corresponding load displacement diagram to determine the dynamic energy absorption. A schematic illustration of this test method is shown in :Eig. 7. In Fig. 5, 4 is a worktop, 5 is a test piece a:nd 6 is a spot welding section.
In Fig. 6, 7 is a hat-shaped test piece and 8 is a spot welding section. In Fig. 7, 9 is a worktop, 10 is a test piece, 11 is a falling weight (150 kg), 12 is a frame, and 13 is a shock absorber. The work hardening coefficient and yield strength of each steel sheet was determined in the following manner. The steel sheet was shaped into a JIS-#5 test piece (gauge length: 50 mm, parallel width: 25 mm), subjected to tensile test at a strain rate of 0.001 (s1) to determine the yield strength and work hardening coefficient (n value at 5%~10% strain). The steel sheet used had a sheet
- 14 -thickness of 1.2 mm and the steel sheet composition contained C at 0.020.25 wt~, either or both Mn and Cr at a total of 0.153.5 wt% and one or more of Si, A1 and P
at a total o.f 0.024.0 wt%, with the remainder Fe as the main component .
Fig. 8 .is a graph showing the relationship between the average value adyn of the deformation stress in the range of 310% of equivalent strain when deformed in a strain rate .range of 5 x 102 - 5 x 103 ( s 1 ) and the static material strength (TS), as an index of the impact energy absorption property upon collision according to the invention, where the static material strength (TS) is the tensile strength (TS: MPa) in the static tensile test as measured .in a strain rate range of 5 x 104 - 5 x 103 ( s-1 ) .
As mentioned above, impact absorbing members such as front side members have a hat-shaped cross-sectional shape, and a;s a result of analysis of deformation of such members upon crushing by high-speed collision, the present inventors have found that despite deformation proceeding up to a high maximum strain of over 40%, at least 70% of the total absorption energy is absorbed in a strain range of 10% or lower in a high-speed stress-strain diagram. Therefore, the dynamic deformation resistance with high-speed deformation at 10% or lower was used as the index of the high-speed collision energy absorption property. In particular, since the amount of strain in the range of 310% is most important, the index used for the impact energy absorption property was the average stress: adyn in the range of 310% of equivalent strain when deformed in a strain rate range of 5 x 10z -5 x 103 (s1) high-speed tensile deformation.
The average stress: edyn of 310% upon high-speed deformation generally increases with increasing static tensile strength {maximum stress (TS: MPa) in a static tensile test measured in a stress rate range of 5 x
at a total o.f 0.024.0 wt%, with the remainder Fe as the main component .
Fig. 8 .is a graph showing the relationship between the average value adyn of the deformation stress in the range of 310% of equivalent strain when deformed in a strain rate .range of 5 x 102 - 5 x 103 ( s 1 ) and the static material strength (TS), as an index of the impact energy absorption property upon collision according to the invention, where the static material strength (TS) is the tensile strength (TS: MPa) in the static tensile test as measured .in a strain rate range of 5 x 104 - 5 x 103 ( s-1 ) .
As mentioned above, impact absorbing members such as front side members have a hat-shaped cross-sectional shape, and a;s a result of analysis of deformation of such members upon crushing by high-speed collision, the present inventors have found that despite deformation proceeding up to a high maximum strain of over 40%, at least 70% of the total absorption energy is absorbed in a strain range of 10% or lower in a high-speed stress-strain diagram. Therefore, the dynamic deformation resistance with high-speed deformation at 10% or lower was used as the index of the high-speed collision energy absorption property. In particular, since the amount of strain in the range of 310% is most important, the index used for the impact energy absorption property was the average stress: adyn in the range of 310% of equivalent strain when deformed in a strain rate range of 5 x 10z -5 x 103 (s1) high-speed tensile deformation.
The average stress: edyn of 310% upon high-speed deformation generally increases with increasing static tensile strength {maximum stress (TS: MPa) in a static tensile test measured in a stress rate range of 5 x
- 15 -4 - 5 x 10 ' ( s-1 ) } of the steel material prior to pre-deformation or baking treatment. Consequently, increasing the static tensile strength (which is synonymous w:Lth the static material strength) of the 5 steel material directly contributes to an improved impact energy absorption property of the member. However, increased strength of the steel results in poorer press formability :into members, making it difficult to obtain members with the necessary shapes. Consequently, steels 10 having a high 6dyn with the same tensile strength TS are preferred. :It was found that, based on this relationship, steel sheets wherein the average value cdyn (MPa) of the deformation stress in the range of 3~10~ of equivalent strain when deformed in a strain rate range of 5 x 102 - 5 ~: 103 ( s 1 ) , after pre-deformation of more than 0~ and .Less than or equal to 10~ of equivalent strain satisfies the inequality: adyn ? 0.766 x TS + 250 as expressed in terms of the tensile strength (TS: MPa) in the static tensile test as measured in a strain rate range of 5 x 10 4 - 5 x 10 3 ( s-1 ) prior to pre-deformation, have higher impact energy absorption properties as actual members compared to other steels, and that the impact energy absorption property is improved without increasing the overall weight of the member, making it possible to provide high-strength steel sheets with high dynamic deformation resistance.
Also, although the details are still unclear, it has been discovered that steel sheets with excellent dynamic deformation properties can be obtained when, as shown in Fig. 9, YS(0)/TS'(5) is no greater than 0.7, which amount is dependent on the initial microstructure, the amount of solid solution elements in the low temperature product phase other than the martensite phase and the main ferrite phase, and the deposited state of carbides, nitrides and carbonitrides. Here, YS(0) is the yield strength, and TS'(5) i_s the tensile strength (TS') in the
Also, although the details are still unclear, it has been discovered that steel sheets with excellent dynamic deformation properties can be obtained when, as shown in Fig. 9, YS(0)/TS'(5) is no greater than 0.7, which amount is dependent on the initial microstructure, the amount of solid solution elements in the low temperature product phase other than the martensite phase and the main ferrite phase, and the deposited state of carbides, nitrides and carbonitrides. Here, YS(0) is the yield strength, and TS'(5) i_s the tensile strength (TS') in the
- 16 -static tensile test with pre-deformation at 5% of equivalent si~rain or after further bake hardening treatment (BH treatment). It was also demonstrated that steel sheets with even more excellent dynamic deformation properties can be obtained when the yield strength: YS(0) x work hardening coefficient is at least 70.
Furthermore, it is known that dynamic deformation strength is usually expressed in the form of the power of the static tensile strength, and as the static tensile strength inc:ceases, the difference between the dynamic deformation atrength and the static deformation strength decreases. l3owever, a small difference between the dynamic deformation strength and the static deformation strength will mean that no greater improvement in the impact absorption properties can be expected. From this standpoint, it is preferred for the value of (od - os) to be in a range which satisfies the following inequality, ( 6d - as ) ? 4 . 1 x os°'a - as .
The microstructure of a steel sheet according to the invention will now be described in detail. As already mentioned, the martensite is at a volume fraction of 350%, and preferably 330%. The average grain size of the martensite is preferably no greater than 5 Vim, and the average grain size of the ferrite is preferably no greater than 10 Vim. That is, the martensite is hard, and contributes to a decrease in the yield ratio and an improvement in the work hardening coefficient, by producing a mobile dislocations primarily in adjacent ferrite grains; however, by satisfying the restrictions mentioned above it is possible to disperse fine martensite in the steel, so that the improvement in the properties spreads throughout the entire steel sheet. In addition, this dispersion of fine martensite in the steel can help to avoid deterioration in the hole expansion ratio and tensile strength x total elongation, which is an adverse effect of the hard martensite. Also, because
Furthermore, it is known that dynamic deformation strength is usually expressed in the form of the power of the static tensile strength, and as the static tensile strength inc:ceases, the difference between the dynamic deformation atrength and the static deformation strength decreases. l3owever, a small difference between the dynamic deformation strength and the static deformation strength will mean that no greater improvement in the impact absorption properties can be expected. From this standpoint, it is preferred for the value of (od - os) to be in a range which satisfies the following inequality, ( 6d - as ) ? 4 . 1 x os°'a - as .
The microstructure of a steel sheet according to the invention will now be described in detail. As already mentioned, the martensite is at a volume fraction of 350%, and preferably 330%. The average grain size of the martensite is preferably no greater than 5 Vim, and the average grain size of the ferrite is preferably no greater than 10 Vim. That is, the martensite is hard, and contributes to a decrease in the yield ratio and an improvement in the work hardening coefficient, by producing a mobile dislocations primarily in adjacent ferrite grains; however, by satisfying the restrictions mentioned above it is possible to disperse fine martensite in the steel, so that the improvement in the properties spreads throughout the entire steel sheet. In addition, this dispersion of fine martensite in the steel can help to avoid deterioration in the hole expansion ratio and tensile strength x total elongation, which is an adverse effect of the hard martensite. Also, because
- 17 -it is possible to reliably achieve work hardening coefficient ._ 0.130, tensile strength x total elongation
18,000 and hole expansion ratio >- 1.2, it is thereby' possible to improve the impact absorption properties and press formability.
With a martensite volume fraction of less than 3%, the yield ratio becomes larger while the press formed member cannot: exhibit an excellent work hardening property (work hardening coefficient >- 0.130) after it has undergone collision deformation, and since the deformation resistance (load) stays at a low level, and the dynamic energy absorption is low preventing improvement in the impact absorption properties. On the other hand, with a martensite volume fraction of greater than 50%, the yield ratio becomes larger while work hardening coefficient is reduced, and deterioration also occurs in the tensile strength x total elongation and the hole expansion ratio. From the standpoint of press formability, the volume fraction of the martensite is preferred to be no greater than 30%.
Also, the ferrite is present at a volume fraction of preferably at; least 50%, and more preferably at least 70%, and its average grain size (mean circle equivalent diameter) is preferably no greater than 10 Vim, and more preferably no greater than 5 Vim, with the martensite preferably adjacent to the ferrite. This aids the fine dispersion oj= the martensite in the ferrite matrix, while effectively extending the property-improving effect, beyond simply a local effect, to the entire steel sheet, favorably acting to prevent the adverse effects of the martensite. The structure of the remainder present with the martensite and ferrite may be a mixed structure comprising a combination of one or. more from among pearlite, ba:inite, retained 'y, etc., and although primarily ba:inite is preferred in cases which require hole expansion properties, since retained y undergoes work-induced transformation into martensite by press forming, experimental results have shown that including retained aust:enite prior to press forming has an-effect even in preferred small amounts (5% or less).
Also, from the standpoint of impact absorption properties and press formability it is preferred for the ratio of the martensite and ferrite particle sizes to be no greater than 0.6, and the ratia of the hardnesses to be at least .1.5.
The rest:rictions on the values for the chemical components of-_ dual-phase type high-strength steel sheets with excellent dynamic deformation properties according to the inveni:ion, and the reasons for those restrictions, will now be Explained.
Dual-phase type high-strength steel sheets with excellent dynamic deformation properties which are used according to the invention are steel sheets containing the following chemical compositions, in terms of weight percentage: c~ at 0.02~0.25%, either or both Mn and Cr at a total of 0.15~3.5%, one or more from among Si, Al and P
at a total o:E 0.02~4.0%, if necessary also one or more from among N:i, Cu and Mo at a total of no more than 3.5%, one or more :from among Nb, Ti and V at no more than 0.30%, and either or both Ca and REM at 0.0005~0.01% for Ca and 0.0050.05% for REM, with the remainder Fe as the primary component. They are also dual-phase type high strength ste~=1 sheets with excellent dynamic deformation properties which contain, if necessary, one or more from among B (<_0.01), S (__<0.01%) and N (<_0.02%). These chemical components and their contents (percent by weight) will now be discussed.
C: C is the element which most strongly affects the microstructure of the steel sheet, and if its content is too low it will become difficult to obtain martensite with the desired amount and strength. Addition in too great an amount leads to unwanted carbide precipitation, inhibited increase in deformation resistance at higher strain rates and overly high strength, as well as poor
With a martensite volume fraction of less than 3%, the yield ratio becomes larger while the press formed member cannot: exhibit an excellent work hardening property (work hardening coefficient >- 0.130) after it has undergone collision deformation, and since the deformation resistance (load) stays at a low level, and the dynamic energy absorption is low preventing improvement in the impact absorption properties. On the other hand, with a martensite volume fraction of greater than 50%, the yield ratio becomes larger while work hardening coefficient is reduced, and deterioration also occurs in the tensile strength x total elongation and the hole expansion ratio. From the standpoint of press formability, the volume fraction of the martensite is preferred to be no greater than 30%.
Also, the ferrite is present at a volume fraction of preferably at; least 50%, and more preferably at least 70%, and its average grain size (mean circle equivalent diameter) is preferably no greater than 10 Vim, and more preferably no greater than 5 Vim, with the martensite preferably adjacent to the ferrite. This aids the fine dispersion oj= the martensite in the ferrite matrix, while effectively extending the property-improving effect, beyond simply a local effect, to the entire steel sheet, favorably acting to prevent the adverse effects of the martensite. The structure of the remainder present with the martensite and ferrite may be a mixed structure comprising a combination of one or. more from among pearlite, ba:inite, retained 'y, etc., and although primarily ba:inite is preferred in cases which require hole expansion properties, since retained y undergoes work-induced transformation into martensite by press forming, experimental results have shown that including retained aust:enite prior to press forming has an-effect even in preferred small amounts (5% or less).
Also, from the standpoint of impact absorption properties and press formability it is preferred for the ratio of the martensite and ferrite particle sizes to be no greater than 0.6, and the ratia of the hardnesses to be at least .1.5.
The rest:rictions on the values for the chemical components of-_ dual-phase type high-strength steel sheets with excellent dynamic deformation properties according to the inveni:ion, and the reasons for those restrictions, will now be Explained.
Dual-phase type high-strength steel sheets with excellent dynamic deformation properties which are used according to the invention are steel sheets containing the following chemical compositions, in terms of weight percentage: c~ at 0.02~0.25%, either or both Mn and Cr at a total of 0.15~3.5%, one or more from among Si, Al and P
at a total o:E 0.02~4.0%, if necessary also one or more from among N:i, Cu and Mo at a total of no more than 3.5%, one or more :from among Nb, Ti and V at no more than 0.30%, and either or both Ca and REM at 0.0005~0.01% for Ca and 0.0050.05% for REM, with the remainder Fe as the primary component. They are also dual-phase type high strength ste~=1 sheets with excellent dynamic deformation properties which contain, if necessary, one or more from among B (<_0.01), S (__<0.01%) and N (<_0.02%). These chemical components and their contents (percent by weight) will now be discussed.
C: C is the element which most strongly affects the microstructure of the steel sheet, and if its content is too low it will become difficult to obtain martensite with the desired amount and strength. Addition in too great an amount leads to unwanted carbide precipitation, inhibited increase in deformation resistance at higher strain rates and overly high strength, as well as poor
- 19 -press formab_Llity and weldability; the content is therefore 0.020.25 wt~.
Mn, Cr: Mn and Cr have an effect of stabilizing austenite and guaranteeing sufficient martensite, and are also solid solution hardening elements; they must therefore be added in a minimum amount of 0.15 wt~, but if added in i~oo much the aforementioned effect becomes saturated thus producing adverse effects such as preventing ferrite transformation, and thus they are added in the maximum amount of 3.5 wt~.
Si, A1, P: Si and A1 are useful elements for producing martensite, and they promote production of ferrite and suppress precipitation of carbides, thus having the e:Efect of guaranteeing sufficient martensite, as well as a solid solution hardening effect and a deoxidization effect. P can also promote martensite formation and solid solution hardening, similar to A1 and Si. From this standpoint, the minimum amount of Si + A1 + p added muat be at least 0.02 wt~. On the other hand, excessive addition will saturate this effect and result instead in b:rittleness, and therefore the maximum amount of addition is no more than 4.0 wt~. In particular, when an excellent surface condition is required, Si scales can be avoided b:y adding Si at no greater than 0.1 wt%, and conversely b:y adding it at 1.0 wto or greater Si scales can be produ~~ed over the entire surface so that they are not conspicuous. Also, when excellent secondary workability, toughness, spot weldability and recycling properties are required, the P content may be kept at no greater than 0.05%, and preferably no greater than 0.02%.
Ni, Cu, Mo: These elements are added when necessary, and are austenite-stabilizing elements similar to Mn, which increase the hardenability of the steel, and are effective for adjustment of the strength. From the standpoint of weldability and chemical treatment, they can be used when the amounts of C, Si, A1 and Mn are restricted, but if the total amount of these elements
Mn, Cr: Mn and Cr have an effect of stabilizing austenite and guaranteeing sufficient martensite, and are also solid solution hardening elements; they must therefore be added in a minimum amount of 0.15 wt~, but if added in i~oo much the aforementioned effect becomes saturated thus producing adverse effects such as preventing ferrite transformation, and thus they are added in the maximum amount of 3.5 wt~.
Si, A1, P: Si and A1 are useful elements for producing martensite, and they promote production of ferrite and suppress precipitation of carbides, thus having the e:Efect of guaranteeing sufficient martensite, as well as a solid solution hardening effect and a deoxidization effect. P can also promote martensite formation and solid solution hardening, similar to A1 and Si. From this standpoint, the minimum amount of Si + A1 + p added muat be at least 0.02 wt~. On the other hand, excessive addition will saturate this effect and result instead in b:rittleness, and therefore the maximum amount of addition is no more than 4.0 wt~. In particular, when an excellent surface condition is required, Si scales can be avoided b:y adding Si at no greater than 0.1 wt%, and conversely b:y adding it at 1.0 wto or greater Si scales can be produ~~ed over the entire surface so that they are not conspicuous. Also, when excellent secondary workability, toughness, spot weldability and recycling properties are required, the P content may be kept at no greater than 0.05%, and preferably no greater than 0.02%.
Ni, Cu, Mo: These elements are added when necessary, and are austenite-stabilizing elements similar to Mn, which increase the hardenability of the steel, and are effective for adjustment of the strength. From the standpoint of weldability and chemical treatment, they can be used when the amounts of C, Si, A1 and Mn are restricted, but if the total amount of these elements
- 20 -added exceeds 3.5 wt% the dominant ferrite phase will tend to be hardened, thus inhibiting the increase in deformation .resistance by a greater strain rate, as well as raising the cost of the steel sheet; the amount of these elements added is therefore 3.50 wt% or lower.
Nb, Ti, V: These elements are added when necessary, and are effe~~tive for strengthening the steel sheet through formation of carbides, nitrides and carbonitridea. However, when added at greater than 0.3 wt% they are deposited in large amounts in the dominant ferrite phase=_ or at the grain boundaries as carbides, nitrides and carbonitrides, becoming a source of the mobile dislocation during high speed deformation, and inhibiting t:he increase in deformation resistance by greater strain rates. In addition, the deformation resistance of the dominant phase becomes higher than necessary, thus wasting the C and leading to higher costs; the maximum amount to be added is therefore 0.3 wt%.
B: B is an element which is effective for strengthening since it improves the hardenability of the steel by suppressing production of ferrite, but if it is added at greater than 0.01 wt% its effect will be saturated, and therefore B is added at a maximum of 0.01 wt%.
Ca, REM: Ca is added to at least 0.0005 wt% for improved press formability (especially hole expansion ratio) by shape control (spheroidizat~ion) of sulfide-based inclusions, and the maximum amount thereof to be added is 0.01 wt% in consideration of effect saturation and the adverse effect due to increase in the aforemention?d inclusions (reduced hole expansion ratio).
For the same reasons, REM is added in an amount of from 0.005% to 0.05 wt%.
S: The amount of S is no greater than 0.01 wt%, and preferably no greater than 0.003 wt%, from the standpoint of press formability (especially hole expansion ratio) by
Nb, Ti, V: These elements are added when necessary, and are effe~~tive for strengthening the steel sheet through formation of carbides, nitrides and carbonitridea. However, when added at greater than 0.3 wt% they are deposited in large amounts in the dominant ferrite phase=_ or at the grain boundaries as carbides, nitrides and carbonitrides, becoming a source of the mobile dislocation during high speed deformation, and inhibiting t:he increase in deformation resistance by greater strain rates. In addition, the deformation resistance of the dominant phase becomes higher than necessary, thus wasting the C and leading to higher costs; the maximum amount to be added is therefore 0.3 wt%.
B: B is an element which is effective for strengthening since it improves the hardenability of the steel by suppressing production of ferrite, but if it is added at greater than 0.01 wt% its effect will be saturated, and therefore B is added at a maximum of 0.01 wt%.
Ca, REM: Ca is added to at least 0.0005 wt% for improved press formability (especially hole expansion ratio) by shape control (spheroidizat~ion) of sulfide-based inclusions, and the maximum amount thereof to be added is 0.01 wt% in consideration of effect saturation and the adverse effect due to increase in the aforemention?d inclusions (reduced hole expansion ratio).
For the same reasons, REM is added in an amount of from 0.005% to 0.05 wt%.
S: The amount of S is no greater than 0.01 wt%, and preferably no greater than 0.003 wt%, from the standpoint of press formability (especially hole expansion ratio) by
- 21 -sulfide-based inclusions, and reduced spot weldability.
The method of applying the pre-deformation according to the inveni~ion will now be explained. The pre-deformation may be press forming for member shaping, or it may be working with a tempering rolling or tension leveler which applied to the steel sheet material prior to its press forming. In this case, either or both a tempering roller and tension leveler may be used. That is, the mean, used may include a tempering rolling, a tension leve:Ler, or a tempering roller and tension leveler. The steel sheet material may also be subjected to press forming after being worked with a tempering rolling or tension leveler. The amount of pre-deformation applied with the tempering rolling and/or tension leve:Ler, i.e. the degree of plastic deformation (T), will differ depending on the initial dislocation density, and T should be small if the initial density is large. Also, with few solid solution elements the introduced dislocations cannot be fixed, and high dynamic deformation :properties cannot be guaranteed.
Consequently, it was found that the plastic deformation (T) is determined based on the ratio between the yield strength YS(0) and the tensile strength TS'(5) in the static tensile test with pre-deformation at 50 of equivalent strain or after further bake hardening treatment (BH treatment), or YS(0)/TS'(5). That is, YS(0)/TS'(5) is an indicator of the sum of the initial dislocation density and the dislocation density introduced by 5% deformation, and the amount of the solid solution elements; it may be concluded that a smaller YS(0)/TS'(5) means a higher initial dislocation density and more of the solid solution elements. YS(0)/TS'(5) is therefore no greater than 0.7, and is preferably provided according to the following equation:
2.5 ~YS(0)/TS'(5) - 0.5} + 15 ? T >_ 2.5 ~YS(0)/TS'(5) - 0.5} + 0.5 wherein the upper limit for T is determined from the
The method of applying the pre-deformation according to the inveni~ion will now be explained. The pre-deformation may be press forming for member shaping, or it may be working with a tempering rolling or tension leveler which applied to the steel sheet material prior to its press forming. In this case, either or both a tempering roller and tension leveler may be used. That is, the mean, used may include a tempering rolling, a tension leve:Ler, or a tempering roller and tension leveler. The steel sheet material may also be subjected to press forming after being worked with a tempering rolling or tension leveler. The amount of pre-deformation applied with the tempering rolling and/or tension leve:Ler, i.e. the degree of plastic deformation (T), will differ depending on the initial dislocation density, and T should be small if the initial density is large. Also, with few solid solution elements the introduced dislocations cannot be fixed, and high dynamic deformation :properties cannot be guaranteed.
Consequently, it was found that the plastic deformation (T) is determined based on the ratio between the yield strength YS(0) and the tensile strength TS'(5) in the static tensile test with pre-deformation at 50 of equivalent strain or after further bake hardening treatment (BH treatment), or YS(0)/TS'(5). That is, YS(0)/TS'(5) is an indicator of the sum of the initial dislocation density and the dislocation density introduced by 5% deformation, and the amount of the solid solution elements; it may be concluded that a smaller YS(0)/TS'(5) means a higher initial dislocation density and more of the solid solution elements. YS(0)/TS'(5) is therefore no greater than 0.7, and is preferably provided according to the following equation:
2.5 ~YS(0)/TS'(5) - 0.5} + 15 ? T >_ 2.5 ~YS(0)/TS'(5) - 0.5} + 0.5 wherein the upper limit for T is determined from the
- 22 -standpoint o:E press formability including impact absorption property and flexibility.
A method of producing a dual-phase type high strength hot rolled steel sheet and a cold rolled steel sheet with e:~ccellent dynamic deformation properties according to the invention will now be explained. In this product:ion method, a continuous cast slab is fed directly from casting to a hot rolling step, or is hot rolled upon ceheating after momentary, cooling. Thin gauge continuous casting and continuous hot rolling techniques (endless hot rolling) may be applied for the hot rolling :in addition to normal continuous casting, but in order to avoid a lower ferrite volume fraction and a coarser average grain size of the thin steel sheet microstructure, the bar (cast strip) thickness at the hot rolling approach side (the initial steel bar thickness) is preferred to be at least 25 mm. At less than 25 mm, the mean circle equivalent size of ferrite of the steel sheet is made=_ coarser, while it is also a disadvantage against obtaining the desired martensite. The final pass rolling speed for the hot rolling is preferred to be at least 500 mpm and more preferably at least 600 mpm, in light of the problems described above. At less than 500 mpm, the mean circle equivalent diameter of ferrite of the steel sheet is made coarser, while it is also a disadvantage against obtaining the desired martensite.
The finishing temperature for the hot rolling is from Ar3 - 50°C to Ar3 + 120°C. At lower than Ar3 - 50°C, deformed ferrite is produced, with inferior work hardening property and press formability. At higher than Ar3 + 120°C, and the mean circle equivalent size of ferrite of the steel sheet is made coarser, while it is also becomes difficult to obtain the desired martensite.
The average cooling rate for cooling in the run-out table is at least 5°C/sec. At less than 5°C/sec it becomes difficult to obtain the desired martensite.
A method of producing a dual-phase type high strength hot rolled steel sheet and a cold rolled steel sheet with e:~ccellent dynamic deformation properties according to the invention will now be explained. In this product:ion method, a continuous cast slab is fed directly from casting to a hot rolling step, or is hot rolled upon ceheating after momentary, cooling. Thin gauge continuous casting and continuous hot rolling techniques (endless hot rolling) may be applied for the hot rolling :in addition to normal continuous casting, but in order to avoid a lower ferrite volume fraction and a coarser average grain size of the thin steel sheet microstructure, the bar (cast strip) thickness at the hot rolling approach side (the initial steel bar thickness) is preferred to be at least 25 mm. At less than 25 mm, the mean circle equivalent size of ferrite of the steel sheet is made=_ coarser, while it is also a disadvantage against obtaining the desired martensite. The final pass rolling speed for the hot rolling is preferred to be at least 500 mpm and more preferably at least 600 mpm, in light of the problems described above. At less than 500 mpm, the mean circle equivalent diameter of ferrite of the steel sheet is made coarser, while it is also a disadvantage against obtaining the desired martensite.
The finishing temperature for the hot rolling is from Ar3 - 50°C to Ar3 + 120°C. At lower than Ar3 - 50°C, deformed ferrite is produced, with inferior work hardening property and press formability. At higher than Ar3 + 120°C, and the mean circle equivalent size of ferrite of the steel sheet is made coarser, while it is also becomes difficult to obtain the desired martensite.
The average cooling rate for cooling in the run-out table is at least 5°C/sec. At less than 5°C/sec it becomes difficult to obtain the desired martensite.
- 23 -The coiling temperature is no higher than 350°C. At higher than :350°C it becomes difficult to obtain the desired martensite.
According to the invention, it was found particularly that a correlation exists between the finishing temperature in the hot rolling step, the finishing approach temperature and the coiling temperature. That is, as shown in Fig. 10 and Fig. 11, specific conditions exist which are determined primarily between the :Finishing temperature, finishing approach temperature and the coiling temperature. Specifically, the hot rolling is carried out so that when the finishing temperature :For hot rolling is in the range of Ar3 - 50°C
to Ar3 + 120°C, the metallurgy parameter A satisfies , inequalities (1) and (2). The above-mentioned metallurgy parameter A may be expressed by the following equation.
A = Ef~ x: exp{ ( 75282 - 42745 x Cep) / [ 1 . 978 x (FT +
273)]}
where FT: finishing temperature (°C) Ceq: carbon equivalents = C + Mneq/6 ( % ) MnE,q: manganese equivalents = Mn + (Ni + Cr +
Cu + Mo)/2 (%) Ef: final pass strain rate (s 1) (v/JRxhl) x (1/Jr) x In X1/(1-r)}
hl: final pass approach sheet thickness hz: final pass exit sheet thickness r . (hl - hz) /hl R . roll radius v . final pass exit speed 0T: fin.ishing temperature (finishing final pass exit temperature) - finishing approach temperature (finish.ing first pass approach temperature) Ar3: 901. - 325 C% + 33 Si % - 92 Mneq Thereafter, it is preferred for the average cooling rate on the :run-out table to be at least 5°C/sec, and the coiling to b~a carried out under conditions such that the
According to the invention, it was found particularly that a correlation exists between the finishing temperature in the hot rolling step, the finishing approach temperature and the coiling temperature. That is, as shown in Fig. 10 and Fig. 11, specific conditions exist which are determined primarily between the :Finishing temperature, finishing approach temperature and the coiling temperature. Specifically, the hot rolling is carried out so that when the finishing temperature :For hot rolling is in the range of Ar3 - 50°C
to Ar3 + 120°C, the metallurgy parameter A satisfies , inequalities (1) and (2). The above-mentioned metallurgy parameter A may be expressed by the following equation.
A = Ef~ x: exp{ ( 75282 - 42745 x Cep) / [ 1 . 978 x (FT +
273)]}
where FT: finishing temperature (°C) Ceq: carbon equivalents = C + Mneq/6 ( % ) MnE,q: manganese equivalents = Mn + (Ni + Cr +
Cu + Mo)/2 (%) Ef: final pass strain rate (s 1) (v/JRxhl) x (1/Jr) x In X1/(1-r)}
hl: final pass approach sheet thickness hz: final pass exit sheet thickness r . (hl - hz) /hl R . roll radius v . final pass exit speed 0T: fin.ishing temperature (finishing final pass exit temperature) - finishing approach temperature (finish.ing first pass approach temperature) Ar3: 901. - 325 C% + 33 Si % - 92 Mneq Thereafter, it is preferred for the average cooling rate on the :run-out table to be at least 5°C/sec, and the coiling to b~a carried out under conditions such that the
- 24 -relationship between the metallurgy parameter A and the coiling temperature (CT) satisfies inequality (3).
9 <_ logA <- 18 ~ (1) eT <- 21 x logA - 61 (2) CT <_ 6 ;K logA + 242 (3) In inequality (1) above, a log A of less than 9 is unacceptable from the viewpoint of production of retained martensite and refinement of the microstructure, while it will also reault in an inferior dynamic deformation resistance adyn and 5~10~ work hardening property. Also, if log A is -to be greater than 18, massive equipment will be required to achieve it. With inequality (2), if the condition of inequality (2) is not satisfied it will be impossible to obtain the desired martensite, and the dynamic deformation resistance ~dyn and 5~10~ work hardening property, etc. will be inferior. The lower limit for eT is more flexible with a lower log A as indicated by inequality (2). Furthermore, if the relationship with the coiling temperature in inequality (3) is not satisfied, there will be an adverse effect on ensuring the amount of martensite, while the retained y will be excessively stable even if retained y can be obtained, it will be impossible to obtain the desired martensite during deformation, and the dynamic deformation resistance 6dyn and 5~10o work hardening property, etc. will be inferior. The limit for the coiling temperature is more flexible with a higher log A.
The cold rolled sheet according to the invention is then subjected to the different steps following hot-rolling and coiling and is cold rolled and subjected to annealing. The annealing is ideally continuous annealing through an annealing cycle such as shown in Fig. 12, and during the annealing of the continuous annealing step, it must be kept. for at least 10 seconds in the temperature range of Acl - Ac3. At less than Acl austenite will not be produced and it will therefore be impossible to obtain
9 <_ logA <- 18 ~ (1) eT <- 21 x logA - 61 (2) CT <_ 6 ;K logA + 242 (3) In inequality (1) above, a log A of less than 9 is unacceptable from the viewpoint of production of retained martensite and refinement of the microstructure, while it will also reault in an inferior dynamic deformation resistance adyn and 5~10~ work hardening property. Also, if log A is -to be greater than 18, massive equipment will be required to achieve it. With inequality (2), if the condition of inequality (2) is not satisfied it will be impossible to obtain the desired martensite, and the dynamic deformation resistance ~dyn and 5~10~ work hardening property, etc. will be inferior. The lower limit for eT is more flexible with a lower log A as indicated by inequality (2). Furthermore, if the relationship with the coiling temperature in inequality (3) is not satisfied, there will be an adverse effect on ensuring the amount of martensite, while the retained y will be excessively stable even if retained y can be obtained, it will be impossible to obtain the desired martensite during deformation, and the dynamic deformation resistance 6dyn and 5~10o work hardening property, etc. will be inferior. The limit for the coiling temperature is more flexible with a higher log A.
The cold rolled sheet according to the invention is then subjected to the different steps following hot-rolling and coiling and is cold rolled and subjected to annealing. The annealing is ideally continuous annealing through an annealing cycle such as shown in Fig. 12, and during the annealing of the continuous annealing step, it must be kept. for at least 10 seconds in the temperature range of Acl - Ac3. At less than Acl austenite will not be produced and it will therefore be impossible to obtain
- 25 -martensite thereafter, while at greater than Ac3 the austenite mon.ophase structure will be coarse, and it will therefore be impossible to obtain the desired average grain size for the martensite. Also, at less than 10 seconds the a.ustenite production will be insufficient, making it impossible to obtain the desired martensite thereafter. The maximum residence time is preferably no greater than 200 seconds, from the standpoint of avoiding addition to the equipment and coarsening of the microstructure. The cooling after this annealing must be at an average cooling rate of at least 5°C/sec. At less than 5°C/sec the desired space factor for the martensite cannot be achieved. Although there is no particular upper limit here, it is preferably 300°C/sec when considering temperature control during the cooling.
According to the invention, the cooled steel sheet is heated to a temperature To from Acl- Ac3 in the continuous annealing cycle shown in Fig. 12, and cooled under cooling conditions provided by a method wherein cooling to a secondary cooling start temperature Tq in the range of 550°C-To at the primary cooling rate of 110°C/sec i=~ followed by cooling to a secondary cooling end temperature Te which is no higher than a temperature Tem which is determined by the chemical compositions of the steel anti annealing temperature To, at a secondary cooling rate of 10200°C/sec. This is a method whereby the cooling end temperature Te in the continuous annealing cycle shown in Fig. 12 is represented as a function of t:he chemical compositions and annealing temperature, and is kept under a given critical value.
After cooling to Te, the temperature is preferably held in a range of Te - 50°C to 400°C for up to 20 minutes prior to cooling to room temperature.
Here, Tem is the martensite transformation start temperature nor the retained austenite at the quenching start point ~Cq. That is, Tem is defined by Tem = T1 -
According to the invention, the cooled steel sheet is heated to a temperature To from Acl- Ac3 in the continuous annealing cycle shown in Fig. 12, and cooled under cooling conditions provided by a method wherein cooling to a secondary cooling start temperature Tq in the range of 550°C-To at the primary cooling rate of 110°C/sec i=~ followed by cooling to a secondary cooling end temperature Te which is no higher than a temperature Tem which is determined by the chemical compositions of the steel anti annealing temperature To, at a secondary cooling rate of 10200°C/sec. This is a method whereby the cooling end temperature Te in the continuous annealing cycle shown in Fig. 12 is represented as a function of t:he chemical compositions and annealing temperature, and is kept under a given critical value.
After cooling to Te, the temperature is preferably held in a range of Te - 50°C to 400°C for up to 20 minutes prior to cooling to room temperature.
Here, Tem is the martensite transformation start temperature nor the retained austenite at the quenching start point ~Cq. That is, Tem is defined by Tem = T1 -
- 26 -T2, or the difference between the value excluding the effect of the C concentration in the austenite (T1) and the value indicating the effect of the C concentration (T2). Here, T1 is the temperature calculated from the solid solution element concentration excluding C, and T2 is the temperature calculated from the C concentration in the retained austenite at Acl and Ac3 determined by the chemical compositions of the steel and Tq determined by ., the annealing temperature To. Ceq represents the carbon equivalents in the retained austenite at the annealing temperature To. Thus, T1 is expressed as:
T1 = 561 - 33 x {Mn% + (Ni + Cr + Cu + Mo)/2}
and T2 is expressed in terms of:
Acl = 72:3 - 0.7 x Mn% - 16.9 x Ni% + 29.1 x Si% +
16.9 x Cr%, Ac3 = 911) - 203 x (C% ) l~z - 15 . 2 x Ni% + 44 . 7 x Si% +
104 x V% + 31.5 x Mo% - 30 x Mn% - 11 x Cr% - 20 x Cu% + 70 x P% + 40 x A1% + 400 x Ti%, and the annealing temperature To, and when Ceq'~ - (Ac3 - Acl) x C/(To - Acl) + (Mn + Si/4 + Ni/7 + Cr + Cu + 1.5 Mo)/6 is greater than 0.6, T2 - 474 x (Ac3 - Acl) x C/(To -Acl) , and when it is 0.6 or less, T2 - 474 x (Ac3 - Acl) x C/~3 x (Ac3 - Ac1) x C + [(Mn + Si/4 + Ni/7 + Cr + Cu + 1.5 Mo)/2 - 0.85)] x (To - Aci).
In other words, when Te is equal to or greater than Tem, the desired martensite cannot be obtained. Also, if Toa is 400°C or higher, the martensite obtained by cooling is tempered, making it impossible to achieve satisfactory dynamic properties and press formability.
On the other hand, if Toa is less than Te - 50°C, additional ccoling equipment is necessary, and greater variation will result in the material due to the difference between the temperature of the continuous annealing furnace and the temperature of the steel sheet;
T1 = 561 - 33 x {Mn% + (Ni + Cr + Cu + Mo)/2}
and T2 is expressed in terms of:
Acl = 72:3 - 0.7 x Mn% - 16.9 x Ni% + 29.1 x Si% +
16.9 x Cr%, Ac3 = 911) - 203 x (C% ) l~z - 15 . 2 x Ni% + 44 . 7 x Si% +
104 x V% + 31.5 x Mo% - 30 x Mn% - 11 x Cr% - 20 x Cu% + 70 x P% + 40 x A1% + 400 x Ti%, and the annealing temperature To, and when Ceq'~ - (Ac3 - Acl) x C/(To - Acl) + (Mn + Si/4 + Ni/7 + Cr + Cu + 1.5 Mo)/6 is greater than 0.6, T2 - 474 x (Ac3 - Acl) x C/(To -Acl) , and when it is 0.6 or less, T2 - 474 x (Ac3 - Acl) x C/~3 x (Ac3 - Ac1) x C + [(Mn + Si/4 + Ni/7 + Cr + Cu + 1.5 Mo)/2 - 0.85)] x (To - Aci).
In other words, when Te is equal to or greater than Tem, the desired martensite cannot be obtained. Also, if Toa is 400°C or higher, the martensite obtained by cooling is tempered, making it impossible to achieve satisfactory dynamic properties and press formability.
On the other hand, if Toa is less than Te - 50°C, additional ccoling equipment is necessary, and greater variation will result in the material due to the difference between the temperature of the continuous annealing furnace and the temperature of the steel sheet;
- 27 -this temperature was therefore determined as the lower limit. Also, the upper limit for the holding time was determined to be 20 minutes, because'when it is longer than 20 minutes it becomes necessary to expand the equipment.
By employing the chemical composition and production method described above, it is possible to produce a dual-phase type high-strength steel sheet with excellent dynamic deformation properties, wherein the microstructure of the steel sheet is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction from 3%~50% after shaping and working at 5% equivalent strain, and wherein t:he difference between the quasi-static deformation ~~trength cs when deformed in a strain rate range of 5 x 104 - 5 x 103 (1/s) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength ~d measured ~n a strain rate range of 5 x 102 - 5 x 10' (1/s) after the aforementioned pre-deformation, i.e.
(6d - as), is at least 60 MPa, and the work hardening coefficient at 510% strain is at least 0.13. The steel sheets according to the invention may be made into any desired product by annealing, tempering rolling, electronic coating or hot-dip coating.
Examples The present invention will now be explained by way of examples.
(Example 1) The 26 steel materials listed in Table 1 (steel nos.
126) were heated to 10501250°C and subjected to hot rolling, cooling and coiling under the production conditions listed in Table 2, to produce hot rolled steel sheets. As shown in Table 3, the steel sheets satisfying the chemical composition conditions and production
By employing the chemical composition and production method described above, it is possible to produce a dual-phase type high-strength steel sheet with excellent dynamic deformation properties, wherein the microstructure of the steel sheet is a composite microstructure wherein the dominating phase is ferrite, and the second phase is another low temperature product phase containing martensite at a volume fraction from 3%~50% after shaping and working at 5% equivalent strain, and wherein t:he difference between the quasi-static deformation ~~trength cs when deformed in a strain rate range of 5 x 104 - 5 x 103 (1/s) after pre-deformation of more than 0% and less than or equal to 10% of equivalent strain, and the dynamic deformation strength ~d measured ~n a strain rate range of 5 x 102 - 5 x 10' (1/s) after the aforementioned pre-deformation, i.e.
(6d - as), is at least 60 MPa, and the work hardening coefficient at 510% strain is at least 0.13. The steel sheets according to the invention may be made into any desired product by annealing, tempering rolling, electronic coating or hot-dip coating.
Examples The present invention will now be explained by way of examples.
(Example 1) The 26 steel materials listed in Table 1 (steel nos.
126) were heated to 10501250°C and subjected to hot rolling, cooling and coiling under the production conditions listed in Table 2, to produce hot rolled steel sheets. As shown in Table 3, the steel sheets satisfying the chemical composition conditions and production
- 28 -conditions according to the invention have a dual-phase structure with a martensite volume fraction of at least 3~ and no greater than 50~, and as shown in Fig. 4, the mechanical properties of the hot rolled steel sheets indicated excellent impact absorption properties as represented by a work hardening coefficient of at least 0.13 at 5~10~: strain, ~d - os >_ 60 MPa, and odyn ? 0.766 x TS + 250, while also having suitable press formability and weldabili.ty.
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- 30 -Table 1 (cont.) Chemical compositions of steels Ste~slTransformation Type No. temperature C
Acl Ac3 Ar3 1. 741 863 793 present invention 2 741 863 793 present invention 3 744 880 805 present invention 9: 756 871 809 present invention 5 709 863 756 present invention E 706 816 731 present invention 7 726 851 794 present invention 8 733 874 791 present invention 712 834 774 present invention 1C1 722 830 787 present invention 17. 733 839 736 present invention 12 741 863 793 present invention la 739 857 775 comparative example 14E 741 863 793 comparative example 1-'i 713 861 806 comparative example 1f> 728 839 732 present: invention 17 740 887 802 present invention 18 767 889 763 present invention 19 735 870 807 present invention 2t) 736 862 798 present invention 2:1 753 860 751 present invention 2:? 704 810 713 present invention 2:3 720 837 801 present invention 24 717 82.6773 present invention 25 752 923 771 present invention 26 722 779 762 comparative example
Acl Ac3 Ar3 1. 741 863 793 present invention 2 741 863 793 present invention 3 744 880 805 present invention 9: 756 871 809 present invention 5 709 863 756 present invention E 706 816 731 present invention 7 726 851 794 present invention 8 733 874 791 present invention 712 834 774 present invention 1C1 722 830 787 present invention 17. 733 839 736 present invention 12 741 863 793 present invention la 739 857 775 comparative example 14E 741 863 793 comparative example 1-'i 713 861 806 comparative example 1f> 728 839 732 present: invention 17 740 887 802 present invention 18 767 889 763 present invention 19 735 870 807 present invention 2t) 736 862 798 present invention 2:1 753 860 751 present invention 2:? 704 810 713 present invention 2:3 720 837 801 present invention 24 717 82.6773 present invention 25 752 923 771 present invention 26 722 779 762 comparative example
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- 35 -(Example 2) The 22 steel materials listed in Table 5 (steel nos.
2748) were heated to 10501250°C and subjected to hot rolling, coo:Ling and coiling, followed by acid pickling and then cold rolling under the conditions listed in Table 6 to p=roduce cold rolled steel sheets.
Temperatures Acl and Ac3 were then calculated from the chemical compositions for each steel, and the sheets were subjected to heating, cooling and holding under the annealing conditions listed in Table 6, prior to cooling to room temperature. As shown in Table 7, the steel sheets satisfying the chemical composition conditions and production conditions according to the invention have a dual-phase structure with a martensite volume fraction of at least 3% and no greater than 50% and, as shown in Fig.
8, the mechanical properties of the hot-rolled steel sheets indicated excellent impact absorption properties as represented by a work hardening coefficient of at least 0.13 at 510% strain, od - 6s :? 60 MPa, and adyn 0.766 x TS + 250, while also having suitable press formability and weldability.
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-~ -x The microstructure was evaluated by the following method.
Identification of the ferrite, bainite, martensite and residual structure, observation of the location and measurement of the average grain size (mean circle equivalent.di.ameter) was accomplished using a 1000 magnification optical micrograph with the thin steel sheet rolling direction cross-section etched with a nital and the reagent disclosed in Japanese Unexamined Patent Publication rlo. 59-219473.
The proF>erties were evaluated by the following methods.
A tensile test was conducted according to JISS
(gauge mark distance: 50 mm, parallel part width: 25 mm) with a strain rate of 0.001/s and, upon determining the tensile strength (TS), yield strength (YS), total elongation (T. E1) and work hardening coefficient (n value for 1%~~5% strain), the YS x work hardening coefficient and TS x T. E1. were calculated.
The stretch flanging property was measured by expanding a 20 mm punched hole from the burrless side with a 30° cone punch, and determining the hole expansion ratio (d/do) between the hole diameter (d) at the moment at which the crack penetrated the plate thickness and the original hollow diameter (do, 20 mm).
The spot. weldability was judged to be unsuitable if a spot welding test piece bonded at a current of 0.9 times the expulsion current using an electrode with a tip radius of 5 times the square root of the steel sheet thickness underwent peel fracture when ruptured with a chisel.
Industrial At?plicability As explained above, the present invention makes it possible to provide, in an economical and stable manner, high-strengt'.z hot rolled steel sheets and cold rolled steel sheets for automobiles which provide previously unobtainable excellent impact absorption properties and press formability and thus offers a markedly wider range of objects and conditions for uses of high-strength steel sheets.
2748) were heated to 10501250°C and subjected to hot rolling, coo:Ling and coiling, followed by acid pickling and then cold rolling under the conditions listed in Table 6 to p=roduce cold rolled steel sheets.
Temperatures Acl and Ac3 were then calculated from the chemical compositions for each steel, and the sheets were subjected to heating, cooling and holding under the annealing conditions listed in Table 6, prior to cooling to room temperature. As shown in Table 7, the steel sheets satisfying the chemical composition conditions and production conditions according to the invention have a dual-phase structure with a martensite volume fraction of at least 3% and no greater than 50% and, as shown in Fig.
8, the mechanical properties of the hot-rolled steel sheets indicated excellent impact absorption properties as represented by a work hardening coefficient of at least 0.13 at 510% strain, od - 6s :? 60 MPa, and adyn 0.766 x TS + 250, while also having suitable press formability and weldability.
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~ 3 p,o +~ f~O 00N M a0N u1o~~ h o~~
-.-icd O O O M O h h -iO ~ ctf <n ~ ~"~
V~ C1 G~M N O~~'~1O v0~1riO h e~N N M v0M Qv . ~-' ~
b M voiow h h ~ ~ m o rwo h ~ h h h h ~ h co~ cs ~ ~
.
a~ ~
C
~i U ~ ~ ---.~ n1 r1 ca o a~
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dy ~-i N O
c0 a) O h I~h Y N O~~!1O WOd Ovr-iO O vDM .-IOvN ~ +
+
~c t0 ~ ~ -i~ M o0a0v0~ u'7N d N M S-r ~ U
u7 r-Iv0N '-IOvO W N t c0r N N N O M N M U (!1 O. i U r~
G
r-IN r-IriN N N N N N M N N r1N -ri H
~1 ~
U G SC
cn O
-~-I -.~ I
cn s-~
t W J t0 G .~
c0 cd >r O
a~ H ~ G
O h ~
O b x N N O N M .-I.-iN I~N Ovr1O~.-iN I~O Ov.-i1~1QJ ~
tb S-i O
U 3 o.m o h ~ ,-I~ avo ~ ~ncoo.0 0 0 0or~~ ovco p-, .-I.-1~ N N H r-IN N N r1r-1N N N r1r-1 N r-IN
'b O
C1.
v p~ N
N I
1-.i V7 >a -.i ~
U H ~ U
3 b C
r-IO U o ~ H N M t tfW I~COO~O r1N M ~tu-1t01~CO b O N
h COO O v M M Y ~ ~ .J~Ysf~Y~Y~ T3....0 ~
O N N N M M M M M M M M ~ W-i h N M
d ~ .-I
~x ~
-~ -x The microstructure was evaluated by the following method.
Identification of the ferrite, bainite, martensite and residual structure, observation of the location and measurement of the average grain size (mean circle equivalent.di.ameter) was accomplished using a 1000 magnification optical micrograph with the thin steel sheet rolling direction cross-section etched with a nital and the reagent disclosed in Japanese Unexamined Patent Publication rlo. 59-219473.
The proF>erties were evaluated by the following methods.
A tensile test was conducted according to JISS
(gauge mark distance: 50 mm, parallel part width: 25 mm) with a strain rate of 0.001/s and, upon determining the tensile strength (TS), yield strength (YS), total elongation (T. E1) and work hardening coefficient (n value for 1%~~5% strain), the YS x work hardening coefficient and TS x T. E1. were calculated.
The stretch flanging property was measured by expanding a 20 mm punched hole from the burrless side with a 30° cone punch, and determining the hole expansion ratio (d/do) between the hole diameter (d) at the moment at which the crack penetrated the plate thickness and the original hollow diameter (do, 20 mm).
The spot. weldability was judged to be unsuitable if a spot welding test piece bonded at a current of 0.9 times the expulsion current using an electrode with a tip radius of 5 times the square root of the steel sheet thickness underwent peel fracture when ruptured with a chisel.
Industrial At?plicability As explained above, the present invention makes it possible to provide, in an economical and stable manner, high-strengt'.z hot rolled steel sheets and cold rolled steel sheets for automobiles which provide previously unobtainable excellent impact absorption properties and press formability and thus offers a markedly wider range of objects and conditions for uses of high-strength steel sheets.
Claims (10)
1. A dual-phase type high-strength steel sheets having high impact energy absorption and dynamic deformation properties, characterized in that:
the steel sheet contains C in a range between 0.02 and 0.25%, at least one of Mn and Cr at a total in a range of between 0.15 and 3.5%, at least one element from among Si, Al and P at a total in a range of between 0.02 and 4.0%;
and further optionally contains at least one from among Ni, Cu and Mo at a total of at most 3.5%, at least one from among Nb, Ti and V at a total of at most 0.3%, and at least one of Ca and REM in a range between 0.0005 and 0.01% for Ca and in a range between 0.005 and 0.05%
for REM, at least one from among F3 at less than 0.01%, S
at less than 0.01% and N at less than 0.02, with a remainder Fe and unavoidable impurities;
the steel sheet having a composite microstructure as a final microstructure, a dominating phase being ferrite and a second phase being a low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, wherein a difference between a quasi-static deformation strength as when deformed in a strain rate range between 5 × 10-4 and 5 × 10-3 (S-1) after pre-deformation of more than 0% and at most 10% of equivalent strain, and a dynamic deformation strength .sigma.d when deformed in a strain rate range between 5 × 10 2 and × 10 3 (S-1) after the pre-deformation, i.e. (.sigma.d - .sigma.s), is at least 60 MPa, and a work hardening coefficient between and 10% strain is at least 0.13.
the steel sheet contains C in a range between 0.02 and 0.25%, at least one of Mn and Cr at a total in a range of between 0.15 and 3.5%, at least one element from among Si, Al and P at a total in a range of between 0.02 and 4.0%;
and further optionally contains at least one from among Ni, Cu and Mo at a total of at most 3.5%, at least one from among Nb, Ti and V at a total of at most 0.3%, and at least one of Ca and REM in a range between 0.0005 and 0.01% for Ca and in a range between 0.005 and 0.05%
for REM, at least one from among F3 at less than 0.01%, S
at less than 0.01% and N at less than 0.02, with a remainder Fe and unavoidable impurities;
the steel sheet having a composite microstructure as a final microstructure, a dominating phase being ferrite and a second phase being a low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, wherein a difference between a quasi-static deformation strength as when deformed in a strain rate range between 5 × 10-4 and 5 × 10-3 (S-1) after pre-deformation of more than 0% and at most 10% of equivalent strain, and a dynamic deformation strength .sigma.d when deformed in a strain rate range between 5 × 10 2 and × 10 3 (S-1) after the pre-deformation, i.e. (.sigma.d - .sigma.s), is at least 60 MPa, and a work hardening coefficient between and 10% strain is at least 0.13.
2. A dual-phase type high-strength steel sheets having high impact energy absorption and dynamic deformation properties, characterized in that:
the steel sheet contains C between 0.02 and 0.25%, at least one of Mn and Cr at a total between 0.15 and 3.5%, at least one element from among S~, Al and P at a total between 0.02 and 4.0%, and further optionally contains at least one from among Ni, Cu and Mo at a total of at most 3.5%, at least one from among Nb, Ti and V at a total of at most 0.3%, and at least one of Ca and REM between 0.0005 and 0.01% for Ca and between 0.005 and 0.05% for REM, at least one from among B at less than 0.01%, S at less than 0.01% and N at less than 0.020, with a remainder Fe and unavoidable impurities;
the steel sheet having a composite microstructure as a final microstructure, a dominating phase being ferrite and a second phase being a low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, an average value .sigma.dyn (MPa) of a deformation stress being in a range between 3 and 10% of equivalent strain when deformed in a strain range between 5 × 10 2 and 5 × 10 3 (S-1), after the pre-deformation of more than 0%
and of at most 10% of equivalent strain, satisfies an inequality: .sigma.dyn >= 0.766 × TS + 250 as expressed in terms of a tensile strength TS (MPa) in a quasi-static tensile test as measured in a strain rate range between 5 × 10-4 and 5 × 10-3 (S-1) prior to pre-deformation, and a work hardening coefficient between 5 and 10% strain being at least 0.13.
the steel sheet contains C between 0.02 and 0.25%, at least one of Mn and Cr at a total between 0.15 and 3.5%, at least one element from among S~, Al and P at a total between 0.02 and 4.0%, and further optionally contains at least one from among Ni, Cu and Mo at a total of at most 3.5%, at least one from among Nb, Ti and V at a total of at most 0.3%, and at least one of Ca and REM between 0.0005 and 0.01% for Ca and between 0.005 and 0.05% for REM, at least one from among B at less than 0.01%, S at less than 0.01% and N at less than 0.020, with a remainder Fe and unavoidable impurities;
the steel sheet having a composite microstructure as a final microstructure, a dominating phase being ferrite and a second phase being a low temperature product phase containing martensite at a volume fraction between 3% and 50% after deformation at 5% equivalent strain of the steel sheet, an average value .sigma.dyn (MPa) of a deformation stress being in a range between 3 and 10% of equivalent strain when deformed in a strain range between 5 × 10 2 and 5 × 10 3 (S-1), after the pre-deformation of more than 0%
and of at most 10% of equivalent strain, satisfies an inequality: .sigma.dyn >= 0.766 × TS + 250 as expressed in terms of a tensile strength TS (MPa) in a quasi-static tensile test as measured in a strain rate range between 5 × 10-4 and 5 × 10-3 (S-1) prior to pre-deformation, and a work hardening coefficient between 5 and 10% strain being at least 0.13.
3. The dual-phase type high-strength steel sheet according to any one of claims 1 and 2, characterized in that a ratio between the yield strength YS(0) and a tensile strength TS'(5) in a static tensile test after one of: the pre-deformation at 5% of equivalent strain and:
bake hardening treatment (BH treatment), satisfies a first inequality: YS(0)/TS'(5) <= 0.7, and also satisfies a second inequality: yield strength YS(0) × work hardening coefficient >= 70.
bake hardening treatment (BH treatment), satisfies a first inequality: YS(0)/TS'(5) <= 0.7, and also satisfies a second inequality: yield strength YS(0) × work hardening coefficient >= 70.
4. The dual-phase type high-strength steel sheet according to any one of claims 1 and 2, characterized in that an average grain size of martensite is at most 5 µm, and an average grain size of ferrite is at most 10 µm.
5. The dual-phase type high-strength steel sheet according to any one of claims 1 and 2, characterized by satisfying a first inequality as: tensile strength (MPa) ×
total elongation (s) >= 18,000, and by satisfying a second inequality as follows: hole expansion ratio (d/d0) >= 1.2.
total elongation (s) >= 18,000, and by satisfying a second inequality as follows: hole expansion ratio (d/d0) >= 1.2.
6. The dual-phase type high-strength steel sheet according to any one of claims 1 and 2, characterized in that a plastic deformation (T) by at least one of: a tempering rolling and: a tension leveller, satisfies a following inequality:
2.5(YS(0)/TS'(5) - 0.5) + 15 >= T >= 2.5 (YS(0)/TS'(5) - 0.5) + 0.5
2.5(YS(0)/TS'(5) - 0.5) + 15 >= T >= 2.5 (YS(0)/TS'(5) - 0.5) + 0.5
7. A method of producing a dual-phase type high strength hot rolled steel sheet having high impact energy absorption properties, characterized in that after a continuous cast slab containing the steel compositions defined in any one of claims 1 and 2 is one of: fed directly from casting to a hot rolling step, and: hot rolled upon repeating after momentary cooling, it is subjected to hot rolling at a finishing temperature between Ar3 - 50°C and Ar3 + 120°C, cooled at an average cooling rate of more than 5°C/sec in a run-out table, and then cooled at a temperature of at: most 350°C.
8. The method according to claim 7, characterized in that at the finishing temperature for hot rolling in the range between Ar3 - 50°C and Ar3 + 120°C, the hot rolling is carried out so that a metallurgy parameter A
satisfies inequalities (1) and (2) below, a subsequent average cooling rate in the run-out table being at least 5°C/sec, and cooling is accomplished so that a relationship between said metallurgy parameter A and a cooling temperature (CT) satisfies inequality (3) below:
9 <= logA <= 18 (1) .DELTA.T <= 21 × logA - 61 (2) CT <= 6 × logA + 242 (3)
satisfies inequalities (1) and (2) below, a subsequent average cooling rate in the run-out table being at least 5°C/sec, and cooling is accomplished so that a relationship between said metallurgy parameter A and a cooling temperature (CT) satisfies inequality (3) below:
9 <= logA <= 18 (1) .DELTA.T <= 21 × logA - 61 (2) CT <= 6 × logA + 242 (3)
9. A method, characterized in that after a continuous cast slab containing the steel compositions defined in any one of claims 1 and 2 is one of: fed directly from casting to a hot rolling step, and: hot rolled upon reheating after momentary cooling, it is hot rolled, the hot rolled and subsequently coiled steel sheet is cold rolled after acid pickling, and during annealing in a continuous annealing step for preparation of a final product, it is heated to a temperature range between Ac1 and Ac3 and subjected to annealing while held in this temperature range for at least 10 seconds, and then cooled at a cooling rate of more than 5°C:/sec,
10. The method according to claim 9, after said continuous annealing step, and for subsequent cooling, it is cooled to a secondary cooling start temperature (Tq) in a range between 550°C-To at a primary cooling rate between 1 and 10°C/sec and then cooled to a secondary cooling end temperature (Te) at most equal to Tem determined by chemical compositions and annealing temperature (To), at a secondary cooling rate between 10 and 200°C/sec.
Applications Claiming Priority (11)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP8243497 | 1997-03-17 | ||
| JP9-82434 | 1997-03-17 | ||
| JP9-190299 | 1997-07-15 | ||
| JP9-190297 | 1997-07-15 | ||
| JP19029997 | 1997-07-15 | ||
| JP19029797A JP3530347B2 (en) | 1997-07-15 | 1997-07-15 | How to select a high-strength steel sheet with excellent dynamic deformation characteristics |
| JP22300897A JP3936440B2 (en) | 1997-08-06 | 1997-08-06 | High-strength steel sheet for automobiles with excellent collision safety and formability and its manufacturing method |
| JP9-223008 | 1997-08-06 | ||
| JP9-258938 | 1997-09-24 | ||
| JP25893897A JP3839928B2 (en) | 1997-07-15 | 1997-09-24 | Dual phase type high strength steel plate with excellent dynamic deformation characteristics |
| PCT/JP1998/001101 WO1998041664A1 (en) | 1997-03-17 | 1998-03-16 | Dual-phase high-strength steel sheet having excellent dynamic deformation properties and process for preparing the same |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| CA2283924A1 CA2283924A1 (en) | 1998-09-24 |
| CA2283924C true CA2283924C (en) | 2006-11-28 |
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ID=27524972
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| CA002283924A Expired - Lifetime CA2283924C (en) | 1997-03-17 | 1998-03-16 | Dual-phase type high-strength steel sheets having high impact energy absorption properties and a method of producing the same |
Country Status (7)
| Country | Link |
|---|---|
| EP (2) | EP2314729B2 (en) |
| KR (1) | KR100334949B1 (en) |
| CN (1) | CN1080321C (en) |
| AU (1) | AU717294B2 (en) |
| CA (1) | CA2283924C (en) |
| TW (1) | TW426742B (en) |
| WO (1) | WO1998041664A1 (en) |
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| JP3039842B2 (en) * | 1994-12-26 | 2000-05-08 | 川崎製鉄株式会社 | Hot-rolled and cold-rolled steel sheets for automobiles having excellent impact resistance and methods for producing them |
| JPH08176732A (en) * | 1994-12-27 | 1996-07-09 | Nkk Corp | Steel for nitriding having excellent machinability |
| JP3529178B2 (en) * | 1994-12-28 | 2004-05-24 | Ｊｆｅスチール株式会社 | Ultra-low carbon steel sheet with excellent shock absorption capacity |
| JP3090421B2 (en) * | 1996-07-22 | 2000-09-18 | 新日本製鐵株式会社 | Hot-rolled high-strength steel sheet for processing with excellent durability fatigue resistance |
-
1998
- 1998-03-16 CN CN98803465A patent/CN1080321C/en not_active Expired - Lifetime
- 1998-03-16 EP EP10181225.3A patent/EP2314729B2/en not_active Expired - Lifetime
- 1998-03-16 KR KR1019997008474A patent/KR100334949B1/en not_active IP Right Cessation
- 1998-03-16 AU AU63118/98A patent/AU717294B2/en not_active Expired
- 1998-03-16 CA CA002283924A patent/CA2283924C/en not_active Expired - Lifetime
- 1998-03-16 EP EP98907247.5A patent/EP0969112B2/en not_active Expired - Lifetime
- 1998-03-16 TW TW087103834A patent/TW426742B/en not_active IP Right Cessation
- 1998-03-16 WO PCT/JP1998/001101 patent/WO1998041664A1/en active IP Right Grant
Also Published As
| Publication number | Publication date |
|---|---|
| AU717294B2 (en) | 2000-03-23 |
| EP2314729B2 (en) | 2017-03-08 |
| CA2283924A1 (en) | 1998-09-24 |
| EP0969112B2 (en) | 2017-03-08 |
| EP0969112A1 (en) | 2000-01-05 |
| CN1080321C (en) | 2002-03-06 |
| AU6311898A (en) | 1998-10-12 |
| CN1251140A (en) | 2000-04-19 |
| KR20000076372A (en) | 2000-12-26 |
| EP2314729B1 (en) | 2012-02-08 |
| EP2314729A1 (en) | 2011-04-27 |
| EP0969112A4 (en) | 2003-05-21 |
| EP0969112B1 (en) | 2011-08-17 |
| KR100334949B1 (en) | 2002-05-04 |
| TW426742B (en) | 2001-03-21 |
| WO1998041664A1 (en) | 1998-09-24 |
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