Technical Field
[0001] The present invention relates to a cold-rolled steel sheet, an electro-galvanized
steel sheet, a hot-dip galvanized steel sheet, and an alloyed hot-dip galvanized steel
sheet, which are suitable as raw material steel sheets for molded products such as
building members, mechanical structural parts, automobile structural parts, etc.,
which are used at positions required to have structural strength, particularly, strength
and/or stiffness in deformation, and which are subjected to heat treatment for increasing
strength after processing such as pressing or the like, and a method of producing
these steel sheets.
[0002] 1. In the present invention, "excellent strain age hardenability" means that in aging
under conditions of holding at a temperature of 170°C for 20 min. after pre-deformation
with a tensile strain of 5%, the increment in deformation stress (represented by the
amount of BH = yield stress after aging - pre-deformation stress before aging) after
aging is 80 MPa or more, and the increment in tensile strength (represented by ΔTS
= tensile strength after aging - tensile strength before pre-deformation) after strain
aging (pre-deformation + aging) is 40 MPa or more.
Background Art
[0003] In producing a press-molded product of a thin steel sheet, a process of coating and
baking at lower than 200°C is used as a method in which a material having low deformation
stress before press forming to facilitate press forming, and then hardened after press
forming to increase the strength of a part. As a steel sheet for such coating and
baking, a BH steel sheet has been developed.
[0004] For example, Japanese Unexamined Patent Application Publication No. 55-141526 discloses
a method in which Nb is added according to the contents of C, N and Al of steel, Nb/(dissolved
C + dissolved N) by at% is limited in a specified range, and the cooling rate after
annealing is controlled to adjust dissolved C and dissolved N in a steel sheet. Also,
Japanese Examined Patent Application Publication No. 61-45689 discloses a method in
which baking hardenability is improved by adding Ti and Nb.
[0005] However, in order to improve deep drawability, strength of the raw material sheets
of the above-described steel sheets is decreased, and thus the steel sheets are not
always sufficient as structural materials.
[0006] Japanese Unexamined Patent Application Publication No. 5-25549 discloses a method
in which baking hardenability is improved by adding W, Cr and Mo to steel singly or
in a combination.
[0007] In the above-described conventional techniques, strength is increased by bake-hardening
due to the functions of small amounts of dissolved C and dissolved N in a steel sheet,
and it is well known that a BH(Bake-Hardening) steel sheet is used for increasing
only the yield strength of a material, not for increasing tensile strength. Therefore,
the conventional techniques have only the effect of increasing the deformation start
stress of a part, and the effect of increasing stress (tensile strength after forming)
required for deformation over the entire deformation region from the deformation start
to the deformation end is not said to be sufficient.
[0008] As a cold-rolled steel sheet having tensile strength increased after forming, for
example, Japanese Unexamined Patent Application Publication No. 10-310847 discloses
an alloying ho-dip galvanized steel sheet having tensile strength increased by 60
MPa or more by heat treatment in the temperature region of 200 to 450°C.
[0009] This steel sheet contains, by mass%, 0.01 to 0.08% of C, and 0.01 to 3.0% of Mn,
and at least one of W, Cr, and Mo in a total of 0.05 to 3.0%, and further contains
at least one of 0.005 to 0.1% of Ti, 0.005 to 0.1% of Nb and 0.005 to 0.1% of V according
to demand, and the microstructure of the steel is composed of ferrite or mainly composed
of ferrite.
[0010] However, this technique comprises forming a fine carbide in the steel sheet by heat
treatment after forming to effectively propagate a dislocation of stress applied during
pressing, thereby increasing the amount of strain. Therefore, heat treatment must
be performed in the temperature range of 220 to 370°C. There is thus the problem of
a necessary heat treatment temperature higher than general bake-hardening temperatures.
[0011] Furthermore, it is a very important problem that the body weight of an automobile
is decreased in relation to the recent regulation of exhaust gases due to global environmental
problems. In order to decrease the body weight of an automobile, it is effective to
increase the strength of the used steel sheet, i.e., use a high-tensile-strength steel
sheet, thinning the steel sheet used.
[0012] An automobile part using a high-tensile-strength thin steel sheet must exhibit a
sufficient property according to its function. The property depends upon the part,
and examples of the property include dent resistance, static strength against deformation
such as bending, twisting, or the like, fatigue resistance, impact resistance, etc.
Namely, the high-tensile-strength steel sheet used for an automobile part is required
to be excellent in such a property after forming. The properties are related to the.
strength of a steel sheet after forming, and thus the lower limit of strength of the
high-tensile-strength steel sheet used must be set for achieving thinning.
[0013] On the other hand, in the process for forming an automobile part, a steel sheet is
press-molded. If the steel sheet has excessively high strength in press forming, the
steel sheet causes the following problems: (1) deteriorating shape fixability; (2)
deteriorating ductility to cause cracking, necking, or the like during forming; and
(3) deteriorating dent resistance (resistance to a dent produced by a local compressive
load) when the sheet thickness is decreased. These problems thus inhibit the extension
of application of the high-tensile-strength steel sheet to automobile bodies.
[0014] As a means for overcoming the problems, a steel sheet composed of ultra-low-carbon
steel is known as a raw material, for example, for a cold-rolled steel sheet for an
external sheet panel, in which the content of C finally remaining in a solid solution
state is controlled to an appropriate range. This type of steel sheet is kept soft
during press forming to ensure shape fixability and ductility, and its yield stress
is increased by utilizing the strain aging phenomenon which occurs in the step of
coating and baking at 170°C for about 20 minutes after press forming, to ensure dent
resistance. This steel sheet is soft during press forming because C is dissolved in
steel, while dissolved C is fixed to a dislocation introduced in press forming in
the coating and baking step after press forming to increase the yield stress.
[0015] However, in this type of steel sheet, the increase in yield stress due to strain
age hardening is kept down from the viewpoint of prevention of the occurrence of stretcher
strain causing a surface defect. This causes the fault that the steel sheet actually
less contributes to a reduction in weight of a part.
[0016] On the other hand, a steel sheet composed of dissolved N to further increase the
amount of bake-hardening, and a steel sheet provided with a composite structure composed
of ferrite and martensite to further improve baking hardenability have been proposed
for applications in which the appearance is not so important.
[0017] For example, Japanese Unexamined Patent Application Publication No. 60-52528 discloses
a method of producing a high-strength steel thin sheet having good ductility and spot
weldability, in which steel containing 0.02 to 0.15% of C, 0.8 to 3.5% of Mn, 0.02
to 0.15% of P, 0.10% or less of Al, and 0.005 to 0.025% of N is hot-rolled by coiling
at a temperature of 550°C or less, cold-rolled, and then annealed by controlled cooling
and heat treatment. A steel sheet produced by the technique disclosed in Japanese
Unexamined Patent Application Publication No. 60-52528 has a mixed structure comprising
a low-temperature transformation product phase mainly composed of ferrite and martensite,
and having excellent ductility, and high strength is achieved by utilizing strain
aging due to positively added N during coating baking.
[0018] Although the technique disclosed in Japanese Unexamined Patent Application Publication
No. 60-52528 greatly increases yield stress YS due to strain age hardening, the technique
less increases tensile strength TS. Also, this technique causes large variations in
the increment in yield stress YS to cause large variations in mechanical properties,
and thus it cannot be expected that a steel sheet can be sufficiently thinned for
contributing to a reduction in weight of an automobile part, which is currently demanded.
[0019] Japanese Examined Patent Application Publication No. 5-24979 discloses a high-tensile-strength
cold-rolled steel thin sheet having baking hardenability which has a composition comprising
0.08 to 0.20% of C, 1.5 to 3.5% of Mn, and the balance composed of Fe and inevitable
impurities, and a structure composed of homogeneous bainite containing 5% or less
of ferrite, or bainite partially containing martensite. The cold-rolled steel sheet
disclosed in Japanese Examined Patent Application Publication No. 5-24979 is produced
by quenching in the temperature range of 200 to 400°C in the cooling process after
continuous annealing, and then slowly cooling to obtain a structure mainly composed
of bainite and having a large amount of bake-hardening which is not obtained by a
conventional method.
[0020] However, in the steel sheet disclosed in Japanese Examined Patent Application Publication
No. 5-24979, yield strength is increased after coating and baking to obtain a large
amount of bake-hardening which is not obtained a conventional method, while tensile
strength cannot be increased. Therefore, in application to a strength member, improvements
in fatigue resistance and impact resistance after forming cannot be expected. Therefore,
there is a problem in which the steel sheet cannot be used for applications greatly
required to have fatigue resistance and impact resistance, etc.
[0021] Also, Japanese Examined Patent Application Publication No. 61-12008 discloses a method
of producing a high-tensile-strength steel sheet having a high r value. This method
is characterized by annealing ultra-low-C steel used as a raw material in a ferrite-austenite
coexistence region after cold rolling. However, the resultant steel sheet has a high
r value and a high degree of baking hardenability (BH property), but the obtained
BH amount is about 60 MPa at most. Also, the yield point of the steel sheet is increased
after strain aging, but TS is not increased, thereby causing the problem of limiting
application to parts.
[0022] Furthermore, the above-described steel sheet exhibits excellent strength after coating
and baking in a simple tensile test, but produces large variations in strength during
plastic deformation under actual pressing conditions. Therefore, it cannot be said
that the steel sheet is sufficiently applied to parts required to have reliability.
[0023] With respect of a hot-rolled steel sheet among coating baked steel sheets for press
molded products, for example; Japanese Examined Patent Application Publication No.
8-23048 discloses a method of producing a hot-rolled steel sheet which is soft during
processing, and has tensile strength increased by coating and baking after processing
to be effective to improve fatigue resistance.
[0024] In this technique, steel contains 0.02 to 0.13 mass % of C, and 0.0080 to 0.0250
mass % of N, and the finisher deliver temperature and the coiling temperature are
controlled to leave a large amount of dissolved N in the steel, thereby forming a
composite structure as a metal structure mainly composed of ferrite and martensite.
Therefore, an increase of 100 MPa or more in tensile strength is achieved at the heat
treatment temperature of 170°C after forming.
[0025] Japanese Unexamined Patent Application Publication No. 10-183301 discloses a hot-rolled
steel sheet having excellent baking hardenability and natural aging resistance, in
which the C and N contents are limited to 0.01 to 0.12 mass % and 0.0001 to 0.01 mass
%, respectively, and the average crystal grain diameter is controlled to 8 µm or less
to ensure a BH amount of as high as 80 MPa or more, and suppress the AI amount to
45 MPa or less.
[0026] However, this steel sheet is a hot-rolled sheet, and is thus difficult to obtain
a high r value because the ferrite aggregation texture is made random due to austeniste-ferrite
transformation. Therefore, the steel sheet cannot be said to have sufficient deep
drawability.
[0027] Furthermore, even if the hot-rolled steel sheet obtained by this technique is used
as a starting material for cold rolling and recrystallization annealing, the increase
in tensile strength obtained after forming and heat treatment is not always equivalent
to a hot-rolled steel sheet, and a BH amount of as high as 80 MPa or more cannot be
always obtained. This is because the microstructure of the cold -rolled steel becomes
different from that of hot-rolled one due to cold rolling and recrystallization annealing,
and strain greatly accumulates during cold rolling to easily form a carbide, a nitride
or a carbonitride, thereby changing the states of dissolved C and dissolved N.
[0028] In consideration of the above-described present conditions, an object of the present
invention is to provide a cold-rolled steel sheet and a hot-dip galvanized steel sheet
(including an alloyed steel sheet) for deep drawing, which have excellent deep drawability,
TS x r value ≥ 750 MPa, and excellent strain aging hardenability (BH ≥ 80 MPa and
ΔTS ≥ 40 MPa), and an effective method of producing these steel sheets.
[0029] Another object of the present invention is to solve the above problems of the conventional
techniques and provide a high-tensile-strength cold-rolled steel sheet which is suitable
for automobile parts required to have high moldability, softness and high moldability,
and stable material properties, and which can easily be molded to an automobile part
having a complicated shape without producing shape defects such as spring back, twisting,
and curving, and cracking, etc., and which has sufficient strength as an automobile
part after heat treatment of a molded automobile part to permit sufficient contribution
to a reduction in body weight of an automobile, a high r value of 1.2 or more, and
excellent strain age hardenability. A further object of the present invention is to
provide an industrial production method capable of producing the steel sheet at low
cost without disturbing its shape.
Disclosure of Invention
[0030] In order to achieve the objects, the inventors produced various steel sheets having
different compositions under various production conditions, and experimentally evaluated
various material properties. As a result, it was found that both moldability and hardenability
after forming can be improved by using as a strengthening element N, which has not
be positively used before in a field requiring high processability, and effectively
using the great strain age hardening phenomenon manifested by the action of the strengthening
element.
[0031] The inventors also found that in order to advantageously use the strain age hardening
phenomenon due to N, the strain age hardening phenomenon due to N must be advantageously
combined with a condition for coating and baking an automobile, or further positively
combined with a heat treatment condition after forming. It was thus found to be effective
to appropriately control the hot rolling condition, the cold rolling and the cold
rolling annealing condition to control the microstructure of a steel sheet and the
amount of dissolved N in certain ranges. It was also found that in order to stably
manifest the strain age hardening phenomenon due to N, it is important to control
the Al content of the composition according to the N content.
[0032] The inventors further found that in order to obtain a high r value, the C content
is decreased, continuous annealing is performed in the ferrite-austenite two-phase
temperature region, and subsequent cooling is controlled to form a structure containing
an acicular ferrite phase at an area ratio of 5% or more in the ferrite phase. Such
a combination of the microstructure and the appropriate amount of dissolved N was
found to enable the achievement of a cold-rolled steel sheet having a high r value,
excellent press moldability, and excellent strain age hardenability. This was also
found to permit sufficient use of N without causing the problem of natural aging deterioration,
which is the problem of a conventional bake-hardening steel sheet.
[0033] Namely, the inventors found that by suing N as a strengthening element, controlling
the Al content according to the N content in an appropriate range, and appropriately
controlling the hot rolling condition and the cold rolling annealing condition to
appropriately control the microstructure and dissolved N, it is possible to obtain
a steel sheet having a high r value and excellent moldability as compared with conventional
solid-solution strengthening-type C-Mn steel sheets and precipitation strengthening-type
steel sheets, and strain age hardenability, which is not possessed by the conventional
steel sheets.
[0034] A steel sheet of the present invention exhibits higher strength after coating and
baking in a simple tensile test, as compared with a conventional steel sheet, and
exhibits small variations in strength in plastic deformation under actual pressing
conditions and stable part strength, thereby enabling application to parts required
to have reliability. For example, a portion where large strain is applied to decrease
the thickness has higher hardenability than other portions, and is considered homogeneous
when being evaluated based on a surcharge load ability of (thickness) x (strength),
thereby stabilizing strength as a part.
[0035] As a result of further intensive research for achieving the objects, the inventors
found the following:
1) In order to increase tensile strength after forming and heat treatment, a new dislocation
must be introduced for progressing tensile deformation. The movement of the dislocation
introduced by pre-deformation must be prevented by interaction between the dislocation
introduced by forming and an interstitial element or a precipitate even when upper
yield stress is attained.
2) In order to obtain the above interaction by forming a carbide, a nitride or a carbonitride
of W, Cr, Mo, Ti, Nb, Al or the like, the heat treatment temperature after forming
must be increased to 200°C or more. Therefore, it is more advantageous to positively
use the interstitial element or a Fe carbide or Fe nitride because the heat treatment
temperature after forming is decreased.
3) Of interstitial elements, dissolved N has the higher interaction with a dislocation
introduced by forming than dissolved C even when the heat treatment temperature after
forming is decreased, and thus a dislocation introduced by pre-deformation less moves
when upper yield stress is attained.
4) Although dissolved N is present in crystal grains and crystal grain boundaries
in steel, the increase in strength after forming and heat treatment increases as the
area of the crystal grain boundaries increases. Namely, the smaller crystal grain
diameter is advantageous.
5) In order to increase the crystal grain boundary area, it is advantageous to add
a combination of Nb and B and cool immediately after the end of hot rolling, suppressing
normal grain growth of ferrite grains after the end of hot rolling and suppressing
grain growth by recrystallization annealing after cold rolling.
[0036] The present invention has been achieved based on the above findings. The findings
were obtained from the experiment described below.
Experiment 1
[0037] A sheet bar (thickness: 30 mm) having a composition containing, by mass %, 0.0015%
of C, 0.0010% of B, 0.015 of Si, 0.5% of Mn, 0.03% of P, 0.08% of S and 0.011% of
N, 0.005 to 0.05% of Nb and 0.005 to 0.03% of Al, and the balance composed of Fe and
inevitable impurities was uniformly heated at 1150°C, hot-rolled by three passes so
that the temperature of the final pass was 900°C higher than the Ar
3 transformation point, and then cooled with water for 0.1 second. Then, the sheet
bar was subjected to heat treatment corresponding to coiling at 500°C for 1 hour.
[0038] The thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with
a rolling reduction ratio of 82,5%, recrystallized and annealed at 800°C for 40 seconds,
and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a tensile test
specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling
direction, and tensile strength was measured with a strain rate of 0.02/s by using
a general tensile testing machine. Also, tensile strain of 10% was applied to a tensile
test specimen of JIS No. 5 separately obtained from the cold-rolled sheet in the rolling
direction, and then the specimen was subjected to a normal tensile test after heat
treatment at 120°C for 20 minutes. The difference between the tensile strength of
the specimen obtained from the cold-rolled sheet and the tensile strength of the specimen
heat treated at 120°C for 20 minutes after application of 10% tensile strain was considered
as the increase in tensile strength after forming (ΔTS).
[0039] Fig. 1 shows the results of measurement of the relation between the steel compositions
(N% - 14/93•Nb% - 14/27•Al%-14/11•B%) and ΔTS.
[0040] The figure indicates that ΔTS becomes 60 MPa or more when the value of (N% - 14/93•Nb%
- 14/27•Al% - 14/11•B%) satisfies 0.0015 mass %.
Experiment 2
[0041] A sheet bar (thickness: 30 mm) having a composition containing, by mass %, 0.0010%
of C, 0.02 of Si, 0.6% of Mn, 0.01% of P, 0.009% of S and 0.012% of N, 0.01% of Al,
0.015% of Nb, 0.00005 to 0.0025% of B, and the balance composed of Fe and inevitable
impurities was uniformly heated at 1100°C, hot-rolled by three passes so that the
temperature of the final pass was 920°C higher than the Ar
3 transformation point, and then cooled with water for 0.1 second. Then, the sheet
bar was subjected to heat treatment corresponding to coiling at 450°C for 1 hour.
[0042] The thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with
a rolling reduction ratio of 82,5%, recrystallized and annealed at 820°C for 40 seconds,
and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a tensile test
specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling
direction, and tensile strength was measured with a strain rate of 0.02/s by using
a general tensile testing machine. Also, tensile strain of 10% was applied to a tensile
test specimen of JIS No. 5 separately obtained from the cold-rolled sheet in the rolling
direction, and then the specimen was subjected to a normal tensile test after heat
treatment at 120°C for 20 minutes.
[0043] Fig. 2 shows the results of measurement of the relation between the B content of
steel and ΔTS. This figure indicates that with a B content of 0.0005 to 0.0015 mass
%, a high ΔTS of 60 MPa or more can be obtained.
[0044] As a result of observation of the microstructure, it was also found that by adding
a combination of Nb and B to make fine crystal grains, a high ΔTS can be obtained.
[0045] Namely, with a B content of less than 0.0005 mass. %, the effect of making fine crystal
grains by adding a combination with Nb is small. On the other hand, with a B content
of over 0.0015 mass %, the amount of B segregated in the grain boundaries and the
vicinities thereof is increased to decrease the amount of effective dissolved N because
of the strong interaction between B atoms and N atoms, thereby possibly decreasing
ΔTS.
Experiment 3
[0046] A sheet bar (thickness: 30 mm) of each of steel A having a composition containing,
by mass %, 0.0010% of C, 0.012% of N, 0.0010% of B, 0.01% of Si, 0.5% of Mn, 0.03%
of P, 0.008% of S, 0.014% of Nb, 0.01% of Al, and the balance composed of Fe and inevitable
impurities, and steel B having a composition containing, by mass %, 0.010% of C, 0.0012%
of N, 0.0010% of B, 0.01% of Si, 0.5% of Mn, 0.03% of P, 0.008% of S, 0.014% of Nb,
0.01% of Al, and the balance composed of Fe and inevitable impurities was uniformly
heated at 1150°C, hot-rolled by three passes so that the temperature of the final
pass was 910°C higher than the Ar
3 transformation point, and then cooled with a gas for 0.1 second. Then, each of the
sheet bars was subjected to heat treatment corresponding to coiling at 600°C for 1
hour.
[0047] Each of the thus-obtained hot-rolled sheets having a thickness of 4 mm was cold-rolled
with a rolling reduction ratio of 82,5%, recrystallized and annealed at 880°C for
40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%.
[0048] Then, a tensile test specimen of JIS No. 5 was obtained from each of the resultant
cold-rolled sheets in the rolling direction, and tensile strength was measured with
a strain rate of 0.02/s by using a general tensile testing machine. Also, tensile
strain of 10% was applied to a tensile test specimen of JIS No. 5 separately obtained
from each of the cold-rolled sheets in the rolling direction, and then the specimen
was subjected to a normal tensile test after heat treatment at various temperatures
for 20 minutes.
[0049] Fig. 3 shows the results of measurement of the influence of the heat treatment temperature
after forming on ΔTS. This figure indicates that in the relatively low temperature
region of heat treatment temperatures of 200°C or less after forming, the ultra-low
carbon steel A having a high N content exhibits higher ΔTS than the semi-ultra low
carbon steel B having a low N content, and while in the high temperature region, both
steel materials exhibit substantially the same ΔTS. There experimental results reveal
that in order to ensure ΔTS in the low temperature region, it is effective to use
dissolved N.
[0050] Fig. 4 shows the results of measurement of the influences of the crystal grain diameter
d and steel compositions (N%-14/93•Nb% - 14/27•Al% - 14/11•B%) on a decrease (ΔE1)
in elongation by natural aging and an increase in tensile strength (ΔTS) after forming.
The decrease (ΔE1) in elongation was evaluated by the difference between the total
elongation measured with the test specimen of JIS NO. 5 obtained from each of the
cold-rolled sheets in the rolling direction, and the total elongation measured with
the separately obtained test specimen after holding at 100°C for 8 hours for accelerating
natural aging.
[0051] Fig. 4 indicates that when the value of (N% - 14/93•Nb%-14/27•A1% - 14/11•B%) is
0.0015 mass % or more, and the crystal grain diameter d is 20 µm or less, both high
ΔTS and low ΔE1 can be achieved.
Experiment 4
[0052] A sheet bar of steel containing 0.0015% of C, 0.30 of Si, 0.8% of Mn, 0.03% of P,
0.005% of S and 0.012% of N, and 0.02 to 0.08% of Al was uniformly heated at 1050°C,
hot-rolled by seven passes so that the temperature of the final pass was 670°C, and
then recrystallized and annealed at 700°C for 5 hours. The thus-obtained hot-rolled
sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of
82,5%, recrystallized and annealed at 875°C for 40 seconds, and then temper-rolled
with a rolling reduction of 0.8%.
Then, a tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled
sheet in the rolling direction, and TS x r value and ΔTS were measured with a strain
rate of 3 x 10
-3/s by using a general tensile testing machine. The results are shown in Fig. 5. In
this figure, when N/Al ≥ 0.03 is satisfied, TS x r value ≥ 750 and ΔTS ≥ 40 MPa are
achieved. It was also confirmed that when N/Al ≥ 0.03, BH ≥ 80 MPa is attained.
Experiment 5
[0053] A sheet bar of steel containing 0.0015% of C, 0.0010% of B, 0.01 of Si, 0.5% of Mn,
0.03% of P, 0.008% of S and 0.011% of N, 0.005 to 0.05% of Nb, and 0.005 to 0.03%
of Al was uniformly heated at 1000°C, hot-rolled by seven passes so that the temperature
of the final pass was 650°C, and then recrystallized and annealed at 800°C for 60
seconds. The thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled
with a rolling reduction ratio of 82,5%, recrystallized and annealing at 880°C for
40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a
tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet
in the rolling direction, and TS x r value, BH and ΔTS were measured with a strain
rate of 3 x 10
- 3/s by using a general tensile testing machine. The relations between the measured
values and N/(Al+Nb+B) are shown in Fig. 5. In this experiment, steel containing 0.005
to 0.05% of Nb and 0.0010% of B was used. This figure indicates that in the range
of N/(Al+Nb+B) ≥ 0.30, BH ≥ 80 MPa, ΔTS ≥ 60 MPa, and TS x r value ≥ 850 are achieved.
Experiment 6
[0054] A sheet bar of steel containing 0.0010% of C, 0.02 of Si, 0.6% of Mn, 0.01% of P,
0.009% of S and 0.015% of N, 0.01% of Al, 0.015% of Nb and 0.0001 to 0.0025% of B
was uniformly heated at 1050°C, hot-rolled by seven passes so that the temperature
of the final pass was 680°C, and then recrystallized and annealed at 850°C for 5 hours.
The thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with
a rolling reduction ratio of 82,5%, recrystallized and annealed at 880°C for 40 seconds,
and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a tensile test
specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling
direction, and TS x r value, BH and ΔTS were measured with a strain rate of 3 x 10
-3/s by using a general tensile testing machine. The relations between the measured
values and the B content are shown in Fig. 6.
[0055] This figure indicates that in the B content range of 0.0003 to 0.0015%, BH ≥ 80 MPa,
ΔTS ≥ 60 MPa, which is higher than the case of B < 0.0003%, and TS x r value ≥ 850
are achieved. As a result of observation of the microstructure, it was also confirmed
that in this B range, crystal grains are significantly made fine.
[0056] The results shown in Figs. 5 and 6 indicate that in the range of N/(Al+Nb+B) ≥ 0.30
wherein B ≥ 0.0003%, the crystal grains are further made fine by combining Nb, and
ΔTS and the level of TS x r value are further improved. When B < 0.0003%, the effect
of making fine crystal grains by combining Nb is not exhibited. On the other hand,
when B > 0.0015%, properties further deteriorate. This is possibly due to the fact
that the amount of B segregated in the grain boundaries and the vicinities thereof
is increased to decrease the amount of effective dissolved N due to the strong interaction
between B and N atoms. The same research as described above was carried out for the
case in which Ti and V were added in place of Nb, and it was confirmed that the same
effect as Nb could be obtained. The present invention has been achieved based on the
above-described findings, and the gist of the invention was follows.
[0057] In a first aspect of the present invention, a cold-rolled steel sheet having excellent
strain age hardenability comprises a composition, by mass %, comprising:
C: 0.15% or less;
Si: 1.0% or less;
Mn: 2.0% or less;
P: 0.1% or less;
S: 0.01% or less;
Al: 0.005 to 0.030%; and
N: 0.0050 to 0.0400%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010% or more, and
the balance is composed of Fe and inevitable impurities.
[0058] In the first aspect of the present invention, the above-described composition preferably
comprises the compositions in the following ranges. Namely, a cold-rolled steel sheet
having excellent strain age hardenability comprises a composition, by mass %, comprising:
C: less than 0.01%;
Si: 0.005 to 1.0%;
Mn: 0.01 to 1.5%;
P: 0.1% or less;
S: 0.01% or less;
Al: 0.005 to 0.030%; and
N: 0.005 to 0.040%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010% or more, and
the balance is composed of Fe and inevitable impurities.
[0059] In the first aspect of the present invention, the composition, by mass %, preferably
further comprises:
B: 0.0001 to 0.0030%; and
Nb: 0.005 to 0.050%;
wherein the ranges of B and Nb satisfy the following equations (1) and (2) :
[0060] In the first aspect of the present invention, the above composition, by mass %, preferably
further comprises at least one of Cu, Ni and Mo in a total of 1.0% or less according
to demand.
[0061] In the first aspect of the present invention, the steel sheet preferably has a crystal
grain diameter of 20 µm or less.
[0062] In the first aspect of the present invention, strength after forming is preferably
increased by 60 MPa or more by heat treatment in the low temperature region of 120
to 200°C.
[0063] In the first aspect of the present invention, the surface of the cold-rolled steel
sheet may be coated by electro-galvanization, hot-dip galvanization, or alloying hot-dip
galvanization.
[0064] In a second aspect of the present invention, a method of producing a cold-rolled
steel sheet having excellent strain age hardenability comprises hot-rolling a steel
slab under conditions in which the slab is cooled immediately after the end of finish
rolling and coiled at a coiling temperature of 400 to 800°C, cold-rolling the hot-rolled
sheet with a rolling reduction ratio of 60 to 95%, and then performing recrystallization
annealing at a temperature of 650 to 900°C, wherein the steel slab has a composition,
by mass %, comprising:
C: less than 0.01%;
Si: 0.005 to 1.0%;
Mn: 0.01 to 1.5%;
P: 0.1% or less;
S: 0.01% or less;
Al: 0.005 to 0.030%; and
N: 0.005 to 0.040%;
wherein N/Al is 0.30 or more, and the balance is substantially composed of Fe.
[0065] In the second aspect of the present invention, the composition, by mass %, preferably
further comprises:
B: 0.0001 to 0.0030%; and
Nb: 0.005 to 0.050%;
wherein the ranges of B and Nb satisfy the following equations (1) and (2):
[0066] In the second aspect of the present invention, in the heating-up step in recrystallization
annealing, the temperature is preferably increased at a rate of 1 to 20°C/s in the
temperature region from 500°C to the recrystallization temperature.
[0067] In the second aspect of the present invention, hot-dip galvanization and heat alloying
may be performed after the recrystallization annealing.
[0068] In a third aspect of the present invention, a cold-rolled deep drawing steel sheet
having excellent strain age hardenability comprises a composition, by mass %, comprising:
C: 0.01% or less;
Si: 1.0% or less;
Mn: 0.01 to 1.5%;
P: 0.1% or less;
S: 0.01% or less;
Al: 0.005 to 0.020%; and
N: 0.0050 to 0.040%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010% or more, the
balance is composed of Fe and inevitable impurities, and Ts x r value is 750 MPa or
more.
[0069] In the third aspect of the present invention, the composition, by mass %, preferably
further comprises:
B: 0.0001 to 0.0030%; and
Nb: 0.005 to 0.050%;
wherein the ranges of B and Nb satisfy the following equations (1) and (2):
[0070] In the third aspect of the present invention, the composition, by mass %, preferably
further comprises at least one of the following:
B: 0.0001 to 0.0030%;
Nb: 0.005 to 0.050%;
Ti: 0.005 to 0.070%; and
V: 0.005 to 0.10%;
wherein N/(Al+Nb+Ti+V+B) is 0.30 or more, and the amount of dissolved N is 0.0010%
or more.
[0071] In a fourth aspect of the present invention, a method of producing a cold-rolled
deep drawing steel sheet having excellent strain age hardenability comprises heating
a steel raw material to 950°C or more, roughly rolling the raw material so that the
finisher delivery temperature is Ar
3 to 1000°C, finish-rolling the material while lubricating it in the temperature region
of 600°C to Ar
3, coiling the rolled sheet in which the total reduction ratio by rolling starting
from rough rolling to finish rolling is 80% or more, recrystallizing and annealing
the hot-rolled sheet, cold-rolling the rolled sheet with a rolling reduction ratio
of 60 to 95%, and then recrystallizing and annealing the resultant cold-rolled sheet,
wherein the steel raw material has a composition, by mass %, comprising:
C: 0.01% or less;
Si: 0.005 to 1.0%;
Mn: 0.01 to 1.0%;
P: 0.1% or less;
S: 0.01% or less;
Al: 0.005 to 0.030%;
N: 0.005 to 0.040%; and
at least one of the following:
B: 0.0003 to 0.0030%;
Nb: 0.005 to 0.050%;
Ti: 0.005 to 0.070%; and
V: 0.005 to 0.10%;
wherein N/(Al+Nb+Ti+V+B) is 0.30 or more.
[0072] In a fifth aspect of the present invention, a high-tensile-strength cold-rolled steel
sheet having excellent moldability, strain age hardenability and natural aging resistance
comprises a composition, by mass %, comprising:
C: 0.0015 to 0.025%;
Si: 1.0% or less;
Mn: 2.0% or less;
P: 0.1% or less;
S: 0.02% or less;
Al: 0.02% or less;
N: 0.0050 to 0.0250%; and
one or both of the following:
B: 0.0005 to 0.0050%; and
Nb: 0.002 to 0.050%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010% or more, the
balance is composed of Fe and inevitable impurities, the structure is composed of
an acicular ferrite phase at an area ratio of 5% or more and a ferrite phase having
an average crystal grain diameter of 20 µm or less, and the r value is 1.2 or more.
[0073] In the fifth aspect of the present invention, the composition preferably further
comprises at least one of the following groups a to c:
Group a: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group b: one or both of Ti and V in a total of 0.1% or less; and
Group c: one or both of Ca and REM in a total of 0.0010 to 0.010%.
[0074] In a sixth aspect of the present invention, a method of producing a high-tensile-strength
cold-rolled steel sheet having a r value of 1.2 or more, and excellent moldability,
strain age hardenability and natural aging resistance comprises:
the hot-rolling step of roughly rolling a steel slab by heating to a slab heating
temperature of 1000°C or more to form a sheet bar, finish-rolling the sheet bar so
that the finisher delivery temperature is 800°C or more, and coiling the finish-rolled
sheet at a coiling temperature of 650°C or less to form a hot-rolled sheet;
the cold rolling step of pickling and cold-rolling the hot-rolled sheet to form a
cold-rolled sheet; and
the cold-rolled sheet annealing step of continuously annealing the cold-rolled sheet
at a temperature in the ferrite-austenite two-phase region, and cooling the annealed
sheet to the temperature region of 500°C or less at a cooling rate of 10 to 300°C/s;
wherein the steel slab has a composition, by mass %, comprising at least one of:
C: 0.0015 to 0.025%;
Si: 1.0% or less;
Mn: 2.0% or less;
P: 0.1% or less;
S: 0.02% or less;
Al: 0.02% or less;
N: 0.0050 to 0.0250%;
B: 0.0003 to 0.0050%; and
Nb: 0.002 to 0.050%;
wherein N/Al is 0.30 or more.
[0075] In the sixth aspect of the present invention, the composition preferably further
comprises, by mass %, at least one of the following groups a to c:
Group a: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group b: one or both of Ti and V in a total of 0.1% or less; and
Group c: one or both of Ca and REM in a total of 0.0010 to 0.010%.
[0076] In a seventh aspect of the present invention, a high-tensile-strength cold-rolled
steel sheet having a high r value and excellent strain age hardenability and natural
aging resistance comprises a composition, by mass %, comprising:
C: 0.025 to 0.15%;
Si: 1.0% or less;
Mn: 2.0% or less;
P: 0.08% or less;
S: 0.02% or less;
Al: 0.02% or less; and
N: 0.0050 to 0.0250%;
wherein N/Al is 0.30 or more, the amount of dissolved N is 0.0010% or more, the
balance is composed of Fe and inevitable impurities, the structure is composed of
a ferrite phase having an average crystal grain diameter of 10 µm or less at an area
ratio of 80% or more and a martensite phase as a second phase at an area ratio of
2% or more, and the r value is 1.2 or more.
[0077] In the seventh aspect of the present invention, the composition preferably further
comprises at least one of the following groups d to g:
Group d: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group e: at least one of Nb, Ti and V in a total of 0.1% or less;
Group f: 0.0030% or less of B; and
Group g: one or both of Ca and REM in a total of 0.0010 to 0.010%.
[0078] In an eighth aspect of the present invention, a method of producing a high-tensile-strength
cold-rolled steel sheet having a r value of as high as 1.2 or more, and excellent
strain age hardenability and natural aging resistance comprises:
the hot-rolling step of roughly rolling a steel slab by heating to a slab heating
temperature of 1000°C or more to form a sheet bar, finish-rolling the sheet bar so
that the finisher delivery temperature is 800°C or more, and coiling the finish-rolled
sheet at a coiling temperature of 650°C or less to form a hot-rolled sheet;
the cold rolling step of pickling and cold-rolling the hot-rolled sheet to form a
cold-rolled sheet; and
the cold-rolled sheet annealing step of box-annealing the cold-rolled sheet at an
annealing temperature of the recrystallization temperature to 800°C, then continuously
annealing the annealed sheet at an annealing temperature of Ac1 transformation point to (Ac3 transformation point - 20°C), and then cooling the sheet to the temperature region
of 500°C or less at a cooling rate of 10 to 300°C/s;
wherein the steel slab has a composition, by mass %, comprising at least one of:
C: 0.025 to 0.15%;
Si: 1.0% or less;
Mn: 2.0% or less;
P: 0.08% or less;
S: 0.02% or less;
Al: 0.02% or less; and
N: 0.0050 to 0.0250%;
wherein N/Al is 0.30 or more.
[0079] In the eighth aspect of the present invention, over aging after the cooling step
of the continuous annealing is preferably carried out in the temperature range of
350°C or more for a residual time of 20 seconds or more.
[0080] In the eighth aspect of the present invention, the composition preferably further
comprises, by mass %, at least one of the following groups d to g:
Group d: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group e: at least one of Nb, Ti and V in a total of 0.1% or less;
Group f: 0.0030% or less of B; and
Group g: one or both of Ca and REM in a total of 0.0010 to 0.010%.
Brief Description of the Drawings
[0081]
Fig. 1 shows the relation between steel compositions (N% - 14/93•Nb% - 14/27•Al% -
14/11•B%) and the increase in tensile strength (ΔTS) after forming.
Fig. 2 shows the relation between the B content and ΔTS of steel containing a combination
Nb and B.
Fig. 3 shows comparison of the difference in increase in tensile strength by heat
treatment after forming in a low temperature region between steel B (conventional
steel) containing a large amount of dissolved C and steel A (steel of this invention)
containing a large amount of dissolved N.
Fig. 4 shows the influence of the crystal grain diameter d and steel compositions
(N% - 14/93•Nb% - 14/27•Al%-14/11•B%) on the decrease in elongation (ΔE1) due to natural
aging and the increase in tensile strength (ΔTS) after forming.
Fig. 5 shows the relations between Ts x r value, BH, ΔTS and N/(Al+Nb+B).
Fig. 6 shows the relations between Ts x r value, BH, ΔTS and the B amount.
Best Mode for Carrying Out the Invention
[0082] Description will now be made of the reasons for limiting compositions to the ranges
below in accordance with a first embodiment of the present invention.
C: less than 0.01 mass %
[0083] From the viewpoint of excellent deep drawability and press moldability, C is advantageously
as small as possible. Also, redissolution of NbC proceeds in the annealing step after
cold rolling to increase the amount of dissolved C in crystal grains, thereby easily
causing deterioration in natural aging resistance. Therefore, the C amount is preferably
suppressed to less than 0.01 mass %, more preferably 0.0050 mass % or less, and most
preferably 0.0030 mass % or less.
Si: 0.005 to 1.0 mass %
[0084] Si is a useful composition for suppressing a decrease in elongation, and improving
strength. However, with a Si content of less than 0.005 mass %, the effect of addition
of Si is insufficient, while with a Si content of over 1.0 mass %, surface properties
deteriorate to deteriorate ductility. Therefore, the Si content is limited to the
range of 0.005 to 1.0 mass %, and preferably the range of 0.01 to 0.75 mass %. Mn:
0.01 to 1.5 mass %
[0085] Mn not only is useful as a strengthening composition for steel, but also has the
function to suppress embrittlement with S due to the formation of MnS. However, with
a Mn content of less than 0.01 mass %, the effect of addition of Mn is insufficient,
while with a Mn content of over 1.5 mass %, surface properties deteriorate to deteriorate
ductility. Therefore, the Mn content is limited to the range of 0.01 to 1.5 mass %,
and preferably the range of 1.10 to 0.75 mass %.
P: 0.01 mass % or less
[0086] P is a solid solution strengthening element which effectively contributes to reinforcement
of steel. However, with a P content of over 0.01 mass %, deep drawability deteriorates
due to the formation of phosphide such as (FeNb)
xP or the like. Therefore, P is limited to 0.10 mass % or less.
S: 0.01 mass % or less
[0087] With a high S content, the amount of inclusions is increased to deteriorate ductility.
Therefore, contamination with S is preferably prevented as much as possible, but an
S content up to 0.01 mass % is allowable.
Al: 0.005 to 0.030 mass %
[0088] Al is added as a deoxidizer for improving the yield of carbonitride forming components.
However, with an Al content of less than 0.005 mass %, the effect is insufficient,
while with an Al content of over 0.030 mass %, the amount of N to be added to steel
is increased to easily cause slab defects during steel making. Therefore, Al is contained
in the range of 0.005 to 0.030 mass %.
N: 0.005 to 0.040 mass %
[0089] In the present invention, N is an important element which plays the role of imparting
strain age hardenability to a steel sheet. However, with an N content of less than
0.005 mass %, a sufficient strain age hardenability cannot be obtained, while with
an N content of as high as over 0.040 mass %, press moldability deteriorates. Therefore,
N is contained in the range of 0.005 to 0.040 mass %, and preferably in the range
of 0.008 to 0.015 mass %.
B: 0.0001 to 0.003 mass %
[0090] B is added in a combination with Nb to exhibit the function to effectively make fine
the hot-rolled structure and the cold-rolled recrystallized annealed structure and
to improve cold-work embrittlement resistance. However, with a B content of less than
0.0001 mass %, the sufficient effect of making fine the structures cannot be obtained,
while with a B content of over 0.003 mass %, the amount of BN precipitate is increased,
and dissolution in the slab heating step is hindered. Therefore, B is contained in
the range of 0.0001 to 0.003 mass %, preferably in the range of 0.0001 to 0.0015 mass
%, and more preferably in the range of 0.0007 to 0.0012 mass %.
Nb: 0.005 to 0.050 mass %
[0091] Nb is added in a combination with B to contribute to refinement of the hot-rolled
structure and the cold-rolled recrystallized annealed structure, and have the function
to fix dissolved C as NbC. Furthermore, Nb forms a nitride NbN to contribute to refinement
of the cold-rolled recrystallized annealed structure. However, with a Nb content of
less than 0.005 mass %, not only it becomes difficult to precipitate and fix dissolved
C, but also the hot-rolled structure and the cold-rolled, recrystallized, annealed
structure are not sufficiently made fine, while with a Nb content of over 0.050 mass
%, ductility deteriorates. Therefore, Nb is contained in the range of 0.005 to 0.050
mass %, and preferably 0.010 to 0.030 mass %.
[0092] As described above, Nb has the function to fix dissolved C as NbC, and forms a nitride
NbN. Similarly, Al and B form A1N and BN, respectively. Therefore, in order to ensure
the sufficient amount of dissolved N and sufficiently decrease the amount of dissolved
C, it is important to satisfy the following relations (1) and (2):
[0093] In the present invention, in order to obtain a high strain aging property and prevent
aging deterioration, the crystal grain diameter is preferably decreased.
[0094] Namely, as described above with reference to Fig. 4, even when (N% - 14/93•Nb% -
14/27•Al%. - 14/11•B%) ≥ 0.0015 mass %, i.e., when a relatively large amount of dissolved
N is contained, ΔE1 can be suppressed to 2.0% or less by decreasing the crystal grain
diameter d to 20 µm or less. The crystal grain diameter d is more preferably decreased
to 15 µm or less. This is because, as shown in Fig. 4, ΔE1 can be suppressed to 2.0%
or less by decreasing the crystal grain diameter d to 15 µm or less.
[0095] The production conditions according to a second embodiment of the present invention
will be described.
[0096] Steel having the above-described suitable composition is melted by a known melting
method such as a converter or the like, and a steel slab is formed by an ingot making
method or a continuous casting method.
[0097] Then, the steel slab is heated and soaked, and then hot-rolled to form a hot-rolled
sheet. In the present invention, the heating temperature of hot rolling is not specified,
but the heating temperature of hot rolling is preferably set to 1300°C or less. This
is because it is advantageous to fix and precipitate dissolved C as a carbide in order
to improve deep drawability. In order to further improve processability, the heating
temperature is preferably set to 1150°C or less. However, with a heating temperature
of less than 900°C, improvement in processability is saturated to conversely increase
the rolling load in hot rolling, thereby increasing the danger of causing a rolling
trouble. Therefore, the lower limit of the heating temperature is preferably 900°C.
[0098] The total rolling reduction ratio of hot rolling is preferably 70% or more. This
is because with a toal rolling reduction ratio of less than 70%, the crystal grains
of the hot-rolled sheet are not sufficiently made fine.
[0099] During hot rolling, finish rolling is preferably finished in the temperature region
of 650 to 960°C, and the finishing temperature of hot-rolling may be in the γ region
above the Ar
3 transformation point, or the α region below the Ar
3 transformation pint. With the finishing temperature in hot-rolling process over 960°C,
the crystal grains of the hot-rolled sheet are coarsened to deteriorate deep drawability
after cold rolling and annealing. On the other hand, with a temperature of less than
650°C, deformation resistance is increased to increase the hot-rolling load, causing
difficulties in rolling.
[0100] Preferably, cooling is started immediately after the end of final rolling in hot-rolling
process to prevent normal grain growth and suppress AlN precipitation in the cooling
step.
[0101] Although the cooling condition is not limited, the starting time of the cooling step
is preferably within 1.5 seconds, more preferably 1.0 second, and most preferably
0.5 second, after the end of finish rolling. This is because when cooling is performed
immediately after the end of rolling, a large amount of ferrite nuclei is produced
due to an increase in the degree of over cooling with accumulated strain to promote
ferrite transformation and suppress the diffusion of dissolved N in the γ phase into
the ferrite grains, thereby increasing the amount of dissolved N present in the ferrite
grain boundaries.
[0102] The cooling rate is preferably 10°C/s or more in order to ensure dissolved N. Particularly,
when the finishing temperature of hot-rolling is the Ar
3 transformation point or more, the cooling rate is preferably 50°C/s or more in order
to ensure dissolved N.
[0103] Then, the hot-rolled sheet is coiled. In order to coarsen a carbide, the coiling
temperature is advantageously as high as possible. However, with a coiling temperature
of over 800°C, the scale formed on the surface of the hot-rolled sheet is thickened
to increase the load of the work of removing the scale, and progress the formation
of a nitride, causing a change in the amount of dissolved N in the coil length direction.
On the other hand, with a coiling temperature of less than 400°C, the coiling work
becomes difficult. Therefore, the coiling temperature of the hot-rolled sheet must
be in the range of 400 to 800°C.
[0104] Then, the hot-rolled sheet is cold-rolled, but the rolling reduction ratio of cold
rolling must be 60 to 95%. This is because with a rolling reduction ratio of cold
rolling of less than 60%, a high r value cannot be expected, while with a rolling
reduction ratio of over 95%, the r value is decreased.
[0105] The cold-rolled sheet subjected to cold rolling is then recrystallized and annealed.
Although the annealing method may be either continuous annealing or batch annealing,
continuous annealing is advantageous. The continuous annealing may be performed either
in a normal continuous annealing line or in a continuous hot-dip galvanization line.
[0106] The preferable annealing conditions include 650°C or more for 5 seconds or more.
This is because with an annealing temperature of less than 650°C, and an annealing
condition of less than 5 seconds, recrystallization is not completed to decrease deep
drawability. In order to improve deep drawability, annealing is preferably performed
in the ferrite single phase region at 800°C or more for 5 seconds or more.
[0107] Annealing in the high-temperature α+γ two-phase region partially produces α → γ transformation
to improve the r value due to the development of the {111} aggregation structure.
However, when α → γ transformation completely proceeds, the aggregation structure
is made random to decrease the r value, thereby deteriorating deep drawability.
[0108] The upper limit of the annealing temperature is preferably 900°C. This is because
with an annealing temperature of over 900°C, redissolution of a carbide proceeds to
excessively increase the amount of dissolved C, thereby deteriorating the natural
aging property. When α → γ transformation occurs, the aggregation structure is made
random to decrease the r value, deteriorating deep drawability.
[0109] Furthermore, in the heating-up step in recrystallization annealing, slow heating
is performed in the temperature region from 500°C to the recrystallization temperature
to sufficiently precipitate AlN, and the like, thereby effectively decreasing the
crystal grain diameter of the steel sheet.
[0110] The temperature region in which controlled heating must be performed is 500°C, at
which AlN or the like starts to precipitate, to the recrystallization temperature
[0111] The heating rate is preferably in the range of 1 to 20°C/s because with a heating
rate of over 20°C/s, the sufficient amount of precipitates cannot be obtained, while
with a heating rate of less than 1°C/s, precipitates are coarsened to weaken the effect
of suppressing grain growth.
[0112] After the recrystallization annealing, temper rolling of 10% or less may be performed
for correcting the shape and controlling surface roughness.
[0113] The cooling rate after soaking in recrystallization annealing is preferably 10 to
50°C/s. This is because with a cooling rate of 10°C/s or less, grains are grown during
cooling to coarsen the crystal grains, thereby deteriorating the strain aging property
and natural aging property. While with a cooling rate of 50°C/s or more, dissolved
N does not sufficiently diffuse into the grain boundaries, deteriorating the natural
aging property. The cooling rate is preferably 10 to 30°C/s.
[0114] After the recrystallization annealing, hot-dip galvanization and alloying by heating
are performed to form a galvannealed steel sheet as occasion demands.
[0115] The hot-dip galvanization and alloying are not limited, and may be performed according
to a conventional known method.
[0116] With a steel sheet subjected to surface treatment generally used for steel thin sheets,
such as a steel sheet (a dull-finished steel sheet, a bright-finished steel sheet,
or a steel sheet having a specified roughness pattern formed on the surface thereof),
which is produced by temper-rolling the alloyed hot-dip galvanized steel sheet, for
improving processability and the appearance after processing, a steel sheet having
an oil film layer of antirust oil or lubricating oil formed on the surface thereof,
or the like, the effect of the present invention can be sufficiently exhibited in
the composition range of the prevent invention.
[0117] Therefore, a cold-rolled steel sheet and a galvannealed steel sheet can be obtained,
which have excellent deep drawability and excellent strain age hardenability, that
tensile strength increased by press forming and heat treatment.
[0118] A description will now be made of the reasons for limiting the components of a steel
sheet in the above ranges according to a third embodiment of the present invention.
C: less than 0.01 mass %
[0119] From the viewpoint of excellent deep drawability and press moldability, C is advantageously
as small as possible. Also, redissolution of NbC proceeds in the annealing step after
cold rolling to increase the amount of dissolved C in crystal grains, thereby easily
causing deterioration in natural aging resistance. Therefore, the C amount is preferably
suppressed to less than 0.01 mass %, more preferably 0.0050 mass % or less, and most
preferably 0.0030 mass % or less. In order to ensure strength and prevent coarsening
of crystal grains, the C content is preferably 0.0005% or more.
Si: 0.005 to 1.0 mass %
[0120] Si is a useful component for suppressing a decrease in elongation, and improving
strength. However, with a Si content of less than 0.005 mass %, the effect of addition
of Si is insufficient, while with a Si content of over 1.0 mass %, the deterioration
of surface properties induce a decrease in elongation. Therefore, the Si content is
limited to the range of 0.005 to 1.0 mass %, and preferably the range of 0.01 to 0.75
mass %.
Mn: 0.01 to 1.5 mass %
[0121] Mn not only is useful as a strengthening component for steel, but also has the function
to suppress embrittlement with S due to the formation of MnS.. However, with a Mn
content of less than 0.01 mass %, the effect of addition of Mn is insufficient, while
with a Mn content of over 1.5 mass %, the deterioration of surface properties induce
a decrease in elongation. Therefore, the Mn content is limited to the range of 0.01
to 1.5 mass %, and preferably the range of 0.10 to 0.75 mass %.
P: 0.01 mass % or less
[0122] P is a solid solution strengthening element which effectively contributes to strengthening
of steel. However, with a P content of over 0.01 mass %, deep drawability deteriorates
due to the formation of phosphide such as (FeNb)
xP or the like. Therefore, P is limited to 0.10 mass % or less.
S: 0.01 mass % or less
[0123] With a high S content, the amount of inclusions is increased to deteriorate ductility.
Therefore, contamination with S is preferably prevented as much as possible, but an
S content up to 0.01 mass % is allowable.
Al: 0.005 to 0.030 mass %
[0124] Al is added as a element for deoxidization for improving the yield of the elements
forming carbonitride. However, with an Al content of less than 0.005 mass %, the effect
is insufficient, while with an Al content of over 0.030 mass %, the amount of N to
be added to steel is increased to easily cause slab defects during steel making. Therefore,
Al is contained in the range of 0.005 to 0.030 mass %.
N: 0.005 to 0.040 mass %
[0125] In the present invention, N is an important element which plays the role of imparting
strain age hardenability to a steel sheet. However, with an N content of less than
0.005 mass %, a sufficient strain age hardenability cannot be obtained, while with
an N content of as high as over 0.040 mass %, press moldability deteriorates. Therefore,
N is contained in the range of 0.005 to 0.040 mass %, and preferably in the range
of 0.008 to 0.015 mass %.
B: 0.0001 to 0.003 mass %
[0126] B is added in a combination with Nb to exhibit the function to effectively make fine
the micro structure of the hot-rolled steel and the cold-rolled steel, annealed for
recrystallization, and to improve cold-work embrittlement resistance. However, with
a B content of less than 0.0001 mass %, the sufficient effect of making fine the structures
cannot be obtained, while with a B content of over 0.003 mass %, the amount of BN
precipitate is increased, and dissolution in the slab heating step is hindered. Therefore,
B is contained in the range of 0.0001 to 0.003 mass %, preferably in the range of
0.0001 to 0.0015 mass %, and more preferably in the range of 0.0007 to 0.0012 mass
%.
Nb: 0.005 to 0.050%, Ti: 0.005 to 0.070%, V: 0.005 to 0.10%
[0127] Nb, Ti and V are added in a combination with B to contribute to refinement of the
the micro structure of the hot-rolled steel and the cold-rolled steel, annealed for
recrystallization, and have the function to precipitate dissolved C as NbC, Tic and
VC, respectively. Therefore, these elements are added together with B according to
demand, but less than 0.005% each of the elements does not sufficiently exhibit the
function. On the other hand, over 0.050% of Nb, over 0.070% of Ti and over 0.10% of
V cause deterioration in ductility. Therefore, Nb, Ti and V are added in the ranges
of 0.005 to 0.050%, 0.005 to 0.070%, and 0.005 to 0.10%, respectively.
[0128] Furthermore, as described above, Nb has the function to fix dissolved C as NbC, and
forms a nitride NbN. Similarly, Al and B form AlN and NB, respectively. Therefore,
in order to ensure the sufficient amount of dissolved N and sufficiently decrease
the amount of dissolved C, it is important to satisfy the following relations (1)
and (2) :
N/Al or N/(Al+Nb+Ti+V+B): 0.30 or more
[0129] Al forms AlN to decrease the amount of dissolved N. In order to ensure an appropriate
amount of dissolved N, N/Al must be 0.30 or more. When Al is added in a combination
with Nb, Ti, V or B, these elements also respectively form NbN, TiN, VN and BN to
decrease the amount of dissolved N. Therefore, in order to ensure an appropriate amount
of dissolved N, (Al+Nb+Ti+V+B) must be 0.30 or more.
Dissolved N: 0.0010% or more
[0130] In order to increase the strain age hardenability of the steel sheet, the content
of dissolved N must be 0.0010% or more.
[0131] The amount of dissolved N is determined by subtracting the amount of precipitated
N from the total N amount of steel. As a result of research of comparison between
various analysis methods, the inventors found that an electrolytic extraction method
using a constant-potential electrolytic method is effective as the method of analyzing
the amount of precipitated N. As a method of dissolving ferrite used for extraction
analysis, an acid digestion method, a halogen method, or an electrolysis method can
be used. Of these methods, the electrolysis method can stably dissolve only ferrite
without decomposing very unstable precipitates such as a carbide, a nitride, etc.
As the electrolyte, an acetyl-acetone system is used for electrolysis at a constant
potential. In the present invention, the results of measurement of the amount of precipitated
N by constant-potential electrolysis showed best correspondence with the actual strength
of parts.
[0132] Therefore, in the present invention, the residue after extraction by constant-potential
electrolysis is chemically analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated N.
[0133] In order to obtain higher BH and ΔTS, the amount of dissolved N is preferably 0.0015%
or more, more preferably 0.0020% or more, and most preferably 0.0030% or more.
[0134] The cold-rolled steel sheet of the present invention is a cold-rolled deep drawing
steel sheet having excellent strain age hardenability and the above-described composition,
wherein TS x r value ≥ 750 MPa.
[0135] A steel sheet having a Ts x r value of less than 750 MPa cannot be widely applied
to members comprising structural member components. In order to extend the application
range, the TS x r value is preferably 850 MPa or more.
[0136] Conventional coating and baking conditions include 170°C for 20 min as standards.
When a strain of 5% is applied to the steel sheet of the present invention, which
contains a large amount of dissolved N, hardening can be achieved even by slow (low
temperature) processing. In other words, the range of aging conditions can be widened.
In order to attain a sufficient amount of hardening, generally, retention at a higher
temperature for a longer time is advantageous as long as softening does not occurs
by over aging.
[0137] Specifically, in the steel sheet of the present invention, the lower limit of the
heating temperature at which hardening significantly takes place after pre-deformation
is about 100°C. On the other hand, with the heating temperature of over 300°C, hardening
peaks, thereby causing the tendency to soften and significantly causing thermal strain
and temper color. With the retention time of about 30 seconds or more, hardening can
be sufficiently achieved at a heating temperature of about 200°C. In order to obtain
more stable hardening, the retention time is preferably 60 seconds or more. However,
retention for over 20 mines is practically disadvantageous because further hardening
cannot be expected, and the production efficiency significantly deteriorates.
[0138] Therefore, in the present invention, the conventional coating and baking conditions,
i.e., the heating temperature of 170°C and the retention time of 20 minutes, are set
as the aging conditions. With the steel sheet of the present invention, hardening
can be stably achieved even under the aging conditions of a low heating temperature
and a short retention time, which fail to achieve sufficient hardening in a conventional
bake-hardening steel sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and other methods
such as induction heating, heating with a nonoxidation flame, a laser, plasma, or
the like, etc. can be preferably used. Alternatively, only a portion in which strength
is desired to be increased may be selectively heated.
[0139] The strength of an automobile part must be sufficient to resist an external complicated
stress load, and thus not only strength in a low strain region but also strength in
a high strain region are important for a raw material steel sheet. In consideration
of this point, in the steel sheet of the present invention used as a raw material
for automobile parts, BH is 80 MPa or more, and ΔTS is 40 MPa or more. More preferably,
BH is 100 MPa or more, and ΔTS is 50 MPa or more. In order to further increase BH
and ΔTS, the heating temperature in aging may be set to a higher temperature, and/or
the retention time may be set to a longer time.
[0140] The steel sheet of the present invention has the advantage that even when the steel
sheet not molded is allowed to stand at room temperature for about one year, natural
aging deterioration does not occur, unlike a conventional steel sheet.
[0141] In the present invention, the cold-rolled steel sheet may be coated by hot-dip galvanization
or alloying hot-dip galvanization without any problem, and TS, BH and ΔTS are equivalent
to those before plating. Besides hot-dip galvanization and alloying hot-dip galvanization,
electro-galvanization, electro-tinning, electric chromium plating, electro-nickeling,
and the like may be preferably used.
[0142] The production conditions according to a fourth embodiment of the present invention
will be described.
[0143] First, steel having the composition, by mass %, comprising less than 0.01% of C,
0.0050 to 0.04% of N, 0.005 to 0.03% of Al, 0.005 to 1.0% of Si, 0.01 to 1.5% of Mn,
0.1% or less of P, and 0.01% or less of S, or further comprising 0.0001 to 0.003%
of B, and at least one of 0.005 to 0.050% of Nb, 0.005 to 0.070% of Ti, and 0.005
to 0.10% of V, wherein N/(Al+Nb+Ti+V+B) is 0.30 or more, is melted by a conventional
melting method such as a converter or the like, and then solidified by an ingot making
method or a continuous casting method to form a steel raw material.
[0144] The steel raw material is heated and soaked, and then hot-rolled to form a hot-rolled
sheet. With an excessively low heating temperature (SRT), the effect of improving
processability is saturated, and the rolling load in hot rolling is increased to cause
a trouble in rolling, and insufficient homogeneity of dissolved N. Therefore, SRT
is preferably 950°C or more. In order to improve deep drawability, it is advantageous
to fix dissolved C and precipitate it as a carbide. Therefore, SRT is preferably 1300°C
or less. In order to further improve processability, SRT is preferably 1150°C or less.
[0145] When the total rolling reduction ratio of hot rolling starting from rough rolling
to finish rolling is less than 80%, the crystal grains of the hot-rolled sheet are
not sufficiently made fine. Therefore, the total rolling reduction ratio is preferably
80% or more.
[0146] With the rough rolling temperature of over 1000°C, γ→α transformed grains are coarsened
to decrease the r value, while with the rough rolling temperature of less than Ar
3 transformation point, α grains are recrystallized and coarsened or grown to decrease
the r value. Therefore, rough rolling is preferably performed in the temperature region
of the Ar
3 transformation point to 1000°C.
[0147] On the other hand, when finish rolling is completed in the temperature region of
over the Ar
3 transformation point, the aggregation structure is made random by γ→α transformation
to fail to obtain excellent deep drawability. When finish rolling is completed in
the temperature region of less than the Ar
3 transformation point, a further improvement in deep drawability cannot be expected,
but the rolling load is increased. Therefore, finish rolling is preferably performed
in the temperature region of 600°C to the Ar
3 transformation point.
[0148] If lubrication rolling is not performed in finish rolling, an additional shear force
acts on the surface of the steel sheet due to frictional force between a roll and
the steel sheet to preferentially form {110} orientation undesirable for deep drawing
in the surface of the steel sheet, deteriorating deep drawability. Therefore, finish
rolling is preferably performed under lubrication.
[0149] Then, the hot-rolled sheet is coiled. The processed material after the coiling step
is referred to as a "coil". The higher coiling temperature (CT) of the hot-rolled
sheet is advantageous for coarsening of the carbide. However, with the coiling temperature
of over 800°C, the scale formed on the surface of the hot-rolled sheet is thickened
to increase the load of a scale removing work, and progress the formation of a nitride,
thereby causing a variation in amount of dissolved N in the longitudinal direction
of the coil. On the other hand, with the coiling temperature of less than 400°C, the
coiling work is difficult. Therefore, CT is preferably 400 to 800°C.
[0150] Then, the thus-obtained hot-rolled sheet is recrystallized and annealed by continuous
annealing or batch annealing. The annealing (hot-rolled sheet annealing) is carried
out for recrystallizing the rolled aggregation structure formed by hot rolling in
the α-phase region in finish rolling to obtain a recrystallized aggregation structure.
[0151] Then, the hot-rolled sheet is cold-rolled to form a cold-rolled sheet. When the rolling
reduction ratio of cold rolling is less than 60%, a high r value cannot be expected.
On the other hand, while a rolling reduction ratio of over 95%, the r value is decreased.
Therefore, the rolling reduction ratio is preferably 60 to 95%.
[0152] Next, the cold-rolled sheet is recrystallized and annealed. The annealing is preferably
carried out in either a continuous annealing line or a continuous hot-dip galvanization
line. The preferable annealing conditions include the annealing temperature of 650°C
or more and the retention time of 5 seconds or more. When either of the annealing
temperature of 650°C or more and the retention time of 5 seconds or more is not satisfied,
recrystallization is not completed to deteriorate deep drawability. In order to obtain
excellent deep drawability, the annealing temperature of 800°C or more and the retention
time of 5 seconds or more are preferred. However, with the annealing temperature of
over 900°C, redissolution of the carbide proceeds to excessively increase the amount
of dissolved C, thereby deteriorating the natural aging property (decreasing elongation
by natural aging). Furthermore, when γ→α transformation occurs, the aggregation structure
becomes random to decrease the r value, deteriorating deep drawability. Therefore,
the annealing temperature is preferably 900°C or less.
[0153] The cold-rolled annealed sheet obtained by recrystallizing and annealing the cold-rolled
steel sheet is further coated by hot-dip galvanization or alloyed. In this case, in
plating, the cooling rate during the time between the completion of recrystallization
annealing and the start of plating is 5°C/s or more, and the sheet temperature in
hot-dip galvanization is preferably 400 to 600°C. In alloying, the processing temperature
is preferably 400 to 600°C, and the processing time is preferably 5 to 40 seconds.
[0154] The cold-rolled steel sheet after recrystallization annealing or the hot-dip galvanized
steel sheet may be temper-rolled for correcting the shape and controlling surface
roughness. The reduction ratio of temper rolling is preferably 10% or less. This is
because with a rolling reduction ratio of over 10%, the r value is decreased.
[0155] Description will now be made of the reasons for limiting the composition of a high-tensile-strength
cold-rolled steel sheet according to a fifth embodiment of the present invention.
C: 0.0015 to 0.025%
[0156] In the present invention, in order to control the structure to a homogeneous fine
structure, and ensure a sufficient amount of an acicular ferrite phase, the C content
must be 0.0015% or more. With a C content of over 0.025%, the ratio of the carbide
in the steel sheet is excessively increased to significantly deteriorate ductility,
the r value and moldability. Therefore, the C content is limited in the range of 0.0015
to 0.025%. From the viewpoint of improvement in moldability, the C content is preferably
0.020% or less, more preferably 0.010% or less. Particularly, from the viewpoint of
stabilization of the BH amount and material properties, the C content preferably exceeds
(12/93) Nb (%) (wherein Nb represents the Nb content (%)).
Si: 1.0% or less
[0157] Si is a useful component capable of increasing the strength of the steel sheet without
significantly deteriorating ductility of steel. Particularly, when high strength is
required, the Si content is preferably 0.10% or more. On the other hand, Si is an
element which greatly changes the transformation point during hot rolling to cause
difficulties in ensuring quality and the shape, or adversely affects surface properties,
chemical conversion properties, and the like, particularly the beauty of the surface
of the steel sheet, and adversely affects plating properties. In the present invention,
therefore, the Si content is limited to 1.0% or less. However, the above-described
adverse effects can be kept down as long as Si is 1.0% or less. Particularly, in applications
required for the steel sheet to have surface beauty, Si is preferably 0.5% or less.
Mn: 2.0% or less
[0158] Mn is an element effective to prevent hot cracking with S, and Mn is preferably added
according to the amount of S contained. Mn also has the great effect of making fine
crystal grains, and is preferably added for improving material properties. In order
to stably fix S, the Mn content is preferably 0.1% or more. Mn is also an element
for increasing the strength of the steel sheet, and is preferably added in an amount
of 0.5% or more when higher strength is required. The Mn content is more preferably
0.8% or more.
[0159] With the Mn content increased to this level, there is the advantage that variations
in the mechanical properties of the steel sheet with respect to variations in the
hot-rolling conditions, particularly strain age hardenability, are significantly improved.
However, with the excessively high Mn content of over 2.0%, the deformation resistance
at elevated temperatures tends to increase to deteriorate weldability and weld moldability.
The detailed mechanism of this is not known. Furthermore, the formation of ferrite
is significantly suppressed, and the r value is significantly decreased. Therefore,
the Mn content is limited to 2.0% or less. In applications required to have good corrosion
resistance and moldability, the Mn content is preferably 1.5% or less.
P: 0.1% or less
[0160] P is a useful element as a solid solution strengthening element for steel, and is
preferably added in an amount of 0.002% or more from the viewpoint of an increase
in strength. Particularly, when high strength is required, the P content is preferably
0.02% or more. On the other hand, when P is excessively added, steel is embrittled,
and stretch-flanging properties of the steel sheet deteriorate. Also, P is liable
to strongly segregate in steel, thereby causing embrittlement of a weld. Therefore,
P is limited to 0.1% or less. In applications in which elongated flange processability
and weld toughness are considered as important, P is preferably 0.08% or less, more
preferably 0.06% or less.
S: 0.02% or less
[0161] S is present as an inclusion in the steel sheet, decreases ductility of the steel
sheet, and causes deterioration in corrosion resistance. Therefore, the Si content
is as low as possible, and in the present invention, the Si content is limited to
0.02% or less. Particularly, in applications required to have good processability,
S is preferably 0.015% or less. Particularly, in applications required to have excellent
stretch-flanging properties, S is preferably 0.010% or less. Although the detailed
mechanism is not known, in order to stably maintain the strain age hardenability of
the steel sheet in a high level, it is effective to decrease the S content to 0.008%
or less.
Al: 0.02% or less
[0162] Al is an element functioning as a element for deoxidation for improving cleanliness
of steel, and making fine the structure of the steel sheet. In the present invention,
the Al content is preferably 0.001% or more. In the present invention, dissolved N
is used as a strengthening element, but aluminum killed steel containing Al in a suitable
range has mechanical properties superior to those of conventional rimmed steel not
containing Al. On the other hand, with an excessively high Al content, the surface
properties of the steel sheet deteriorate, and the amount of dissolved N is significantly
decreased to cause difficulties in obtaining a large amount of strain age hardening,
which is the main object of the present invention. Therefore, in the present invention,
Al is limited to 0.02% or less. From the viewpoint of stability of material quality,
Al is more preferably 0.001 to 0.015%. Although a decrease in the Al content possibly
causes coarsening of crystal grains, in the present invention, the amounts of other
alloy elements are appropriately determined, and the annealing conditions are set
in appropriately ranges, thereby effectively preventing coarsening.
N: 0.0050 to 0.0250 mass %
[0163] N is an element for increasing the strength of the steel sheet by solid solution
strengthening and strain age hardening, and in the present invention, N is the most
important element. In the present invention, an appropriate amount of N is contained,
the Al content is controlled to the appropriate value, and production conditions such
as the hot-rolling conditions, and the annealing conditions are controlled to ensure
necessary and sufficient dissolved N in a cold-rolled product or a coated product.
This exhibits the. sufficient effect of increasing strength (yield stress and tensile
strength) by solid solution strengthening and strain age hardening, to stably obtain
the target values of the mechanical properties of the steel sheet of the present invention,
such as a tensile strength of 340 MPa or more, an amount (BH amount) of bake-hardening
of 80 MPa or more, and an increase is tensile strength ΔTS of 40 MPa or more after
strain aging. Since N also has the function to decrease the. transformation point,
N is effective for rolling of a thin material for which rolling at a temperature greatly
over the transformation point is undesirable.
[0164] With a N content of less than 0.0050%, the effect of increasing strength is less
stably exhibited, while with a N content of over 0.0250%, the rate of occurrence of
internal defects in the steel sheet is increased, and slab cracking frequently occurs
during continuous casting. Therefore, N is limited to the range of 0.0050 to 0.0250%.
From the viewpoint of improvement in stability of material properties and yield over
the entire production process, N is preferably in the range of 0.0070 to 0.0200%,
and more preferably in the range of 0.0100 to 0.0170%. With the N amount.in the range
of the present invention, there is no adverse effect on weldability, and the like.
Dissolved N: 0.0010% or more
[0165] In order to ensure sufficient strength of a cold-rolled product, and effectively
exhibit strain age hardening with N, it is necessary that the content of dissolved
N (solid solution N) in the steel sheet is at least 0.0010% or more.
[0166] The amount of dissolved N is determined by subtracting the amount of precipitated
N from the total N amount of steel. As a result of comparison research of various
methods, the inventors found that electrolytic extraction analysis using constant-potential
electrolysis is effective as the method of analyzing the amount of precipitated N.
As the method of dissolving ferrite used for extraction analysis, an acid digestion
method, a halogen method, or an electrolysis method can be used. Of these methods,
the electrolysis method can stably dissolve only ferrite without decomposing very
unstable precipitates such as a carbide, a nitride, etc. As the electrolyte, an acetyl-acetone
system is used for electrolyzing at a constant potential. In the present invention,
the results of measurement of the amount of precipitated N by constant-potential electrolysis
showed best correspondence with changes in actual material properties.
[0167] Therefore, in the present invention, the residue after extraction by constant-potential
electrolysis is chemically analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated N.
[0168] In order to obtain higher BH and ΔTS, the amount of dissolved N is preferably 0.0020%
or more. In order to obtain further high values, the amount of dissolved N is preferably
0.0030% or more. Although the upper limit of the amount of dissolved N is not limited,
the mechanical properties less deteriorate even when the all amount of N remains.
N/Al (the content ratio of N to Al): 0.3 or more
[0169] In order to cause 0.0010% or more of dissolved N to stably remain in a product state,
the amount of Al which is an element for strongly binding to N, must be limited. As
a result of research of steel sheets in which the combination of the N content (0.0050
to 0.0250%) and the Al content (0.02% or less) were widely changed in the composition
range of the present invention, it was found that with N/Al of 0.3 or more, the amount
of dissolved N of a cold-rolled product or coated product can be stably set to 0.0010%
or more. Therefore, N/Al is limited to 0.3 or more. In order to stably increase the
strain age hardenability, N/Al is preferably 0.6 or more, and more preferably 0.8
or more.
Nb: 0.002 to 0.050%
[0170] Nb effectively functions to form an acicular ferrite phase in combination with B.
In the present invention, the Nb content must be 0.002% or more. On the other hand,
a Nb content of over 0.050%, the effect is saturated, and deformation resistance at
elevated temperatures is significantly increased to cause difficulties in hot rolling.
Therefore, Nb is limited in the range of 0.002 to 0.050%, and preferably 0.005 to
0.040%.
B: 0.0001 to 0.0050%
[0171] B effectively functions to form an acicular ferrite phase in combination with Nb.
In the present invention, the B content must be 0.001% or more. On the other hand,
a B content of over 0.0050%, the amount of dissolved N contributing strain age hardenability
is decreased. Therefore, B is limited in the range of 0.0001 to 0.0050%, and preferably
0.0003 to 0.0030%.
[0172] In the present invention, the above composition preferably further contains at least
one of the following groups a to c:
Group a: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group b: one or both of Ti and V in a total of 0.1% or less; and
Group c: one or both of Ca and REM in a total of 0.0010 to 0.010%.
[0173] Element of group a: Cu, Ni, Cr and Mo are all contribute to an increase in strength
of the steel sheet, and can be contained singly or in a combination according to demand.
The effect is recognized by containing 0.01% or more each of Cu, Ni, Cr and Mo. However,
with an excessively high content, deformation resistance at elevated temperatures
in hot rolling is increased, or chemical conversion properties and surface treatment
properties in a wide sense deteriorate, and a welded portion is hardened to deteriorate
weld moldability. Therefore, Cu, Ni, Cr, and Mo are preferably contained singly at
1.0% or less, 1.0% or less, 0.5% or less, and 0.2% or less, respectively, and preferably
contained in a combination at a total of 1.0% or less.
[0174] Element of group b: Both Ti and V are elements contributing refinement and homegenization
of crystal grains, and may be added singly or in a combination according to demand.
The effect can be recognized by containing 0.005% or more each of Ti and V. However,
with an excessively high content, deformation resistance at elevated temperatures
in hot rolling is increased, or chemical conversion properties and surface treatment
properties in a wide sense deteriorate. Furthermore, there is the adverse effect of
decreasing the amount of dissolved N. Therefore, Ti and V are preferably contained
singly at 1.0% or less and 1.0% or less, respectively, and preferably contained in
a combination at a total of 0.1% or less.
[0175] Elements of group c: Both Ca and REM are elements useful for controlling the form
of inclusions. Particularly, when the stretch flanging property is required, these
elements are preferably added singly or in a combination. When the total of the elements
of group d is less than 0.0010%, the effect of controlling the form of inclusions
is insufficient, while when the total exceeds 0.010%, surface defects significantly
occur. Therefore, the total of the elements of group d is preferably limited to the
range of 0.0010 to 0.010%. This permits improvement in the stretch flanging property
without causing surface defects.
[0176] The structure of the steel sheet of the present invention is described below.
[0177] The steel sheet of the present invention has the structure composed of an acicular
ferrite phase at an area ratio of 5% or more and a ferrite phase having an average
crystal grain diameter of 20 µm or less.
[0178] The area ratio of the acicular ferrite phase: 5% or more
[0179] The cold-rolled steel sheet of the present invention contains the acicular ferrite
phase at an area ratio of 5% or more. The presence of the acicular ferrite phase at
5% or more permits the achievement of good ductility and a larger amount of strain
age hardening. Although the detailed mechanism is not known, it is thought that strain
is effectively accumulated in the steel sheet during pre-strain processing before
aging by the presence of the acicular ferrite phase. Furthermore, the presence of
the acicular ferrite phase improves natural aging deterioration at room temperature
to be effective to obtain natural non-aging properties. In order to obtain a good
balance of strength and ductility, and higher strength, the area ratio of the acicular
ferrite phase is preferably 10% or more. The presence of a large amount of acicular
ferrite phase of over 20% has the problem of deteriorating the r value. Therefore,
the area ratio of the acicular ferrite phase is 5% or more, and preferably 10% to
20%.
[0180] In the present invention, "the acicular ferrite phase" is a low-temperature transformation
phase peculiar to ultra low carbon steel having the composition of the present invention,
in which no carbide is contained therein. This phase can be clearly discriminated
from normal polygonal ferrite by observation on an optical microscope, and is harder
than the polygonal ferrite phase because of the high internal dislocation density.
[0181] According to observation on an optical microscope, the acicular ferrite phase has
a distribution of any one of (1) crystal grains having irregularly angular boundaries,
(2) crystal grains present along crystal grains of precipitates or the like, and (3)
crystal grains or crystal grain groups (many sub-boundaries are observed in relatively
second phase grains) having a scratch-like pattern, singly or in a combination. This
acicular ferrite can be clearly distinguished from general polygonal ferrite. The
color tone of the corroded insides of the grains is different from martensite and
bainite, and substantially the same as ordinary polygonal ferrite, and thus acicular
ferrite can be clearly distinguished from martensite and bainite. According to observation
on a transmission electron microscope, the acicular ferrite phase has a very high
density of dislocation in the vicinities of the crystal grains and/or in the grains,
and particularly the above form (3) comprises a layer having a very high dislocation
density and a layer having a relatively low dislocation density.
[0182] The cold-rolled steel sheet of the present invention is directed to use as an automobile
steel sheet required to have high moldability, and comprises a ferrite phase other
than the acicular ferrite phase in order to ensure ductility. When the area ratio
of the ferrite phase is less than 80%, it is difficult to ensure ductility and a high
r value which are necessary for the automobile steel sheet required to have processability.
When further improved ductility is required, the area ratio of the ferrite phase is
80% or more, and preferably 85% or more. In the present invention, ferrite means so-called
polygonal ferrite in which no strain remains.
Average crystal grain diameter of ferrite phase: 20 µm or less
[0183] In the present invention, the value used as the average crystal grain diameter is
a higher one of the value calculated from a photograph of a sectional structure by
a quadrature method defined by ASTM, and the nominal value determined by an intercept
method defined by ASTM (refer to, for example, Umemoto et al.: Heat Treatment, 24
(1984), p334).
[0184] The cold-rolled steel sheet of the present invention maintains a predetermined amount
of dissolved N in the product step. However, as a result of experiment and research
conducted by the inventors, it was found that variations in strain age hardenability
occur in steel sheets containing the same amount of dissolved N, and one of the main
causes of the variations is a crystal grain diameter. In the structure of the present
invention, in order to stably obtain a high BH amount and ΔTS, the average crystal
grain diameter is at least 20 µm or less, and preferably 15 µm or less. Although the
detailed mechanism is not known, this is supposed to be related to the segregation
and precipitation of alloy elements in the crystal grain boundaries, and the influences
of processing and heat history on the segregation and precipitation.
[0185] Therefore, in order to achieve stability of strain age hardenability, the average
crystal grain diameter of the ferrite phase is 20 µm or less, and preferably 15 µm
or less.
[0186] The cold-rolled steel sheet of the present invention, which has the above-described
composition and structure, has a tensile strength (TS) of about 340 MPa to 590 MPa,
a r value of as high as 1.2 or more, an excellent strain age hardenability. The steel
sheet having TS of less than 340 MPa cannot be widely applied to members each comprising
a structural component. In order to extend the application range, TS is preferably
400 MPa or more. With a r value of less than 1.2, the steel sheet cannot be applied
to a wide range of press-molded products. The preferred range of the r value is 1.3
or more.
[0187] The conventional coating and baking conditions include 170°C and 20 min as standards.
When a strain of 5% or more is applied to the steel sheet of the present invention
containing a large amount of dissolved N, hardening can be achieved even by aging
at low temperature. In other words, the aging conditions can be selected from a wide
range. Generally, in order to earn the amount of hardening, it is advantageous to
hold the steel sheet at a higher temperature for a longer time as long as softening
does not occur due to over aging.
[0188] Specifically, in the steel sheet of the present invention, the lower limit of the
heating temperature at which hardening significantly takes place after pre-deformation
is about 100°C. On the other hand, with the heating temperature of over 300°C, hardening
peaks, thereby causing the tendency to soften and significantly causing thermal strain
and temper color. With the retention time of about 30 seconds or more, hardening can
be sufficiently achieved at a heating temperature of about 200°C. In order to obtain
more stable hardening, the retention time is preferably 60 seconds or more. However,
retention for over 20 minutes is practically disadvantageous because further hardening
cannot be expected, and the production efficiency significantly deteriorates.
[0189] Therefore, in the present invention, the conventional coating and baking conditions,
i.e., the heating temperature of 170°C and the retention time. of 20 minutes, are
set as the aging conditions. With the steel sheet of the present invention, hardening
can be stably achieved even under the aging conditions of a low heating temperature
and a short retention time, which fail to achieve sufficient hardening in a conventional
bake-hardening steel sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and other methods
such as induction heating, heating with a nonoxidation flame, a laser, plasma, or
the like, etc. can be preferably used.
[0190] The strength of an automobile part must be sufficient to resist an external complicated
stress load, and thus not only strength in a low strain region but also strength in
a high strain region are important for a raw material steel sheet. In consideration
of this point, in the steel sheet of the present invention used as a raw material
for automobile parts, BH is 80 MPa or more (corresponding to strength in a relatively
low strain region), and ΔTS is 40 MPa or more (corresponding to strength in a relatively
high strain region). More preferably, BH is 100 MPa or more, and ΔTS is 50 MPa or
more. In order to further increase BH and ΔTS, the heating temperature in aging may
be set to a higher temperature, and/or the retention time may be set to a longer time.
[0191] The effect of the present invention is exhibited by a product having a relatively
large thickness. However, with a product having a thickness of over 3.2 mm, a sufficient
cooling rate necessary for the cold-rolled sheet annealing step cannot be ensured
to cause strain aging at the time of continuous annealing, thereby failing to obtain
the target strain age hardenability as a product. Therefore, the steel sheet of the
present invention preferably has a thickness of 3.2 mm or less.
[0192] In the present invention, the surface of the cold-rolled steel sheet may be coated
by hot-dip galvanization or alloying hot-dip galvanization without any problem. These
coated steel sheets also exhibit TS, BH and ΔTS which are equivalent to those before
plating. As the plating type, any one of electro-galvanization, hot-dip galvanization,
alloying hot-dip galvanization, electro-tinning, electric chromium plating, electro-nickeling,
and the like may be preferably used.
[0193] The method of producing a steel sheet according to a sixth embodiment of the present
invention is described.
[0194] The steel sheet of the present invention is basically produced by performing the
hot rolling step in which a steel slab having the above-described composition is heated,
and then roughly rolled to form a sheet bar, and the sheet bar is finish-rolled and
cooled to form a coiled hot-rolled sheet, the cold rolling step in which the hot-rolled
sheet is pickled and cold-rolled to form a cold-rolled sheet, and the cold-rolled
sheet annealing step in which the cold-rolled sheet is continuously annealed.
[0195] Although the slab used in the production method of the present invention is preferably
produced by a continuous casting method in order to prevent macro-segregation of components,
an ingot making method or a thin slab casing method may be used. Alternatively, a
conventional method comprising cooling the produced slab to room temperature and then
again heating the slab, or an energy-saving process of direct rolling, in which a
hot slab is charged into a heating furnace without cooling and then rolled, or the
slab is rolled directly immediately after being slightly kept warm, may be used without
a problem. Particularly, direct rolling is a useful technique for effectively ensuring
dissolved N.
[0196] Description will be now be made of the reasons for limiting the conditions of the
hot rolling step.
Slab heating temperature: 1000°C or more
[0197] The slab heating temperature is preferably 1000°C or more in order to ensure a necessary
and sufficient amount of dissolved N in an initial state and satisfy the target amount
of dissolved N in a product. Since a loss is increased by an increase in the oxide
weight, the heating temperature is preferably 1280°C or less.
[0198] The slab heated under the above condition is roughly rolled to form a sheet bar.
The condition of the rough rolling is not defined, and rough rolling may be performed
according to a normal method. However, in order to ensure the amount of dissolved
N, rough rolling is preferably performed within a as short time as possible. Then,
the sheet bar is finish-rolled to form a hot-rolled sheet.
[0199] In the present invention, the adjacent sheet bars. are preferably bonded together
during the time between rough rolling and finish rolling, and then continuously rolled.
As the bonding means, a pressure welding method, a laser welding method, an electron
beam welding method, or the like is preferably used.
[0200] By continuous rolling, so-called non-stationary portions at the front end and the
rear end of a coil (processed material) are removed to permit hot rolling over the
entire length and the entire width of the coil (processed material) under stable conditions.
This is very effective to improve the sectional shape and dimensions of not only the
hot-rolled steel sheet but also the cold-rolled steel sheet. When the steel sheet
is cooled on a hot run table after rolling, the shape of the steel sheet can be sufficiently
maintained because tension can be always applied.
[0201] By continuous rolling, the ends of the coil can be stably passed, and it is thus
possible to use lubricating rolling, which cannot be easily applied to ordinary single
rolling for each sheet bar because of the problem of continuous rolling processes
and biting property. Therefore, the rolling load can be decreased, and at the same
time, the surface pressure of the roll can be decreased, thereby increasing the life
time of the roll.
[0202] In the present invention, at the entrance of a finisher between rough rolling and
finish rolling, the temperature distributions of the sheet bar in the width direction
and the long direction thereof are preferably made uniform by using any one or both
of a sheet bar edge heater for heating the ends of the sheet bar in the width direction
and a sheet bar heater for heating the ends of the sheet bar in the long direction.
This can further decrease the variations in material properties of the steel sheet.
The sheet bar edge heater and the sheet bar heater are preferably of an induction
heating type.
[0203] As the procedure of use of the heaters, preferably, the difference in temperature
in the width direction is first corrected by the sheet bar edge heater. Although the
heating amount depends upon the steel composition, the temperature distribution in
the width direction at the finisher entrance is preferably set in the range of about
20°C or less. Next, the difference in temperature in the long direction is corrected
by the sheet bar heater. At this time, the heating amount is preferably set so that
the temperatures at the ends in the long direction are about 20°C higher than the
temperature at the center.
Finisher delivery temperature: 800°C or more
[0204] In order to obtain a homogeneous fine hot-rolled base sheet structure, the finisher
deliver temperature FDT is 800°C or more. With a FDT of lower than 800°C, the structure
of the steel sheet becomes inhomogeneous, and the processed structure partially remains
to leave heterogeneity of the structure after the cold-rolled sheet annealing step.
Therefore, the danger of producing various troubles in press forming is increased.
When a high coiling temperature is used for avoiding the processed micro structure
from remaining, coarse crystal grains are produced to cause the same troubles as described
above. With the high coiling temperature, the amount of dissolved N is significantly
decreased to cause difficulties in obtaining a target tensile strength of 340 MPa
or more. Therefore, the finisher deliver temperature FDT is 800°C or more. In order
to further improve the mechanical properties, the FDT is preferably 820°C or more.
In order to improve the r value, the FDT is preferably the Ac
3 transformation point or more. Although the upper limit of FDT is not limited, a scale
scar significantly occurs at excessively high FDT. The FDT is preferably up to about
1000°C.
Coiling temperature: 800°C or less
[0205] The strength of the steel sheet is liable to increase as the coiling temperature
CT decreases. In order to ensure the target tensile strength TS of 340 MPa or more,
the CT is preferably 800°C or less. With a CT of less than 200°C, the shape of the
steel sheet is readily disturbed to increase the. danger of causing troubles in a
practical operation, thereby deteriorating homogeneity of material properties. Therefore,
the CT is preferably 200°C or more. When the homogeneity of the material properties
is required, the CT is preferably 300°C or more, and more preferably 350°C or more.
[0206] In the present invention, in finish rolling, lubricating rolling may be performed
for decreasing the hot rolling load. The lubricating rolling has the effect of further
making homogeneous the shape and material properties of the hot-rolled sheet. During
the lubricating rolling, the frictional coefficient is preferably in the range of
0.25 to 0.10. By combining lubricating rolling and continuous rolling, the operation
of hot rolling is further stabilized.
[0207] The hot-rolled sheet subjected to the above-described hot rolling step is then pickled
and cold-rolled in the cold rolling step to form a cold-rolled sheet.
[0208] The pickling conditions may be the same as conventional known conditions, and are
not limited. When the scale of the hot-rolled sheet is extremely small, cold rolling
may be immediately after hot rolling without pickling.
[0209] The cold rolling conditions may be the same as conventional known conditions, and
are not limited. In order to ensure homogeneity of the structure, the reduction ratio
of cold rolling is preferably 60% or more. The reasons for limiting the conditions
of the cold-rolled sheet annealing step are described below.
[0210] The cold-rolled sheet is then subjected to the cold-rolled sheet annealing step comprising
continuous annealing and cooling.
Continuous annealing temperature: temperature in the ferrite-austenite two-phase
coexistence region
[0211] By annealing at a temperature in the ferrite-austenite two-phase coexistence region,
the acicular ferrite phase is formed. In addition, the (111) aggregation structure
is strongly developed in the ferrite phase to obtain a high r value. On the other
hand, with a high temperature where an austenite single phase is formed beyond the
ferrite-austenite two-phase coexistence region, the aggregation structure of the steel
sheet is made random by reverse transformation and transformation to decrease the
r value. Therefore, in the present invention, the annealing temperature of continuous
annealing is limited to the recrystallization temperature or more in the ferrite-austenite
two-phase coexistence region. From the viewpoint of stability of the r value, the
temperature is preferably set so that the fraction of austenite is 10% to 50%. With
the continuous annealing temperature of less than the recrystallization temperature,
ductility deteriorates, and the steel sheet can be applied as an automobile part only
to limited special applications. Therefore, the annealing temperature is preferably
the recrystallization temperature or more.
[0212] The retention time of continuous annealing is preferably as short as possible in
order to ensure the production efficiency, the fine structure and the amount of dissolved
N. From the viewpoint of stability of the operation, the retention time is preferably
10 seconds or more. Also, in order to ensure the fine structure and the amount of
dissolved N, the retention time is preferably 90 seconds or less. From the viewpoint
of stability of material properties, the retention time is preferably 20 seconds or
more.
Cooling after continuous annealing: cooling to the temperature region of 500°C
or less at a cooling rate of 10 to 300°C/s
[0213] Cooling after soaking by continuous annealing is important for making fine the structure,
forming the acicular ferrite phase, and ensuring the amount of dissolved N. In the
present invention, cooling is continuously carried out to the temperature region of
at least 500°C or less at a cooling rate of 10°C/s or more. With a cooling rate of
less than 10°C/s, a necessary amount of acicular ferrite phase, a homogeneous fine
structure and a sufficient amount of dissolved N cannot be obtained. On the other
hand, with a cooling rate of over 300°C/s, homogeneity in material properties of the
steel sheet in the width direction is insufficient. When the stop temperature of cooling
at a cooling rate of 10 to 300°C/s after continuous annealing exceeds 500°C, refinement
of the structure cannot be attained.
Temper rolling or lever processing: elongation of 0.5 to 10%
[0214] In the present invention, in order to correct the shape and control roughness, temper
rolling or leveler processing may be carried out subsequent to the cold rolling step.
When the total elongation of temper rolling or leveler processing is less than 0.5%,
the desired purpose of correcting the shape and controlling roughness cannot be achieved.
On the other hand, with a total elongation of over 10%, ductility deteriorates. In
order to ensure ductility, the elongation is preferably 5% or less. It is confirmed
that the processing system of temper rolling is different from that of leveler processing,
but the effects of both processes are substantially the same. Temper rolling and leveler
processing are effective even after plating.
[0215] Description will be now be made of the reasons for limiting the composition of a
high-tensile-strength steel sheet according to a seventh embodiment of the present
invention.
C: 0.025 to 0.15%
[0216] C is an element for increasing the strength of a steel sheet, and 0.025% or more
of C must be contained for controlling the structure to a homogeneous fine structure,
which is an important requirement of the present invention, and ensuring a sufficient
amount of a martensite phase. With a C content of over 0.15%, the ratio of the carbide
in the steel sheet is excessively increased to significantly deteriorate ductility
and moldability. With a C content of over 0.15%, there is the more important problem
of significantly deteriorating spot weldability and arc weldability. Therefore, the
C content is limited in the range of 0.025 to 0.15%. From the viewpoint of improvement
in moldability, the C content is preferably 0.08% or less. Particularly, when good
ductility is required, the C content is preferably 0.05% or less.
Si: 1.0% or less
[0217] Si is a useful component capable of increasing the strength of the steel sheet without
significantly deteriorating ductility of steel. The Si content is preferably 0.005%
or more, and more preferably 0.10% or more. On the other hand, Si is an element which
greatly changes the transformation point during hot rolling to cause difficulties
in ensuring quality and the shape, or adversely affects surface properties, chemical
conversion properties, and the like, particularly the beauty of the surface of the
steel sheet, and adversely affects plating properties. In the present invention, therefore,
the Si content is limited to 1.0% or less. However, the above-described adverse effects
can be kept down as long as Si is 1.0% or less. Particularly, in applications required
for the steel sheet to have a low level of strength and, particularly, surface. beauty,
Si is preferably 0.5% or less.
Mn: 2.0% or less
[0218] Mn is an element effective to prevent hot cracking with S, and Mn is preferably added
according to the amount of S contained. Mn also has the great effect of making fine
crystal grains, and is preferably added for improving material properties. Furthermore,
Mn is an element effective to stably form martensite during rapid cooling after continuous
annealing. In order to stably fix S, the Mn content is preferably 0.2% or more. Mn
is also an element for increasing the strength of the steel sheet, and is preferably
added in an amount of 1.2% or more when a strength TS of over 500 MPa is required.
The Mn content is more preferably 1.5% or more.
[0219] With the Mn content increased to this level, there is the advantage that variations
in the mechanical properties of the steel sheet with respect to variations in the
hot-rolling conditions, particularly strain age hardenability, are significantly improved.
However, with the excessively high Mn content of over 2.0%, a high r value, which
is an important requirement of the present invention, cannot be easily obtained, and
ductility significantly deteriorates.
Therefore, the Mn content is limited to 2.0% or less. In applications required to
have good corrosion resistance and moldability, the Mn content is preferably 1.7%
or less.
P: 0.08% or less
[0220] P is a useful element as a solid solution strengthening element for steel, and is
preferably added in an amount of 0.001% or more, and more preferably 0.015% or more,
from the viewpoint of an increase in strength. On the other hand, when P is excessively
added, steel is embrittled, and stretch-flanging properties of the steel sheet deteriorate.
Also, P is liable to strongly segregate in steel, thereby causing embrittlement of
a weld. Therefore, P is limited to 0.08% or less. In applications in which elongated
flange processability and weld toughness are considered as important, P is preferably
0.04% or less.
S: 0.02% or less
[0221] S is present as an inclusion in the steel sheet, decreases ductility of the steel
sheet, and causes deterioration in corrosion resistance. Therefore, the Si content
is as low as possible, and in the present invention, the S content is limited to 0.02%
or less. Particularly, in applications required to have good processability, S is
preferably 0.015% or less. Particularly, in applications required to have excellent
stretch-flanging properties, S is preferably 0.008% or less. Although the detailed
mechanism is not known, in order to stably maintain the strain age hardenability of
the steel sheet in a high level, it is effective to decrease the S content to 0.008%
or less.
Al: 0.02% or less
[0222] Al is an element functioning as a deoxidization for improving cleanliness of steel,
and making fine the structure of the steel sheet. In the present invention, the Al
content is preferably 0.001% or more. In the present invention, dissolved N is used
as a strengthening element, but aluminum killed steel containing Al in a suitable
range has mechanical properties superior to those of conventional rimmed steel not
containing Al. On the other hand, with an excessively high Al content, the surface
properties of the steel sheet deteriorate, and the amount of dissolved N is significantly
decreased to cause difficulties in obtaining a large amount of strain age hardening,
which is the main object of the present invention. Therefore, in the present invention,
Al is limited to 0.02% or less. From the viewpoint of stability of material properties,
Al is more preferably 0.001 to 0.015%. Although a decrease in the Al content possibly
causes coarsening of crystal grains, in the present invention, the amounts of other
alloy elements are appropriately determined to appropriately set the annealing conditions,
thereby effective preventing coarsening.
N: 0.0050 to 0.0250%
[0223] N is an element for increasing the strength of the steel sheet by solid solution
strengthening and strain age hardening, and in the present invention, N is the most
important element. In the present invention, an appropriate amount of N is contained,
the Al content is controlled to the appropriate value, and production conditions such
as the hot-rolling conditions, and the annealing conditions are controlled to ensure
necessary and sufficient dissolved N in a cold-rolled product or a coated product.
This exhibits the sufficient effect of increasing strength (yield stress and tensile
strength) by solid solution strengthening and strain age hardening, to stably obtain
the target values of the mechanical properties of the steel sheet of the present invention,
such as a tensile strength of 440 MPa or more, an amount (BH amount) of bake-hardening
of 80 MPa or more, and an increase is tensile strength ΔTS of 40 MPa or more after
strain aging. Since N also has the function to decrease the transformation point,
N is effective for rolling of a thin material for which rolling at a temperature greatly
over the transformation point is undesirable.
[0224] With a N content of less than 0.0050%, the effect of increasing strength is less
stably exhibited, while with a N content of over 0.0250%, the rate of occurrence of
internal defects in the steel sheet is increased, and slab cracking frequently occurs
during continuous casting. Therefore, N is limited to the range of 0.0050 to 0.0250%.
From the viewpoint of improvement in stability of material properties and yield over
the entire production process, N is preferably in the range of 0.0070 to 0.0170%.
With the N amount in the range of the present invention, there is no adverse effect
on weldability, and the like.
Dissolved N: 0.0010% or more
[0225] In order to ensure sufficient strength of a cold-rolled product, and effectively
exhibit strain age hardening with N, it is necessary that the content of dissolved
N (solid solution N) in the steel sheet is at least 0.0010% or more.
[0226] The amount of dissolved N is determined by subtracting the amount of precipitated
N from the total N amount of steel. As a result of comparison research of various
methods, the inventors found that electrolytic extraction analysis using constant-potential
electrolysis is effective as the method of analyzing the amount of. precipitated N.
As the method of dissolving ferrite used for extraction analysis, an acid digestion
method, a halogen method, or an electrolysis method can be used. Of these methods,
the electrolysis method can stably dissolve only ferrite without decomposing very
unstable precipitates such as a carbide, a nitride, etc. As the electrolyte, an acetyl-acetone
system is used for electrolysis at a constant potential. In the present invention,
the results of measurement of the amount of precipitated N by constant-potential electrolysis
showed best correspondence with changes in actual material properties.
[0227] Therefore, in the present invention, the residue after extraction by constant-potential
electrolysis is chemically analyzed to determine the amount of N in the residue. The
thus-determined value is considered as the amount of precipitated N.
[0228] In order to obtain higher BH and ΔTS, the amount of dissolved N is preferably 0.0020%
or more, more preferably 0.0020% or more. In order to obtain further high values,
the amount of dissolved N is preferably 0.0030% or more. Although the upper limit
of the amount of dissolved N is not limited, the mechanical properties less deteriorate
even when the all amount of N added remains.
N/Al (the content ratio of N to Al): 0.3 or more
[0229] In order to cause 0.0010% or more of dissolved N to stably remain in a product state,
the amount of Al which is an element for strongly fixing N, must be limited. As a
result of research of steel sheets in which the combination of the N content (0.0050
to 0.0250%) and the Al content (0.02% or less) were widely changed in the composition
range of the present invention, it was found that with N/Al of 0.3 or more, the amount
of dissolved N of a cold-rolled product or coated product can be stably set to 0.0010%
or more. Therefore, N/Al is limited to 0.3 or more.
[0230] In the present invention, the above component preferably further contains at least
one of the following groups d to g:
Group d: at least one of Cu, Ni, Cr and Mo in a total of 1.0% or less;
Group e: at least one of Nb, Ti and V in a total of 0.1% or less;
Group f: 0.00305 of B; and
Group g: one or both of Ca and REM in a total of 0.0010 to 0.010%.
[0231] Elements of group d: Cu, Ni, Cr and Mo are all contribute to an increase in strength
of the steel sheet, and can be contained singly or in a combination according to demand.
The effect is recognized by containing 0.005% or more each of Cu, Ni, Cr and Mo. However,
with an excessively high content, deformation resistance in hot rolling at elevated
temperatures is increased, or chemical conversion properties and surface treatment
properties in a wide sense deteriorate, and a welded portion is hardened to deteriorate
weld moldability. Also, the r value is liable to decrease. Therefore, the elements
of group a are preferably contained in a total of 1.0% or less. With a Mo content
of 0.05% or more, the r value is significantly decreased in some cases. In the present
invention, therefore, the Mo content is preferably limited to less than 0.05%.
[0232] Elements of group e: All of Nb, Ti and V are elements contributing refinement and
homegenization of crystal grains, and may be added singly or in a combination according
to demand. The effect can be recognized by containing 0.005% or more each of Nb, Ti
and V. However, with an excessively high content, deformation resistance in hot rolling
at elevated temperatures is increased, or chemical conversion properties and surface
treatment properties in a wide sense deteriorate. Therefore, the elements in group
b are preferably contained at a total of 0.1% or less.
[0233] Elements of group f: B is an element having the effect of improving hardenability
of steel, and can be contained for increasing the fraction of a low-temperature transformation
phase other than the ferrite phase to increase strength of steel according to demand.
This effect is recognized with a B content of 0.0005% or more. However, with an excessively
high B content, deformability at elevated temperatures in hot rolling deteriorates
to produce BN, decreasing the amount of dissolved N. Therefore, the B content is preferably
0.0030% or less.
[0234] Elements in group g: Both Ca and REM are elements useful for controlling the form
of inclusions. Particularly, when the stretch flanging property is required, these
elements are preferably added singly or in a combination. When the total of the elements
of group d is less than 0.0010%, the effect of controlling the form of inclusions
is insufficient, while when the total exceeds 0.010%, surface defects significantly
occur. Therefore, the total of the elements of group d is preferably limited to the
range of 0.0010 to 0.010%. This permits improvement in the stretch flanging property
without causing surface defects.
[0235] The structure of the steel sheet of the present invention is described below.
Area ratio of ferrite phase: 80% or more
[0236] The cold-rolled steel sheet of the present invention is directed to use as an automobile
steel sheet required to have some extent of moldability, and has a structure containing
the ferrite phase at an area ratio of 80% or more in order to ensure ductility. With
the ferrite phase at an area ratio of less than 80%, it is difficult to ensure ductility
required for an automobile steel sheet required to have moldability. When good ductility
is required, the area ratio of the ferrite phase is preferably 85% or more. In the
present invention, "ferrite" means so-called polygonal ferrite in which no strain
remains.
Average crystal grain diameter of ferrite phase: 10 µm or less
[0237] In the present invention, the value used as the average crystal grain diameter is
a higher value of the value calculated from a photograph of a sectional structure
by a quadrature method defined by ASTM, and the nominal value determined by an intercept
method defined by ASTM (refer to, for example, Umemoto et al.: Heat Treatment, 24
(1984), p334).
[0238] The cold-rolled steel sheet of the present invention maintains a predetermined amount
of dissolved N in the product step. However, as a result of experiment and research
conducted by the inventors, it was found that variations in strain age hardenability
occur in steel sheets containing the same amount of dissolved N, and one of the main
causes of the variations is a crystal grain diameter. In the structure of the present
invention, in order to stably obtain a high BH amount and ΔTS, the average crystal
grain diameter is at least 10 µm or less, and preferably 8 µm or less. Although the
detailed mechanism is not known, this is supposed to be related to the segregation
and precipitation of alloy elements in the crystal grain boundaries, and the influences
of processing and heat history on the segregation and precipitation.
[0239] Therefore, in order to achieve stability of strain age hardenability, the average
crystal grain diameter of the ferrite phase is 10 µm or less, and preferably 8 µm
or less.
[0240] In order to ensure ductility of an automobile steel sheet, and stability of strain
age hardenability, the structure of the present invention contains the ferrite phase
with an average crystal grain diameter of 10 µm or less at an area ratio of 80% or
more.
Area ratio of martensite phase: 2% or more
[0241] The cold-rolled steel sheet of the present invention contains the martensite phase
as a second phase at an area ratio of 2% or more. The presence of 2% or more of the
martensite phase can produce good ductility and a large amount of strain age hardening.
Although the detailed mechanism is not known, this effect is supposed to be due to
the effective accumulation of strain in the steel sheet due to the presence of the
martensite phase during pre-strain processing before aging. Furthermore, the presence
of the martensite phase is effective to improve aging deterioration. In order to a
good balance between strength and ductility and a low yield ratio, the area ratio
of the martensite phase is preferably 5% or more. The presence of the martensite phase
at an area ratio of over 20% causes the problem of deteriorating ductility. Therefore,
the area ratio of the martensite phase is 2% or more, and preferably 5% to 20%.
[0242] Besides the above-described martensite phase, pearlite, bainite, residual austenite
are present as second phases without causing any problem. However, in the present
invention, it is necessary that the fraction of the ferrite phase is 80% or more,
and the fraction of the martensite phase is 2% or more. Therefore, the total area
ratios of pearlite, bainiate and residual austenite are limited to less than 18%.
[0243] The cold-rolled steel sheet of the present invention, which has the above-described
composition and structure, has a tensile strength (TS) of 440 MPa to about 780 MPa,
a high r value of 1.2 or more obtained by controlling the aggregation structure of
the ferrite base phase, and excellent strain age hardenability. A steel sheet having
TS of less than 440 MPa cannot be widely applied to members having structural components.
Furthermore, in order to extend the application range, TS is preferably 500 MPa or
more. With the r value of less than 1.2, the steel sheet cannot be applied to a wide
range of press forming parts. The preferable range of the r value is 1.4 or more.
[0244] As described above, in the present invention, "excellent strain age hardenability"
means that in aging under conditions of holding at a temperature of 170°C for 20 min.
after pre-deformation with a tensile strain of 5%, the increment in deformation stress
(represented by the amount of BH = yield stress after aging - pre-deformation stress
before aging) after aging is 80 MPa or more, and the increment in tensile strength
(represented by ΔTS = tensile strength after aging - tensile strength without strain
aging) after strain aging (pre-deformation + aging) is 40 MPa or more.
[0245] In defining the strain age hardenability, the amount of pre-strain (pre-deformation)
is an important factor. As a result of research of the influence of the amount of
pre-strain on strain age hardenability, the inventors found that (1) the deformation
stress in the above-described deformation system can be referred to as an amount of
approximately uniaxial strain (tensile strain) except the case of excessive deep drawing,
(2) the amount of uniaxial strain of an actual part exceeds 5%, and (3) the strength
of a part sufficiently corresponds to the strength (YS and TS) obtained after strain
aging with a pre-strain of 5%. In the present invention, based on these findings,
the pre-deformation of strain aging is defined to a tensile strain of 5%.
[0246] Conventional coating and baking conditions include 170°C and 20 min as standards.
When a strain of 5% is applied to the steel sheet of the present invention, which
contains a large amount of dissolved N, hardening can be achieved even by aging at
low temperature. In other words, the range of aging conditions can be widened. In
order to attain a sufficient amount of hardening, generally, retention at a higher
temperature for a longer time is advantageous as long as softening does not occurs
by over aging.
[0247] Specifically, in the steel sheet of the present invention, the lower limit of the
heating temperature at which hardening significantly takes place after pre-deformation
is about 100°C. On the other hand, with the heating temperature of over 300°C, hardening
peaks, thereby causing the tendency to soften and significantly causing thermal strain
and temper color. With the retention time of about 30 seconds or more, hardening can
be sufficiently achieved at a heating temperature of about 200°C. In order to obtain
more stable hardening, the retention time is preferably 60 seconds or more. However,
retention for over 20 mines is practically disadvantageous because further hardening
cannot be expected, and the production efficiency significantly deteriorates.
[0248] Therefore, in the present invention, the conventional coating and baking conditions,
i.e., the heating temperature of 170°C and the retention time of 20 minutes, are set
as the aging conditions. With the steel sheet of the present invention, hardening
can be stably achieved even under the aging conditions of a low heating temperature
and a short retention time, which fail to achieve sufficient hardening in a conventional
bake-hardening steel sheet. The heating method is not limited, and atmospheric heating
with a furnace, which is generally used for coating and baking, and other methods
such as induction heating, heating with a nonoxidation flame, a laser, plasma, or
the like, etc. can be preferably used.
[0249] The strength of an automobile part must be sufficient to resist an external complicated
stress load, and thus not only strength in a low strain region but also strength in
a high strain region are important for a raw material steel sheet. In consideration
of this point, in the steel sheet of the present invention used as a raw material
for automobile parts, BH is 80 MPa or more, and ΔTS is 40 MPa or more. More preferably,
BH is 100 MPa or more, and ΔTS is 50 MPa or more. In order to further increase BH
and TS, the heating temperature in aging may be set to a higher temperature, and/or
the retention time may be set to a longer time.
[0250] The steel sheet of the present invention has the advantage that when the steel sheet
is allowed to stand at room temperature for about one week without heating after forming,
an increase in strength of about 40% of that at the time of complete aging can be
expected.
[0251] The steel sheet of the present invention also has the advantage that even when it
is allowed in an unmolded state at room temperature for a long time, aging deterioration
(an increase in YS and a decrease in El (elongation)) does not occurs, unlike a conventional
aging steel sheet. In order to prevent the occurrence of a trouble in actual press
forming, it is necessary that in aging at room temperature for 3 months before press
forming, an increase in YS is 30 MPa or less, a decrease in elongation is 2% or less,
and a recovery of yield point elongation is 0.2% or less.
[0252] In the present invention, the surface of the cold-rolled steel sheet may be coated
by hot-dip galvanization or alloying hot-dip galvanization without any problem, and
TS, BH and ΔTS are equivalent to those before plating. AS the plating method, electro-galvanization,
hot-dip galvanization, alloying hot-dip galvanization, electro-tinning, electric chromium
plating, electro-nickeling, and the like may be preferably used.
[0253] The method of producing a steel sheet according to an eighth embodiment of the present
invention will be described.
[0254] The steel sheet of the present invention is basically produced by performing the
hot rolling step in which a steel slab having the above-described composition is heated,
and then roughly rolled to form a sheet bar, and the sheet bar is finish-rolled and
cooled to form a coiled hot-rolled sheet, the cold rolling step in which the hot-rolled
sheet is pickled and cold-rolled to form a cold-rolled sheet, and the cold-rolled
sheet annealing step in which the cold-rolled sheet is box-annealing and then continuously
annealed.
[0255] Although the slab used in the production method of the present invention is preferably
produced by a continuous casting method in order to prevent macro-segregation of components,
an ingot making method or a thin slab casing method may be used. Alternatively, a
conventional method comprising cooling the produced slab to room temperature and then
again heating the slab, or an energy-saving process of direct rolling comprising charging
a hot slab into a heating furnace without cooling and then rolling it, or rolling
directly the slab immediately after slightly keeping it warm may be used without a
problem. Particularly, direct rolling is a useful technique for effectively ensuring
dissolved N.
[0256] Description will be now be made of the reasons for limiting the conditions of the
hot rolling step.
Slab heating temperature: 1000°C or more
[0257] The slab heating temperature is preferably 1000°C or more in order to ensure a necessary
and sufficient amount of dissolved N in an initial state and satisfy the target amount
of dissolved N in a product. Since a loss is increased by an increase in the oxide
weight, the heating temperature is preferably 1280°C or less.
[0258] The slab heated under the above condition is roughly rolled to form a sheet bar.
The condition of the rough rolling is not defined, and rough rolling may be performed
according to a normal method. However, in order to ensure the amount of dissolved
N, rough rolling is preferably performed within a as short time as possible. Then,
the sheet bar is finish-rolled to form a hot-rolled sheet.
[0259] In the present invention, the adjacent sheet bars are preferably bonded together
during the time between rough rolling and finish rolling, and then continuously rolled.
As the bonding means, a pressure welding method, a laser welding method, an electron
beam welding method, or the like is preferably used.
[0260] By continuous rolling, so-called non-stationary portions at the front end and the
rear end of a coil (processed material) are removed to permit hot rolling over the
entire length and the entire width of the coil (processed material) under stable conditions.
This is very effective to improve the sectional shape and dimensions of not only the
hot-rolled steel sheet but also the cold-rolled steel sheet. When the steel sheet
is cooled on a hot run table after rolling, the shape of the steel sheet can be sufficiently
maintained because tension can be always applied.
[0261] By continuous rolling, the ends of the coil can be stably passed, and it is thus
possible to use lubricating rolling, which cannot be easily applied to ordinary single
rolling for each sheet bar because of the problem of continuous rolling processes
and biting property. Therefore, the rolling load can be decreased, and at the same
time, the surface pressure of the roll can be decreased, thereby increasing the life
time of the roll.
[0262] In the present invention, at the entrance of a finisher between rough rolling and
finish rolling, the temperature distributions of the sheet bar in the width direction
and the long direction thereof are preferably made uniform by using any one or both
of a sheet bar edge heater for heating the ends of the sheet bar in the width direction
and a sheet bar heater for heating the ends of the sheet bar in the long direction.
This can further decrease the variations in material properties of the steel sheet.
The sheet bar edge heater and the sheet bar heater are preferably of an induction
heating type.
[0263] As the procedure of use of the heaters, preferably, the difference in temperature
in the width direction is first corrected by the sheet bar edge heater. Although the
heating amount depends upon the steel composition, the temperature distribution in
the width direction at the finisher entrance is preferably set in the range of about
20°C or less. Next, the difference in temperature in the long direction is corrected
by the sheet bar heater. At this time, the heating amount is preferably set so that
the temperatures at the ends in the long direction are about 20°C higher than the
temperature at the center.
Finisher delivery temperature: 800°C or more
[0264] In order to obtain a homogeneous fine hot-rolled base sheet structure, the finisher
deliver temperature FDT is 800°C or more. With a FDT of lower than 800°C, the structure
of the steel sheet becomes inhomogeneous, and the processed structure partially remains
to leave heterogeneity of the structure after the cold-rolled sheet annealing step.
Therefore, the danger of causing various troubles in press forming is increased. When
a high coiling temperature is used for avoiding the processed structure from remaining,
coarse crystal grains are produced to cause the same troubles as described above.
With the high coiling temperature, the amount of dissolved N is significantly decreased
to cause difficulties in obtaining a target tensile strength of 440 MPa or more. Therefore,
the finisher deliver temperature FDT is 800°C or more. In order to further improve
the mechanical properties, the FDT is preferably 820°C or more. Although the upper
limit of FDT is not limited, a scale scar significantly occurs at excessively high
FDT. The FDT is preferably up to about 1000°C.
[0265] Although cooling after finish rolling is not strictly limited, the conditions described
below are preferable from the viewpoint of homogeneity in material properties of the
steel sheet in the long direction and the width direction thereof. In the present
invention, cooling is preferably started immediately after (within 0.5 seconds after)
finish rolling, and the mean cooling rate in cooling is preferably 40°C/s or more.
By satisfying these conditions, the steel sheet can be rapidly cooled in the high
temperature region where AlN precipitates to effectively ensure N in a solid solution
state. When the starting time of cooling or the cooling rate does not satisfy the
above condition, grain growth excessively proceeds to fail to achieve fine crystal
grains, and promote AlN precipitation due to stain energy introduced by rolling. Therefore,
the amount of dissolved N tends to decrease, and the structure tends to be made inhomogeneous.
In order to ensure homogeneity in material properties and shape, the cooling rate
is preferably kept at 300°C/s or less.
Coiling temperature: 800°C or less
[0266] The strength of the steel sheet is liable to increase as the coiling temperature
CT decreases. In order to ensure the target tensile strength TS of 440 MPa or more,
the CT is preferably 800°C or less. With a CT of less than 200°C, the shape of the
steel sheet is readily disturbed to increase the danger of causing troubles in a practical
operation, thereby deteriorating homogeneity of material properties. Therefore, the
CT is preferably 200°C or more. When the homogeneity of the material properties is
required, the CT is preferably 300°C or more, and more preferably 350°C or more. In
the present invention, in finish rolling, lubricating rolling may be performed for
decreasing the hot rolling load. The lubricating rolling has the effect of further
making homogeneous the shape and material properties of the hot-rolled sheet. During
the lubricating rolling, the frictional coefficient is preferably in the range of
0.25 to 0.10. By combining lubricating rolling and continuous rolling, the operation
of hot rolling is further stabilized.
[0267] The hot-rolled sheet subjected to the above-described hot rolling step is then pickled
and cold-rolled in the cold rolling step to form a cold-rolled sheet.
[0268] The pickling conditions may be the same as conventional known conditions, and are
not limited. When the scale of the hot-rolled sheet is extremely small, cold rolling
may be immediately after hot rolling without pickling.
[0269] The cold rolling conditions may be the same as conventional known conditions, and
are not limited. In order to ensure homogeneity of the structure, the reduction ratio
of cold rolling is preferably 40% or more. The reasons for limiting the conditions
of the cold rolling step are described below.
[0270] The cold-rolled sheet is then subjected to the cold-rolled sheet annealing step comprising
box annealing and continuous annealing.
Box annealing temperature: the recrystallization temperature to 800°C
[0271] In the present invention, the cold-rolled sheet is subjected to box annealing to
control the aggregation structure of the ferrite phase as a base. By controlling the
aggregation structure of the ferrite phase, the r value of the produced sheet can
be increased. By box annealing, the (111) aggregation structure suitable for increasing
the r value is readily formed in the produced sheet.
[0272] With the box annealing temperature less than the recrystallization temperature, recrystallization
is not completed to fail to control the aggregation structure of the ferrite phase,
thereby failing to increase the r value. On the other hand, with the box annealing
temperature of over 800°C, surface defects significantly occur in the steel sheet,
thereby failing to achieve the initial purpose. Box annealing is preferably performed
in an annealing atmosphere containing a nitrogen gas as a main component and 3 to
5% of hydrogen gas. In this case, the heating and cooling rates may be the same as
normal box annealing, and are about 30°C/hr. By using 100% hydrogen gas as an annealing
atmosphere gas, the higher heating and cooling rates may be used.
Continuous annealing temperature: Ac
1 transformation point to (Ac
3 transformation point - 20°C)
[0273] With the continuous annealing temperature of less than the Ac
1 transformation point, the martensite phase is not formed after annealing, while with
the continuous annealing temperature of over (Ac
3 transformation point - 20°C), the desirable aggregation structure formed in box annealing
is lost due to transformation, thereby failing to obtain the produced sheet having
a high r value. Therefore, the continuous annealing temperature is preferably Ac
1 transformation point to (Ac
3 transformation point - 20°C). The retention time of continuous annealing is preferably
as short as possible in order to ensure the production efficiency, the fine structure
and the amount of dissolved N. -From the viewpoint of stability of the operation,
the retention time is preferably 10 seconds or more. Also, in order to ensure the
fine structure and the amount of dissolved N, the retention time is preferably 120
seconds or less. From the viewpoint of stability of material properties, the retention
time is preferably 20 seconds or more.
Cooling after continuous annealing: cooling to the temperature region of 500°C
or less at a cooling rate of 10 to 300°C/s
[0274] Cooling after soaking by continuous annealing is important for making fine the structure,
forming the martensite phase, and ensuring the amount of dissolved N. In the present
invention, cooling is continuously carried out to the temperature region of at least
500°C or less at a cooling rate of 10°C/s or more. With a cooling rate of less than
10°C/s, a necessary amount of martensite phase, a homogeneous fine structure and a
sufficient amount of dissolved N cannot be obtained. On the other hand, with a cooling
rate of over 300°C/s, homogeneity in material properties of the steel sheet in the
width direction deteriorates due to a significant increase in the amount of supersaturated
dissolved C. When the stop temperature of cooling at a cooling rate of 10 to 300°C/s
after continuous annealing exceeds 500°C, refinement of the structure cannot be attained.
Over aging condition: retention in the temperature region of 350°C to the cooling
stop temperature for 20 seconds or more subsequent to cooling after continuous annealing
Over aging may be performed by retention in the temperature region of 350°C to
the cooling stop temperature for 20 seconds or more subsequent to the stop of cooling
after soaking by continuous annealing. By over aging, the amount of dissolved C can
be selectively decreased, while the amount of dissolved N is maintained. With the
retention temperature region of less than 350°C, a long time is required for decreasing
the amount of dissolved C to cause a reduction in productivity. Therefore, the temperature
region is preferably 350°C or more.
[0275] By retention in the temperature region of 350°C to the cooling stop temperature for
20 seconds or more, the amount of dissolved C can be decreased to achieve a higher
degree of non-aging properties at room temperature. By increasing the retention time,
further improvement can be expected, but the effect is saturated with the retention
time of about 120 seconds. Therefore, the retention time is preferably 120 second
or less.
[0276] In order to obtain a larger amount of strain age hardening, it is advantageous to
use either of dissolved C and dissolved N. However, by using dissolved C, aging deterioration
at room temperature becomes significant, thereby limiting parts to which the steel
sheet is applied.
Therefore, in order to produce a strain age hardenable steel sheet having versatility,
over aging is preferably performed with the sufficient amount of dissolved N being
ensured.
[0277] In producing a high-tensile-strength cold-rolled coated steel sheet comprising a
high-tensile-strength cold-rolled steel sheet and a hot-dip coated layer formed on
the surface thereof, continuous annealing after box annealing can be performed in
a continuous hot-dip coating line comprising hot-dip galvanization subsequent to cooling
after continuous annealing or further alloying to produce a hot-dip galvanized steel
sheet.
Temper rolling or lever processing: elongation of 0.2 to 15%
[0278] In the present invention, in order to correct the shape and control roughness, temper
rolling or leveler processing may be carried out subsequent to the cold rolling step.
When the total elongation of temper rolling or leveler processing is less than 0.2%,
the desired purpose of correcting the shape and controlling roughness cannot be achieved.
On the other hand, with a total elongation of over 15%, ductility significantly deteriorates.
It is confirmed that the processing system of temper rolling is different from that
of leveler processing, but the effects of both processes are substantially the same.
Temper rolling and leveler processing are effective after plating.
[0279] For reference, description will now be made of forming conditions and conditions
for subsequent heat treatment for increasing strength when the steel sheet of the
present invention is molded, for example, press-molded. When the steel sheet of the
present invention is subjected to press working, for example, deep drawing, the strain
introduced by press working is several % to several tens %. Although the amount of
strain changes with molded parts, a strain of about 5 to 10% is introduced into an
inner plate and a structural member in the automobile field.
[0280] These automobile parts are heat-treated by coating and baking. However, with the
steel sheet of the present invention, strength of a molded product can be effectively
increased after heat treatment. In the present invention, as a method of evaluating
burning hardenability in a laboratory, a tensile test specimen of JIS No. 5 size is
obtained from the steel sheet in the rolling direction, and tensile strain of 10%
is applied to the tensile test specimen by a tensile testing machine. Then, the specimen
is heat-treated and again subjected to a tensile test. Particularly, when properties
are evaluated after heat treatment in a low temperature region, the heat treatment
conditions include 120°C and 20 minutes. In this test, the properties of the completed
portion after heat treatment subsequent to press forming are evaluated.
[0281] Namely, in the present invention, the difference (ΔTS) between the tensile strength
of the specimen after application of tensile strain and heat treatment and the tensile
strength of a product is defined as the strength increasing ability of heat treatment.
[0282] In order to increase the strength of the molded product, the amount of strain introduced
by forming, or the heat treatment temperature after processing is preferably as high
as possible.
[0283] However, with the steel sheet of the present invention, when the amount of applied
strain is about 5 to 10%, the strength can be sufficiently increased even by heat
treatment at a temperature lower than conventional heat treatment, i.e., a temperature
of 200°C or less, after forming. However, with a heat treatment temperature of less
than 120°C, the strength cannot be sufficiently increased with the low train applied.
On the other hand, with the heat treatment temperature of over 350°C after forming,
softening proceeds. Therefore, the temperature of heat treatment after forming is
preferably about 120 to 350°C.
[0284] The heating method is not limited, and hot gas heating, infrared furnace heating,
hot-bath heating, direct current heating, induction heating, and the like can be used.
Alternatively, only a portion where strength is desired to be increased is selectively
heated.
Examples
[0285] In the examples below, the amount of dissolved N, the microstructure, tensile properties,
the r value, strain age hardenability, and aging property were examined. The examination
methods were as follows:
(1) Amount of dissolved N
[0286] The amount of dissolved N was determined by subtracting the amount of precipitated
N from the total N amount of steel determined by chemical analysis. The amount of
precipitated N was determined by an analysis method using a constant-potential electrolytic
method.
(2) Microstructure
[0287] A test specimen was obtained from each of cold-rolled annealed steel sheets, and
the microstructure of a section (C section) perpendicular to the rolling direction
was imaged with an optical microscope or a scanning electron microscope. Then, the
fraction of the ferrite texture and the type and the structure fraction of a second
phase were determined by an image analysis apparatus.
(3) Crystal grain diameter
[0288] In the present invention, the value used as the average crystal grain diameter was
a higher one of the value calculated from a photograph of a sectional structure by
a quadrature method defined by ASTM, and the nominal value determined from a photograph
of a sectional structure by an intercept method defined by ASTM (refer to, for example,
Umemoto et al.: Heat Treatment, 24 (1984), p334).
(4) Tensile properties
[0289] A test specimen of JIS No. 5 was obtained from each of cold-rolled annealed steel
sheets in the rolling direction, and a tensile test was carried out with a strain
rate of 3 x 10
-3/s according to the regulations of JIS Z 2241 to determine yield stress YS, tensile
strength TS, and elongation El.
(5) Strain age hardenability
[0290] A test specimen of JIS No. 5 was obtained from each of cold-rolled annealed steel
sheets in the rolling direction, and a tensile strain of 5% was applied as pre-deformation.
Then, the specimen was subjected to heat treatment corresponding to coating and baking
at 170°C for 20 minutes, and a tensile test with a strain rate of 3 x 10
-3/s was performed to determine the tensile properties (yield stress TSBH, tensile strength
TSBH). Then, BH amount = YSBH - YS5%, and ΔTS = TSBH - TS were calculated. YS5% represents
deformation stress in 5% pre-deformation of the produced sheet, YSBH and TSBH represent
yield stress and tensile strength, respectively, after pre-deformation and heat treatment,
and TS represents the tensile strength of the produced sheet.
(6) Measurement of r value
[0291] A test specimen of JIS No. 5 was obtained from each of the cold-rolled annealed steel
sheets in each of the rolling direction (L direction), the direction (D direction)
at 45° with the rolling direction, and the direction (C direction) at 90° with the
rolling direction. The width-direction strain and thickness-direction strain of each
of the test specimens were determined when a uniaxial tensile strain of 15% was applied
to each specimen, and the r value of each specimen in each of the directions was determined
from the following ratio of width-direction strain to thickness-direction strain:
(wherein w0 and t0 represent the width and thickness of a specimen before the
test, and wand t represent the width and thickness of a specimen after the test).
The mean value was determined by the following equation:
wherein rL represents the r value in the rolling direction (L direction), rD represents
the r value in the direction (D direction) at 45° with the rolling direction, and
rL represents the r value in the direction (C direction) at 90° with the rolling direction.
In order to improve the precision of experiment, calculation was made by using changes
in elongation strain and strain in the width direction on the assumption that the
volume was constant.
(7) Aging properties
[0292] A test specimen of JIS No. 5 was obtained from each of cold-rolled annealed steel
sheets in the rolling direction, and then subjected to aging at 50°C for 200 hours,
followed by a tensile test. The difference in yield elongation ΔY-E1 between before
and after aging was determined from the obtained results to evaluate aging properties
at normal temperature. When ΔY-El was zero, it was evaluated that the specimen has
non-aging properties and excellent natural aging resistance.
(8) Tensile strength after forming and heat treatment
[0293] A test specimen of JIS No. 5 was obtained from each of produced sheets in the rolling
direction, and then a pre-strain of 10% was applied thereto. Then, heat treatment
was conducted for 20 minutes at a conventional heat treatment temperature of 120°C
and a temperature of 170°C corresponding to coating and baking, and then tensile strength
was determined.
(9) Decrease (ΔE1) in total elongation by natural aging
[0294] The decrease (ΔE1) in total elongation by natural aging was determined as the difference
between the total elongation measured with a specimen of JIS N0 5 obtained from the
produced sheet in the rolling direction, and the total elongation measured with a
specimen of JIS N0 5 separately obtained from the produced sheet in the rolling direction
after accelerated aging (retention at 100°C for 8 hours) of natural aging.
Example 1
[0295] A steel slab having each of the compositions shown in Table 1 was hot-rolled sheet
having a thickness of 3.5 mm, and then cold-rolled to a cold-rolled sheet having a
thickness of 0.7 mm under the conditions shown in Table 2. Then, the cold-rolled sheet
was recrystallized, annealed and further galvannealed in a continuous annealing line
or a continuous annealing and galvanizing line. Then, the annealed sheet was temper-rolled
with a rolling reduction ratio of 1.0 % to produce a cold-rolled steel sheet and a
galvannealed steel sheet having both sides coated with a weight of 45 g/m
2 per side. In Table 2, the finisher deliver temperatures of the others are the Ar
3 transformation point or more.
[0296] The thus -obtained cold-rolled steel sheets and the galvannealed steel sheets were
measured with respect to tensile strength, the r value, and a change in tensile strength
after forming and heat treatment. The results are shown in Table 3.
[0297] Table 3 indicates that with all the cold-rolled steel sheets and the galvannealed
steel sheets obtained according to the present invention, a high r value and excellent
strain age hardenability are obtained, as compared with comparative examples. Particularly,
in the suitable examples in which the crystal grain diameter is 20 µm or less, the
decrease in elongation due to natural aging is also as low as 2.0% or less.
Example 2
[0298] A slab of steel symbol B shown in Table 1 was hot-rolled under the same production
conditions as No. 2 shown in Table 2 in which the heating temperature was 1100°C,
and the finisher deliver temperature of hot rolling was 900°C, and then coiled at
coiling temperature of 550°C into a coil. The thus-obtained coil was cold-rolled with
a reduction ratio of 80%, and then recrystallized and annealed at 840°C. With respect
to the product properties of the resultant cold-rolled steel sheet, tensile strength
TS was 365 MPa, and the r value was 1.7. A test specimen of JIS No. 5 was obtained
from the cold-rolled steel sheet in the rolling direction, and a tensile strain of
10% was applied by a tensile test machine. Then, the specimen was subjected to heat
treatment under the heat treatment conditions (temperature and time) shown in Table
4, and a tensile test was again performed. Table 4 also shows the increase in tensile
strength (ΔTS) from the tensile strength (TS = 365 MPa) of a product before application
of strain.
[0299] Table 4 indicates that the increase in strength increases as the heat treatment temperature
increases, and the heat treatment time increases. However, with the steel sheet of
the present invention, a sufficient increase in tensile strength of 82 MPa (85% or
more of an increase in heat treatment for 20 minutes) can be obtained even by heat
treatment at low temperature of 120°C for a short retention time of 2 minutes. It
is thus found that with the steel sheet of the present invention, good strain age
hardenability can be obtained even by heat treatment at a low temperature for a short
time. In order to obtain the stable effect of increasing strength of an automobile
structural member, or the like, heat treatment at a normal temperature for a normal
time causes no problem. It was confirmed that with the galvanized steel sheets and
the galvannealed steel sheets obtained by hot-dip galvanizing and heat alloying the
cold-rolled sheets, the same results as shown in Table 4 are obtained.
Example 3
[0300] A steel slab having each of the compositions shown in Table 6 was hot-rolled under
the conditions shown Table 7 to form a hot-rolled sheet having a thickness of 3.5
mm. Each of the thus-obtained hot-rolled sheets was cold-rolled under the conditions
shown in Table 7 to form a cold-rolled sheet having a thickness of 0.7 mm, and then
recrystallized and annealed under the conditions shown in the same table. Some of
the annealed sheets were further coated by hot-dip galvanization or alloying hot-dip
galvanization under the conditions shown in the same table. The thus-obtained produced
sheets were examined with respect to the amount of dissolved N, the microstructure,
tensile properties, and strain age hardenability.
[0301] The results are shown in Table 8. Table 8 indicates that with the steel sheets of
the present invention, TS x r value ≥ 750 MPa (in a combination with at least one
of B, Nb, Ti and V, Ts x r value 850 ≥ MPa), BH ≥ 80 MPa and ΔTS ≥ 40 MPa are satisfied,
while in the comparative examples, at least one of the three properties does not reach
the level of the present invention.
Example 4
[0302] Examples of the present invention are described below.
[0303] Melted steel having each of the compositions shown in Table 9 was formed in an ingot
by a converter, and then formed in a steel slab by a continuous casting method. Each
of the resultant steel slabs was heated under the conditions shown in Table 10, and
roughly rolled to form a sheet bar. The sheet bar was hot-rolled by the hot rolling
step comprising finish rolling under the conditions shown in Table 10 to form a hot-rolled
sheet. The Ar
3 transformation point was measured by simulating the conditions for hot rolling finish
rolling using a processing transformation measuring apparatus (produced by Fuji Denpa
Kouki). The results are shown in Table 10.
[0304] Each of the thus-obtained hot-rolled sheets was cold-rolled by the cold rolling step
under the conditions shown in Table 10 to form a cold-rolled sheet. Then, each of
the cold-rolled sheets was continuously annealed under the conditions shown in Table
10. Some of the cold-rolled sheets were further temper-rolled after the cold-rolled
sheet annealing step.
[0305] The thus-obtained cold-rolled annealed sheets were examined with respect to the amount
of dissolved N, the microstructure, tensile properties, the r value, strain age hardenability
and aging properties.
[0306] The surface of each of the steel sheets of Nos. 4 and 10 was hot-dip galvanized to
form a coated steel sheet, and evaluated with respect to the same properties as described
above.
[0307] The results are shown in Table 11.
[0308] All examples of the present invention have excellent ductility, an extremely high
BH amount and ΔTS, excellent strain age hardenability, a mean r value of as high as
1.2 or more, and non-aging properties at natural aging (excellent natural aging resistance).
With respect to the properties of the coated steel sheets obtained by hot-dip galvanizing
the surfaces of the steel sheets of Nos. 4 and 10, the mean r value is decreased by
0.2, and the elongation El is decreased by about 1%, as compared with the cold-rolled
steel sheets, because of shrinkage restriction of the coated layer in the width direction.
However, strain age hardenability and natural aging resistance are substantially the
same as those before coating. On the other hand, in the comparative examples out of
the range of the present invention, ductility deteriorates, the BH amount and ΔTS
are low, or natural aging deterioration significantly occurs. Therefore, the comparative
examples do not have all the intended properties, and thus cannot be said as steel
sheets having sufficient properties.
[0309] Steel sheet No. 11 contains C, Al, N, and N/Al out of the range of the present invention,
and thus the r value, the BH amount, ΔTS and natural aging resistance deteriorate.
Steel sheet No. 12 contains B and Nb out of the range of the present invention, and
thus the amount of acicular ferrite is greatly deviated from the range of the present
invention, deteriorating the BH amount, ΔTS, and natural aging resistance.
[0310] Steel sheet No. 13 contains B out of the range of the present invention, and thus
the amount of acicular ferrite is greatly deviated from the range of the present invention,
deteriorating the r value, the BH amount, ΔTS, and natural aging resistance. Steel
sheet No. 14 contains Nb out of the range of the present invention, and thus the amount
of dissolved N is greatly lower than the range of the present invention, deteriorating
strain age hardenability.
[0311] Steel sheet No. 15 contains N out of the range of the present invention, and thus
the amount of dissolved N is low, deteriorating strain age hardenability. In Steel
Nos. 17 to 20, the hot rolling conditions and the cold-rolled sheet annealing conditions
are deviated from the suitable ranges, and thus the microstructure is out of the range
of the present invention, decreasing the BH amount and ΔTS and deteriorating strain
age hardenability and natural aging resistance.
Example 5
[0312] Steel having the composition shown in Table 12 was formed in a slab by the same method
as Example 4, and then heated and temper-rolled under the conditions shown in Table
13 to form a sheet bar. The sheet bar was then hot-rolled by the hot rolling step
comprising finish rolling under the conditions shown in Table 13 to form a hot-rolled
sheet. The adjacent sheet bars on the finisher entrance side after rough rolling were
bonded together by a melt welding method and then continuously rolled. The temperatures
of the ends of the sheet bar were controlled in the width direction and the length
direction by using an induction heating-type sheet bar edge heater and a sheet bar
heater.
[0313] The thus-obtained hot-rolled sheet was cold-rolled by the cold rolling step comprising
pickling and cold rolling under the conditions shown in Table 13 to form a cold-rolled
sheet having a thickness of 1.6 mm. Then, the cold-rolled sheet was continuously annealed
under the conditions shown in Table 13.
[0314] The thus-obtained cold-rolled annealed sheet was examined with respect to the amount
of dissolved N, the microstructure, tensile properties, the r value, and strain age
hardenability by the same methods as Example 4. The tensile property of each cold-rolled
annealed sheet was measured at ten positions in each of the width direction and the
long direction to examine variations in yield strength, tensile strength and elongation.
[0315] The results are shown in Table 14.
[0316] All the examples of the present invention have excellent strain age hardenability
and a high r value, and exhibit extremely high stable BH amount, ΔTS and mean value
regardless of variations in production conditions. It was also recognized that in
the examples of the present invention, by performing continuous rolling and controlling
the temperature of the sheet bar in the long direction and the width direction, the
thickness precision and the shape of the produced steel sheet are improved, and variations
in material properties are decreased to 1/2. Even when the elongation of temper rolling
is changed to 0.5 to 2%, and the elongation of the leveler is changed to 0 to 1%,
strain age hardenability does not deteriorate.
Example 6
[0317] Examples of the present invention are described below.
[0318] Melted steel having each of the compositions shown in Table 15 was formed in an ingot
by a converter, and then formed in a slab by a continuous casting method. Each of
the steel slabs was heated (in some cases, a hot slab was charged) and roughly rolled
under the conditions shown in Table 16 to form a sheet bar. The sheet bar was then
hot-rolled by the hot rolling step comprising finish rolling under the conditions
shown in Table 16 to form a hot-rolled sheet. With some of the sheet bars, the adjacent
sheet bars were bonded by the melt welding method, and then continuously rolled.
[0319] Each of the resultant hot-rolled sheets was cold-rolled in the cold rolling step
comprising pickling and cold rolling under the conditions shown in Table 16 to form
a cold-rolled sheet. Each of the thus-obtained cold-rolled sheet was box-annealed
and then continuously annealed under the conditions shown in Table 16. Some of the
cold-rolled sheets were temper-rolled after the cold-rolled sheet annealing step.
Box annealing may not be carried out. In all cases, the annealing temperature of box
annealing was the recrystallization temperature or more.
[0320] The thus-obtained cold-rolled annealed sheets were examined with respect to the amount
of dissolved N, the microstructure, tensile properties, the r value, strain age hardenability,
and the aging property.
[0321] The surfaces of the steel sheets of Nos. 17 and 18 were coated by hot-dip galvanization
in an in-line after continuous annealing shown in Table 16 to form coated steel sheets.
The coated steel sheets were also examined with respect to the same properties as
described above.
[0322] The results are shown in Table 17.
[0323] All the examples of the present invention have excellent ductility, extremely high
stable BH amount and ΔTS, excellent strain age hardenability, a mean r value of as
high as 1.2, and natural non-aging properties. The properties of the hot-dip galvanized
steel sheets of Nos. 17 and 18 shown in Table 17 are substantially the same as the
cold-rolled steel sheets subjected to continuous annealing. On the other hand, in
the comparative examples out of the range of the present invention, ductility deteriorates,
the BH amount and TS are low, or aging deterioration significantly occurs. Therefore,
the comparative examples do not have all the intended properties, and thus cannot
be said as steel sheets having sufficient properties.
[0324] Steel sheet No. 11 contains C and N in amounts out of the range of the present invention,
and has an amount of dissolved N and a martensite amount lower than the range of the
present invention. Therefore, the BH amount and ΔTS are decreased, and ΔY-E1 is increased.
Steel sheet No. 12 contains Al, N/Al and N out of the range of the present invention,
and has an amount of dissolved N lower than the range of the present invention, and
the average crystal grain diameter of ferrite higher than the range of the present
invention. Therefore, the BH amount and ΔTS are decreased, and ΔY-E1 is increased.
[0325] In steel sheet No. 13, the slab heating temperature and finisher delivery temperature
FDT are out of the range of the present invention, the amount of dissolved N and the
amount of martensite are lower than the range of the present invention, and the average
crystal grain diameter of ferrite is higher than the range of the present invention.
Therefore, the r value, the BH amount and ΔTS are decreased. In steel sheet No. 14,
the coiling temperature after hot rolling is out of the range of the present invention,
the amount of dissolved N and the amount of martensite are lower than the range of
the present invention, and the average crystal grain diameter of ferrite is higher
than the range of the present invention. Therefore, the r value, the BH amount and
ΔTS are decreased.
[0326] In steel sheet No. 15, the continuous annealing temperature is out of the range of
the present invention, martensite is not formed, and the average crystal grain diameter
of ferrite is higher than the range of the present invention. Therefore, the BH amount
and ΔTS are decreased, and ΔY-E1 is increased. In steel sheet No. 16, box annealing
is not performed to fail to develop the desirable aggregation structure, deteriorating
the r value. Also, the average crystal grain diameter of ferrite, and the area ratio
of martensite are out of the range of the present invention.
Example 7
[0327] Steel having the composition shown in Table 18 was formed in a slab by the same method
as Example 1, and then heated and' roughly rolled under the conditions shown in Table
19 to form a sheet bar having a thickness of 30 mm. The sheet bar was hot-rolled by
the hot rolling step comprising finish rolling under the conditions shown in Table
19 to form a hot-rolled sheet. With some of the sheet bars, the adjacent sheet bars
on the finisher entrance side after rough rolling were bonded together' by the melt
welding method, and then continuously rolled. The temperatures of the ends of the
sheet bar were controlled in the width direction and the length direction by using
an induction heating-type sheet bar edge heater and a sheet bar heater.
[0328] The thus-obtained hot-rolled sheet was cold-rolled by the cold rolling step comprising
pickling and cold rolling under the conditions shown in Table 19 to form a cold-rolled
sheet having a thickness of 1.6 mm. Then, the cold-rolled sheet was box-annealed and
then continuously annealed by a continuous annealing furnace under the conditions
shown in Table 19. In all cases, the annealing temperature of box annealing are the
recrystallization temperature or more.
[0329] The thus-obtained cold-rolled annealed sheet was examined with respect to the amount
of dissolved N, the microstructure, tensile properties, the r value, and strain age
hardenability by the same methods as Example 1. The tensile property of each cold-rolled
annealed sheet was measured at ten positions in each of the width direction and the
long direction to examine variations in yield strength, tensile strength and elongation.
The variation is shown by a difference between the maxim and minimum of all measurements,
for example, δYS = (maximum of YS ) - (minimum of YS). The results are shown in Table
20.
[0330] All the examples of the present invention have excellent strain age hardenability
and a high r value, and exhibit extremely high stable BH amount, ΔTS and mean r value
regardless of variations in production conditions. It was also recognized that in
the examples of the present invention, by performing continuous rolling and controlling
the temperature of the sheet bar in the long direction and the width direction, the
thickness precision and the shape of the produced steel sheet are improved, and variations
in material properties are decreased.
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