U.S. patent application number 10/578525 was filed with the patent office on 2008-05-22 for high-stiffness high-strength thin steel sheet and method for producing the same.
This patent application is currently assigned to JFE Steel Corporation. Invention is credited to Yoshihiro Hosoya, Taro Kizu, Kaneharu Okuda, Toshiaki Urabe, Hiromi Yoshida.
Application Number | 20080118390 10/578525 |
Document ID | / |
Family ID | 35063805 |
Filed Date | 2008-05-22 |
United States Patent
Application |
20080118390 |
Kind Code |
A1 |
Kizu; Taro ; et al. |
May 22, 2008 |
High-Stiffness High-Strength Thin Steel Sheet and Method For
Producing the Same
Abstract
There is provided a high-stiffness high-strength thin steel
sheet having a tensile strength of not less than 590 MPa and a
Young's modulus of not less than 225 GPa, which comprises C:
0.02-0.15%, Si: not more than 1.5%, Mn: 1.5-4.0%, P: not more than
0.05%, S: not more than 0.01%, Al: not more than 1.5%, N: not more
than 0.01% and Nb: 0.02-0.40% as mass %, provided that C, N and Nb
contents satisfy
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06 and
N.ltoreq.(14/92.9).times.(Nb-0.01) and the remainder being
substantially iron and inevitable impurities, and has a texture
comprising a ferrite phase as a main phase and having a martensite
phase at an area ratio of not less than 1%.
Inventors: |
Kizu; Taro; (Tokyo, JP)
; Okuda; Kaneharu; (Tokyo, JP) ; Urabe;
Toshiaki; (Tokyo, JP) ; Yoshida; Hiromi;
(Tokyo, JP) ; Hosoya; Yoshihiro; (Tokyo,
JP) |
Correspondence
Address: |
IP GROUP OF DLA PIPER US LLP
ONE LIBERTY PLACE, 1650 MARKET ST, SUITE 4900
PHILADELPHIA
PA
19103
US
|
Assignee: |
JFE Steel Corporation
Chiyoda-ku
JP
|
Family ID: |
35063805 |
Appl. No.: |
10/578525 |
Filed: |
March 31, 2005 |
PCT Filed: |
March 31, 2005 |
PCT NO: |
PCT/JP05/06288 |
371 Date: |
May 8, 2006 |
Current U.S.
Class: |
420/85 ;
420/84 |
Current CPC
Class: |
C21D 2211/008 20130101;
C21D 8/0205 20130101; C22C 38/02 20130101; C21D 9/46 20130101; C21D
2211/005 20130101; C22C 38/04 20130101 |
Class at
Publication: |
420/85 ;
420/84 |
International
Class: |
C22C 38/02 20060101
C22C038/02; C22C 38/04 20060101 C22C038/04; C22C 38/06 20060101
C22C038/06; C22C 38/14 20060101 C22C038/14 |
Foreign Application Data
Date |
Code |
Application Number |
Mar 31, 2004 |
JP |
2004-106,721 |
Nov 30, 2004 |
JP |
2004-347,025 |
Claims
1. A high-stiffness high-strength thin steel sheet comprising C:
0.02-0.15%, Si: not more than 1.5%, Mn: 1.5-4.0%, P: not more than
0.05%, S: not more than 0.01%, Al: not more than 1.5%, N: not more
than 0.01% and Nb: 0.02-0.40% as mass %, provided that C, N and Nb
contents satisfy the relationships of the following equations (1)
and (2):
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06 (1)
N.ltoreq.(14/92.9).times.(Nb-0.01) (2) and the remainder being
substantially iron and inevitable impurities, and having a texture
comprising a ferrite phase as a main phase and having a martensite
phase at an area ratio of not less than 1%, and having a tensile
strength of not less than 590 MPa and a Young's modulus of not less
than 225 GPa.
2. A high-stiffness high-strength thin steel sheet according to
claim 1, which further contains one or two of Ti: 0.01-0.50% and V:
0.01-0.50% as mass % in addition to the above composition and
satisfy the relationships of the following equations (3) and (4)
instead of the equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.T-
i*-(12/50.9).times.V.ltoreq.0.06 (3)
N*.ltoreq.(14/92.9).times.(Nb-0.01) (4) provided that N* in the
equations (3) and (4) is N*=N-(14/47.9).times.Ti at
N-(14/47.9).times.Ti>0 and N*=0 at
N-(14/47.9).times.Ti.ltoreq.0, and Ti* in the equation (3) is
Ti*=Ti-(47.9/14).times.N-(47.9/32.1).times.S at
Ti-(47.9/14.times.N-(47.9/32.1).times.S>0 and Ti*=0 at
Ti-(47.9/14).times.N-(47.9/32.1).times.S.ltoreq.0.
3. A high-stiffness high-strength thin steel sheet according to
claim 1 or 2, which further contains one or more of Cr: 0.1-1.0%,
Ni: 0.1-1.0%, Mo: 0.1-1.0%, Cu: 0.1-2.0% and B: 0.0005-0.0030% as
mass % in addition to the above composition.
4. A method for producing a high-stiffness high-strength thin steel
sheet comprising subjecting a starting material of steel comprising
C: 0.02-0.15%, Si: not more than 1.5%, Mn: 1.5-4.0%, P: not more
than 0.05%, S: not more than 0.01%, Al: not more than 1.5%, N: not
more than 0.01% and Nb: 0.02-0.40% as mass %, provided that C, N
and Nb contents satisfy the relationships of the following
equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06 (1)
N.ltoreq.(14/92.9).times.(Nb-0.01) (2) to a hot rolling step under
conditions that a total rolling reduction below 950.degree. C. is
not less than 30% and a finish rolling is terminated at
Ar.sub.3-900.degree. C., coiling the hot rolled sheet below
650.degree. C., pickling, subjecting to a cold rolling at a rolling
reduction of not less than 50%, raising a temperature to
780-900.degree. C. at a temperature rising rate from 500.degree. C.
of 1-40.degree. C./s to conduct soaking, and then cooling at a
cooling rate up to 500.degree. C. of not less than 5.degree. C./s
to conduct annealing.
5. A method for producing a high-stiffness high-strength thin steel
sheet according to claim 4, wherein the starting material of steel
further contains one or two of Ti: 0.01-0.50% and V: 0.01-0.50% as
mass % in addition to the above composition and satisfies the
relationships of the following equations (3) and (4) instead of the
equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.Ti*-(12-
/50.9).times.V.ltoreq.0.06 (3) N*.ltoreq.(14/92.9).times.(Nb-0.01)
(4) provided that N* in the equations (3) and (4) is
N*=N-(14/47.9).times.Ti at N-(14/47.9).times.Ti>0 and N*=0 at
N-(14/47.9).times.Ti.ltoreq.0, and Ti* in the equation (3) is
Ti*=Ti-(47.9/14).times.N-(47.9/32.1).times.S at
Ti-(47.9/14).times.N-(47.9/32.1).times.S>0 and Ti*=0 at
Ti-(47.9/14).times.N-(47.9/32.1).times.S.ltoreq.0.
6. A method for producing a high-stiffness high-strength thin steel
sheet according to claim 4 or 5, wherein the staring material of
steel further contains one or more of Cr: 0.1-1.0%, Ni: 0.1-1.0%,
Mo: 0.1-1.0%, Cu: 0.1-2.0% and B: 0.0005-0.0030% as mass % in
addition to the above composition.
Description
TECHNICAL FIELD
[0001] This invention relates to a high-stiffness high-strength
thin steel sheet suitable mainly as a vehicle body for automobiles
and a method for producing the same. Moreover, the high-stiffness
high-strength thin steel sheet according to the invention is a
column-shaped structural member having a thickness susceptibility
index of the stiffness near to 1 such as a center pillar, locker,
side flame, cross member or the like of the automobile and is
widely suitable for applications requiring a stiffness.
RELATED ART
[0002] As a result of recent heightened interest in global
environment problems, the exhaust emission control is conducted
even in the automobiles, and hence the weight reduction of the
vehicle body in the automobile is a very important matter. For this
end, it is effective to attain the weight reduction of the vehicle
body by increasing the strength of the steel sheet to reduce the
thickness thereof.
[0003] Recently, the increase of the strength in the steel sheet is
considerably advanced, and hence the use of thin steel sheets
having a thickness of less than 2.0 mm is increasing. In order to
further reduce the weight by the increase of the strength, it is
indispensable to simultaneously control the deterioration of the
stiffness in parts through the thinning of the thickness. Such a
problem of deteriorating the stiffness of the parts through the
thinning of the thickness in the steel sheet is actualized in steel
sheets having a tensile strength of not less than 590 MPa, and
particularly this problem is serious in steel sheets having a
tensile strength of not less than 700 MPa.
[0004] In general, in order to increase the stiffness of the parts,
it is effective to change the shape of the parts, or to increase
the number of welding points or change the welding condition such
as changeover to laser welding or the like in the spot-welded
parts. However, when these parts are used in the automobile, there
are problems that it is not easy to change the shape of the parts
in a limited space inside the automobile, and the change of the
welding conditions causes the increase of the cost and the
like.
[0005] Consequently, in order to increase the stiffness of the
parts without changing the shape of the parts or the welding
conditions, it becomes effective to increase the Young's modulus of
the material used in the parts.
[0006] In general, the stiffness of the parts under the same shape
of parts and welding conditions is represented by a product of
Young's modulus of the material and geometrical moment of inertia
of the part. Further, the geometrical moment of inertia can be
expressed so as to be approximately proportionate to t.sup..lamda.
when the thickness of the material is t. In this case, .lamda. is a
thickness susceptibility index and is a value of 1-3 in accordance
with the shape of the parts. For example, in case of one plate
shape such as panel parts for the automobile, .lamda. is a value
near to 3, while in case of column-shape such as structural parts,
.lamda. is a value near to 1.
[0007] When .lamda. of the parts is 3, if the thickness is made
small by 10% while equivalently maintaining the stiffness of the
parts, it is required to increase the Young's modulus of the
material by 37%, while when .lamda. of the parts is 1, if the
thickness is made small by 10%, it may be enough to increase the
Young's modulus by 11%.
[0008] That is, in case of the parts having .lamda. near to 1 such
as column-shaped parts, it is very effective to increase the
Young's modulus of the steel sheet itself for the weight reduction.
Particularly, in case of steel sheets having a high strength and a
small thickness, it is strongly demanded to highly increase the
Young's modulus of the steel sheet.
[0009] In general, the Young's modulus is largely dependent upon
the texture and is known to become high in a closest direction of
atom. Therefore, it is effective to develop {112}<110> in
order to develop an orientation advantageous for the Young's
modulus of steel being a body-centered cubic lattice in a steel
making process comprising the rolling through rolls and the heat
treatment, whereby the Young's modulus can be increased in a
direction perpendicular to the rolling direction.
[0010] There have hitherto been variously examined steel sheets by
controlling the texture to increase the Young's modulus.
[0011] For example, the patent article 1 discloses a technique
wherein a steel obtained by adding Nb or Ti to an extremely low
carbon steel is hot-rolled at a rolling reduction at
Ar.sub.3-Ar.sub.3+150.degree. C.) of not less than 85% to promote
transformation from non-crystallized austenite to ferrite to
thereby render the texture of ferrite at the stage of the
hot-rolled sheet into {311}<011> and {332}<113>, which
is an initial orientation and is subjected to a cold rolling and a
recrystallization annealing to render {211}<011> into a main
orientation to thereby increase the Young's modulus in a direction
perpendicular to the rolling direction.
[0012] Also, the patent article 2 discloses a method for producing
a hot rolled steel sheet having an increased Young's modulus in
which Nb, Mo and B are added to a low carbon steel having a C
content of 0.02-0.15% and the rolling reduction at
Ar.sub.3-950.degree. C. is made to not less than 50% to develop
[211]<011>.
[0013] Further, the patent article 3 discloses a method for
producing. a hot rolled steel sheet having a high stiffness in
which Nb is added to a low carbon steel having a C content of not
more than 0.05% and a finish rolling start temperature is made to
not higher than 950.degree. C. and a finish rolling end temperature
is made to (Ar.sub.3-50.degree. C.)-(Ar.sub.3+100.degree. C.) to
control the development of {100} decreasing the Young's
modulus.
[0014] Moreover, the patent article 4 discloses a method for
producing a hot rolled steel sheet in which Si and Al are added to
a low carbon steel having a C content of not more than 0.05% to
enhance Ar.sub.3 transformation point and the rolling reduction
below Ar.sub.3 transformation point in the hot rolling is made to
not less than 60% to increase Young's modulus in a direction
perpendicular to the rolling direction.
[0015] Patent article 1: JP-A-H05-255804
[0016] Patent article 2: JP-A-H08-311541
[0017] Patent Article 3: JP-A-H05-247530
[0018] Patent article 4: JP-A-H09-53118
DISCLOSURE OF THE INVENTION
Problems to be Solved in the Invention
[0019] However, the aforementioned techniques have the following
problems.
[0020] In the technique disclosed in the patent article 1, the
Young's modulus of the steel sheet is increased by using the
extremely low carbon steel having a C content of not more than
0.01% to control the texture, but the tensile strength is low as
about 450 MPa at most, so that there is a problem in the increase
of the strength by applying this technique.
[0021] In the technique disclosed in the patent article 2, since
the C content is as high as 0.02-0.15%, it is possible to increase
the strength, but as the target steel sheet is the hot rolled steel
sheet, the control of the texture through cold working can not be
utilized, and hence there are problems that it is difficult to
further increase the Young's modulus but also it is difficult to
stably produce high-strength steel sheets having a thickness of
less than 2.0 mm through low-temperature finish rolling.
[0022] Also, the technique disclosed in the patent article 3 is the
production of the hot rolled steel sheet, so that it has the same
problems as mentioned above.
[0023] Further, in the technique disclosed in the patent article 4,
the crystal grains are coarsened by conducting the rolling at the
ferrite zone, so that there is a problem that the workability is
considerably deteriorated.
[0024] Thus, the increase of the Young's modulus in the steel sheet
by the conventional techniques is targeted to hot rolled steel
sheets having a thick thickness or soft steel sheets, so that it is
difficult to increase the Young's modulus of high-strength thin
steel sheet having a thickness of not more than 2.0 mm by using the
above conventional techniques.
[0025] As a strengthening mechanism for increasing the tensile
strength of the steel sheet to not less than 590 MPa, there are
mainly a precipitation strengthening mechanism and a transformation
texture strengthening mechanism.
[0026] When the precipitation strengthening mechanism is used as
the strengthening mechanism, it is possible to increase the
strength while suppressing the lowering of the Young's modulus of
the steel sheet as far as possible, but the following difficulty is
accompanied. That is, when utilizing the precipitation
strengthening mechanism for finely precipitating, for example, a
carbonitride of Ti, Nb or the like, in the hot rolled steel sheet,
the increase of the strength is attained by conducting the fine
precipitation in the coiling after the hot rolling, but in the cold
rolled steel sheet, the coarsening of the precipitate can not be
avoided at the step of recrystallization annealing after the cold
rolling and it is difficult to increase the strength through the
precipitation strengthening.
[0027] When utilizing the transformation texture strengthening
mechanism as the strengthening mechanism, there is a problem that
the Young's modulus of the steel sheet lowers due to strain
included in a low-temperature transformation phase such as bainite
phase, martensite phase or the like.
[0028] It is, therefore, an object of the invention to solve the
above problems and to provide a high-stiffness high-strength thin
steel sheet having a tensile strength of not less than 590 MPa,
preferably not less than 700 MPa, a Young's modulus of not less
than 225 GPa, preferably not less than 230 GPa, more preferably not
less than 240 GPa and a thickness of not more than 2.0 mm as well
as an advantageous method for producing the same.
Means for Solving Problems
[0029] In order to achieve the above object, the gist and
construction of the invention are as follows.
[0030] (I) A high-stiffness high-strength thin steel sheet
comprising C: 0.02-0.15%, Si: not more than 1.5%, Mn: 1.5-4.0%, P:
not more than 0.05%, S: not more than 0.01%, Al: not more than
1.5%, N: not more than 0.01% and Nb: 0.02-0.40% as mass %, provided
that C, N and Nb contents satisfy the relationships of the
following equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06
(1)
N.ltoreq.(14/92.9).times.(Nb-0.01) (2)
and the remainder being substantially iron and inevitable
impurities, and having a texture comprising a ferrite phase as a
main phase and having a martensite phase at an area ratio of not
less than 1%, and having a tensile strength of not less than 590
MPa and a Young's modulus of not less than 225 GPa.
[0031] (II) A high-stiffness high-strength thin steel sheet
according to the item (I), which further contains one or two of Ti:
0.01-0.50% and V: 0.01-0.50% as mass % in addition to the above
composition and satisfy the relationships of the following
equations (3) and (4) instead of the equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.Ti*-(1-
2/50.9).times.V.ltoreq.0.06 (3)
N*.ltoreq.(14/92.9).times.(Nb-0.01) (4)
provided that N* in the equations (3) and (4) is
N*=N-(14/47.9).times.Ti at N-(14/47.9).times.Ti>0 and N*=0 at
N-(14/47.9).times.Ti.ltoreq.0, and Ti* in the equation (3) is
Ti*=Ti-(47.9/14).times.N-(47.9/32.1).times.S at
Ti-(47.9/14).times.N-(47.9/32.1).times.S>0 and Ti*=0 at
Ti-(47.9/14).times.N-(47.9/32.1).times.S.ltoreq.0.
[0032] (III) A high-stiffness high-strength thin steel sheet
according to the item (I) or (II), which further contains one or
more of Cr: 0.1-1.0%, Ni: 0.1-1.0%, Mo: 0.1-1.0%, Cu: 0.1-2.0% and
B: 0.0005-0.0030% as mass % in addition to the above
composition.
[0033] (IV) A method for producing a high-stiffness high-strength
thin steel sheet comprising subjecting a starting material of steel
comprising C: 0.02-0.15%, Si: not more than 1.5%, Mn: 1.5-4.0%, P:
not more than 0.05%, S: not more than 0.01%, Al: not more than
1.5%, N: not more than 0.01% and Nb: 0.02-0.40% as mass %, provided
that C, N and Nb contents satisfy the relationships of the
following equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06
(1)
N.ltoreq.(14/92.9).times.(Nb-0.01) (2)
to a hot rolling step under conditions that a total rolling
reduction below 950.degree. C. is not less than 30% and a finish
rolling is terminated at Ar.sub.3-900.degree. C., coiling the hot
rolled sheet below 650.degree. C., pickling, subjecting to a cold
rolling at a rolling reduction of not less than 50%, raising a
temperature to 780-900.degree. C. at a temperature rising rate from
500.degree. C. of 1-40.degree. C./s to conduct soaking, and then
cooling at a cooling rate up to 500.degree. C. of not less than
5.degree. C./s to conduct annealing.
[0034] (V) A method for producing a high-stiffness high-strength
thin steel sheet according to the item (IV), wherein the starting
material of steel further contains one or two of Ti: 0.01-0.50% and
V: 0.01-0.50% as mass % in addition to the above composition and
satisfies the relationships of the following equations (3) and (4)
instead of the equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.Ti*-(1-
2/50.9).times.V.ltoreq.0.06 (3)
N*.ltoreq.(14/92.9).times.(Nb-0.01) (4)
provided that N* in the equations (3) and (4) is
N*=N-(14/47.9).times.Ti at N-(14/47.9).times.Ti>0 and N*=0 at
N-(14/47.9).times.Ti.ltoreq.0, and Ti* in the equation (3) is
Ti*=Ti-(47.9/14).times.N-(47.9/32.1).times.S at
Ti-(47.9/14).times.N-(47.9/32.1).times.S>0 and Ti*=0 at
Ti-(47.9/14).times.N-(47.9/32.1).times.S.ltoreq.0.
[0035] (VI) A method for producing a high-stiffness high-strength
thin steel sheet according to the item (IV) or (V), wherein the
staring material of steel further contains one or more of Cr:
0.1-1.0%, Ni: 0.1-1.0%, Mo: 0.1-1.0%, Cu: 0.1-2.0% and B:
0.0005-0.0030% as mass % in addition to the above composition.
Effect of the Invention
[0036] According to the invention, it is possible to provide a
high-stiffness high-strength thin steel sheet having a tensile
strength of not less than 590 MPa, preferably not less than 700 MPa
and a Young's modulus of not less than 225 GPa, preferably not less
than 230 GPa, more preferably not less than 240 GPa.
[0037] That is, the starting material of low carbon steel added
with Mn and Nb is roll-reduced below 950.degree. C., preferably
below 900.degree. C. (strictly speaking, just above Ar.sub.3 point)
in the hot rolling to promote the transformation from
non-recrystallized austenite to ferrite and then cold rolled to
develop a crystal orientation useful for the improvement of Young's
modulus and thereafter a low-temperature transformation phase
suppressing the lowering of the Young's modulus is produced and a
greater amount of ferrite phase useful for the improvement of the
Young's modulus is retained in the cooling stage by the control of
the heating rate in the annealing step and the soaking at two-phase
region, whereby the thin steel sheet satisfying higher strength and
higher Young's modulus can be produced, which develops an effective
effect in industry.
[0038] Further explaining in detail, the starting material of low
carbon steel added with Mn and Nb is roll-reduced just above
Ar.sub.3 transformation point in the hot rolling to increase the
non-recrystallized austenite texture having a crystal orientation
of {112}<111>, and subsequently the transformation from the
non-recrystallized austenite of {112}<111> to ferrite is
promoted in the cooling stage to develop ferrite orientation of
{113}<110>.
[0039] In the cold rolling after the coiling and pickling, the
rolling is carried out at a rolling reduction of not less than 50%
to turn the crystal orientation of {113}<110> to
{112}<110> useful for the improvement of the Young's modulus,
and in the temperature rising stage at the subsequent annealing
step, the temperature is raised from 500.degree. C. to the soaking
temperature at a heating rate of 1-40.degree. C./s to promote the
recrystallization of ferrite having an orientation of
{112}<110> and provide a two-phase region at a state of
partly retaining the non-recrystallized grains of {112}<110>,
whereby the transformation from the non-recrystallized ferrite of
{112}<110> to austenite can be promoted.
[0040] Further, in the transformation from austenite phase to
ferrite phase at the cooling after the soaking, ferrite grains
having an orientation of {112}<110> is grown to enhance the
Young's modulus, while the steel enhancing the hardenability by the
addition of Mn is cooled at a rate of not less than 5.degree. C./s
to produce the low-temperature transformation phase, whereby it is
attempted to increase the strength.
[0041] Moreover, the low-temperature transformation phase is
produced by retransforming the austenite phase transformed from
ferrite having an orientation of {112}<110> during the
cooling, so that {112}<110> can be also developed even in the
crystal orientation of the low-temperature transformation
phase.
[0042] Thus, the Young's modulus is enhanced by developing
{112}<110> of ferrite phase, and particularly
{112}<110> is increased in the orientation of the
low-temperature transformation phase largely exerting on the
lowering of the Young's modulus, whereby the strength can be
increased by the formation of the low-temperature transformation
phase and the lowering of the Young's modulus accompanied with the
formation of the low-temperature transformation phase can be
largely suppressed.
BRIEF DESCRIPTION OF THE DRAWINGS
[0043] FIG. 1 is a graph showing an influence of a total rolling
reduction below 950.degree. C. or below 900.degree. C. on Young's
modulus;
[0044] FIG. 2 is a graph showing an influence of a final
temperature in hot finish rolling on Young's modulus;
[0045] FIG. 3 is a graph showing an influence of a coiling
temperature on Young's modulus;
[0046] FIG. 4 is a graph showing an influence of a rolling
reduction in cold rolling on Young's modulus; and
[0047] FIG. 5 is a graph showing an influence of an average
temperature rising rate from 500.degree. C. to soaking temperature
in annealing on Young's modulus.
BEST MODE FOR CARRYING OUT THE INVENTION
[0048] The high-stiffness high-strength thin steel sheet according
to the invention is a steel sheet having a tensile strength of not
less than 590 MPa, preferably not less than 700 MPa, a Young's
modulus of not less than 225 GPa, preferably not less than 230 GPa,
more preferably not less than 240 GPa, and a thickness of not more
than 2.0 mm. Moreover, the steel sheet to be targeted in the
invention includes steel sheets subjected to a surface treatment
such as galvanization inclusive of alloying, zinc electroplating or
the like in addition to the cold rolled steel sheet.
[0049] The reason of limiting the chemical composition in the steel
sheet of the invention will-be described below. Moreover, the unit
for the content of each element in the chemical composition of the
steel sheet is "% by mass", but it is simply shown by "%" unless
otherwise specified.
[0050] C: 0.02-0.15%
[0051] C is an element stabilizing austenite and can largely
contribute to increase the strength by enhancing the hardenability
at the cooling stage in the annealing after the cold rolling to
largely promote the formation of the low-temperature transformation
phase. Further, the Ar.sub.3 transformation point is lowered in the
hot rolling and it is possible to conduct the rolling at a lower
temperature region when the rolling is conducted just above
Ar.sub.3, whereby the transformation from the non-recrystallized
austenite to ferrite can be promoted to develop {113}<110>,
and the Young's modulus can be improved at the subsequent cold
rolling and annealing steps. Moreover, C can contribute to increase
the Young's modulus by promoting the transformation of ferrite
grains having {112}<110> from the non-recrystallized ferrite
to austenite after the cold rolling.
[0052] In order to obtain such effects, the C content is required
to be not less than 0.02%, preferably not less than 0.05%, more
preferably not less than 0.06%. On the other hand, when the C
content exceeds 0.15%, the fraction of hard low-temperature
transformation phase becomes large, and the strength of the steel
is extremely increased but also the workability is deteriorated.
Also, the greater amount of C suppresses the recrystallization of
the orientation useful for the increase of the Young's modulus at
the annealing step after the cold rolling. Further, the greater
amount of C brings about the deterioration of the weldability.
[0053] Therefore, the C content is required to be not more than
0.15%, preferably not more than 0.10%.
[0054] Si: not more than 1.5%
[0055] Si raises the Ar.sub.3 transformation point in the hot
rolling, so that when the rolling is carried out just above
Ar.sub.3, the recrystallization of worked austenite is promoted.
Therefore, when Si is contained in an amount exceeding 1.5%, the
crystal orientation required for the increase of the Young's
modulus can not be obtained. Also, the greater amount of Si
deteriorates the weldability of the steel sheet but also promotes
the formation of fayalite on a surface of a slab in the heating at
the hot rolling step to accelerate the occurrence of surface
pattern so-called as a red scale. Furthermore, in case of using as
a cold rolled steel sheet, Si oxide produced on the surface
deteriorates the chemical conversion processability, while in case
of using as a galvanized steel sheet, Si oxide produced on the
surface induces non-plating. Therefore, the Si content is required
to be not more than 1.5%. Moreover, in case of steel sheets
requiring the surface properties or the galvanized steel sheet, the
Si content is preferable to be not more than 0.5%.
[0056] Also, Si is an element stabilizing ferrite and promotes the
ferrite transformation at the cooling stage after the soaking of
two-phase region in the annealing step after the cold rolling to
enrich C in austenite, whereby austenite can be stabilized to
promote the formation of the low-temperature transformation phase.
For this end, the strength of steel can be increased, if necessary.
In order to obtain such an effect, the Si content is desirable to
be not less than 0.2%.
[0057] Mn: 1.5-4.0%
[0058] Mn is one of important elements in the invention. Mn is an
element suppressing the recrystallization of worked austenite in
the hot rolling and stabilizing austenite, and since Mn lowers the
Ar.sub.3 transformation point, when the rolling is carried out just
above Ar.sub.3, it is possible to conduct the rolling at a lower
temperature region, and further Mn has an action of suppressing the
recrystallization of the worked austenite. Moreover, Mn can promote
the transformation from the non-recrystallized austenite to ferrite
to develop {113}<110> and improve the Young's modulus in the
subsequent cold rolling and annealing steps.
[0059] Furthermore, Mn as an austenite stabilizing element lowers
Ac.sub.1 transformation point in the temperature rising stage at
the annealing step after the cold rolling to promote the
transformation from the non-recrystallized ferrite to austenite,
and can develop the orientation useful for the improvement of the
Young's modulus to control the lowering of the Young's modulus
accompanied with the formation of the low-temperature
transformation phase with respect to the orientation of the
low-temperature transformation phase produced in the cooling stage
after the soaking.
[0060] Also, Mn enhances the hardenability in the cooling stage
after the soaking and annealing at the annealing step to largely
promote the formation of the low-temperature transformation phase,
which can largely contribute to the increase of the strength.
Further, Mn acts as a solid-solution strengthening element, which
can contribute to the increase of the strength in steel. In order
to obtain such an effect, the Mn content is required to be not less
than 1.5%.
[0061] On the other hand, when the Mn content exceeds 4.0%,
Ac.sub.3 transformation point is excessively lowered in the
temperature rising stage at the annealing step after the cold
rolling, so that the recrystallization of ferrite phase at the
two-phase region is difficult and it is required to raise the
temperature up to an austenite single-phase region above Ac.sub.3
transformation point. As a result, ferrite of {112}<110>
orientation useful for the increase of the Young's modulus obtained
by the recrystallization of worked ferrite can not be developed to
bring about the lowering of the Young's modulus. Further, the
greater amount of Mn deteriorates the weldability of the steel
sheet. Therefore, the Mn content is not more than 4.0%, preferably
not more than 3.5%.
[0062] P: not more than 0.05%
[0063] Since P segregates in the grain boundary, if the P content
exceeds 0.05%, the ductility and toughness of the steel sheet lower
but also the weldability is deteriorated. In case of using the
alloyed galvanized steel sheet, the alloying rate is delayed by P.
Therefore, the P content is required to be not more than 0.05%. On
the other hand, P is an element effective for the increase of the
strength as a solid-solution strengthening element and has an
action of promoting the enrichment of C in austenite as a ferrite
stabilizing element. In the steel added with Si, it has also an
action of suppressing the occurrence of red scale. In order to
obtain these actions, the P content is preferable to be not less
than 0.01%.
[0064] S: not more than 0.01%
[0065] S considerably lowers the hot ductility to induce hot
tearing and considerably deteriorate the surface properties.
Further, S hardly contributes to the strength but also forms coarse
MnS as an impurity element to lower the ductility and
drill-spreading property. These problems become remarkable when the
S content exceeds 0.01%, so that it is desirable to reduce the S
content as far as possible. Therefore, the S content is not more
than 0.01%. From a viewpoint of improving the drill-spreading
property, it is preferable to be not more than 0.005%.
[0066] Al: not more than 1.5%
[0067] It is an element useful for deoxidizing steel to improve the
cleanness of the steel. However, Al is a ferrite stabilizing
element, and largely raises the Ar.sub.3 transformation of the
steel, so that when the rolling is carried out just above Ar.sub.3,
the recrystallization of worked austenite is promoted to suppress
the development of the crystal orientation required for the
increase of the Young's modulus. Further, when the Al content
exceeds 1.5%, the austenite single-phase region disappears and it
is difficult to terminate the rolling at austenite region in the
hot rolling step. Therefore, the Al content is required to be not
more than 1.5%. From this viewpoint, Al is preferable to be made
lower, and further preferable to be limited to not more than 0.1%.
On the other hand, Al as a ferrite forming element promotes the
formation of ferrite in the cooling stage after the soaking at the
two-phase region in the annealing step after the cold rolling to
enrich C in austenite, whereby austenite can be stabilized to
promote the formation of the low-temperature transformation phase.
As a result, the strength of the steel can be enhanced, if
necessary. In order to obtain such an effect, the Al content is
desirable to be not less than 0.2%.
[0068] N: not more than 0.01%
[0069] N is a harmful element because slab breakage is accompanied
in the hot rolling to cause surface defect. When the N content
exceeds 0.01%, the occurrence of slab breakage and surface defect
becomes remarkable. Therefore, the N content is required to be not
more than 0.01%.
[0070] Nb: 0.02-0.40%
[0071] Nb is a most important element in the invention. That is, Nb
suppresses the recrystallization of worked austenite at the finish
rolling step in the hot rolling to promote the transformation from
the non-recrystallized austenite to ferrite and develop
{113}<110> and can increase the Young's modulus at the
subsequent cold rolling and annealing steps. Also, the
recrystallization of worked ferrite is suppressed at the
temperature rising stage in the annealing step after the cold
rolling to promote the transformation from the non-recrystallized
ferrite to austenite. As to the orientation of the low-temperature
transformation phase produced in the cooling stage after the
soaking, the orientation useful for the increase of the Young's
modulus can be developed to suppress the lowering of the Young's
modulus accompanied with the formation of the low-temperature
transformation phase. Also, a fine carbonitride of Nb can
contribute to the increase of the strength. In order to obtain such
an action, the Nb content is required to be not less than 0.02%,
preferably not less than 0.05%.
[0072] On the other hand, when the Nb content exceeds 0.40%, the
all carbonitride can not be solid-soluted in the re-heating at the
usual hot rolling step and hence coarse carbonitride remains, so
that the effect of suppressing the recrystallization of worked
austenite in the hot rolling step and the effect of suppressing the
recrystallization of worked ferrite in the annealing step after the
cold rolling can not be obtained. Also, even if the hot rolling of
the slab after the continuous casting is started as it is without
conducting the re-heating after the continuously cast slab is
cooled, when Nb is included in an amount exceeding 0.40%, the
improvement of the effect of suppressing the recrystallization is
not recognized and the increase of the alloy cost is caused.
Therefore, the Nb content is 0.02-0.40%, preferably 0.05-0.40%.
[0073] In the invention, the contents of C, N and Nb are required
to satisfy the relationship of the following equations (1) and
(2):
0.01.ltoreq.C+(12/14).times.N-(12/92.9).times.Nb.ltoreq.0.06
(1)
N.ltoreq.(14/92.9).times.(Nb-0.01) (2)
[0074] If C not fixed as a carbonitride is existent in an amount
exceeding 0.06%, the introduction of strain in the cold rolling
becomes non-uniform and further the recrystallization of the
orientation useful for the increase of the Young's modulus is
suppressed, so that the C amount not fixed as the carbonitride
calculated by (C+(12/14).times.N-(12/92.9).times.Nb) is required to
be not more than 0.06%, preferably not more than 0.05%. At this
moment, N is preferentially fixed and precipitated as compared with
C, so that the C amount not fixed as the carbonitride can be
calculated by (C+(12/14).times.N-(12/92.9).times.Nb). On the other
hand, when the C amount not fixed as the carbonitride is less than
0.01%, the C content in austenite decreases in the annealing at the
two-phase region after the cold rolling and the formation of
martensite phase after the cooling is suppressed, so that it is
difficult to increase the strength of the steel. Therefore, the
amount of (C+(12/14).times.N-(12/92.9).times.Nb), which is the C
amount not fixed as the carbonitride, is 0.01-0.06%, preferably
0.01-0.05%. Further, N coarsely precipitates a nitride of Nb at a
high temperature, and hence the effect of suppressing the
recrystallization by Nb is reduced. In order to control this
action, the N content is required to be limited to
N.ltoreq.(14/92.9).times.(Nb-0.01) in relation with the Nb content,
preferably N.ltoreq.(14/92.9).ltoreq.(Nb-0.02).
[0075] Moreover, the term "the remainder being substantially iron
and inevitable impurities" used herein means that steels containing
slight amounts of other elements without damaging the action and
effect of the invention are included within the scope of the
invention. In case of further increasing the strength, one or two
of Ti and V and one or more of Cr, Ni, Mo, Cu and B may be added,
if necessary, in addition to the above definition of the chemical
composition.
[0076] Ti: 0.01-0.50%
[0077] Ti is an element contributing to the increase of the
strength by forming a fine carbonitride. Also, it is an element
contributing to the increase of the Young's modulus by suppressing
the recrystallization of worked austenite in the finish rolling
step of the hot rolling to promote the transformation from the
non-recrystallized austenite to ferrite. Since Ti has the above
actions, the content is preferable to be not less than 0.01%. On
the other hand, when the Ti content exceeds 0.50%, all the
carbonitride can not be solid-soluted in the re-heating at the
usual hot rolling step and a coarse carbonitride remains, and hence
the effect of increasing the strength and the effect of suppressing
the recrystallization can not be obtained. Also, even if the hot
rolling of the slab after the continuous casting is started as it
is without conducting the re-heating after the continuously cast
slab is cooled, the Ti content exceeding 0.50% is small in the
contribution to the effect of increasing the strength and the
effect of suppressing the recrystallization and also the increase
of the alloy cost is caused. Therefore, the Ti content is
preferably not more than 0.50%, more preferably not more than
0.20%.
[0078] V: 0.01-0.50%
[0079] V is an element contributing to the increase of the strength
by forming a fine carbonitride. Since V has such an action, the V
content is preferable to be not less than 0.01%. On the other hand,
when the V content exceeds 0.50%, the effect of increasing the
strength by the amount exceeding 0.50% is small and the increase of
the alloy cost is caused. Therefore, the V content is preferably
not more than 0.50%, more preferably not more than 0.20%.
[0080] In the invention, when Ti and/or V are included in addition
to Nb, the contents of C, N, S, Nb, Ti and V are required to
satisfy the relationship of the following equations (3) and (4)
instead of the equations (1) and (2):
0.01.ltoreq.C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.Ti*-(1-
2/50.9).times.V.ltoreq.0.06 (3)
N*.ltoreq.(14/92.9).times.(Nb-0.01) (4)
provided that N* in the equations (3) and (4) is
N*=N-(14/47.9).times.Ti at N-(14/47.9).times.Ti>0 and N*=0 at
N-(14/47.9).times.Ti.ltoreq.0, and Ti* in the equation (3) is
Ti*=Ti-(47.9/14).times.N-(47.9/32.1).times.S at
Ti-(47.9/14).times.N-(47.9/32.1).times.S>0 and Ti*=0 at
Ti-(47.9/14).times.N-(47.9/32.1).times.S.ltoreq.0.
[0081] Further, N coarsely precipitates the nitride of Nb at a high
temperature as previously mentioned, so that the effect of
suppressing the recrystallization through Nb is decreased. In case
of Ti-containing steel, N is preferentially fixed as a nitride of
Ti, N* as a N amount not fixed as a nitride of Ti is required to be
limited to N* s (14/92.9).times.(Nb-0.01), preferably
N*.ltoreq.(14/92.9).times.((Nb-0.02).
[0082] Ti and V form the carbonitride to decrease the C content not
fixed as the carbonitride. Further, Ti is fixed by the formation of
a sulfide, so that the value of
C+(12/14).times.N*-(12/92.9).times.Nb-(12/47.9).times.Ti*-(12/50.9).times-
.V is required to be 0.01-0.06%, preferably 0.01-0.05% when Ti
and/or V are added in order that the C content not fixed as the
carbonitride is made to 0.01-0.06%.
[0083] Cr: 0.1-1.0%
[0084] Cr is an element enhancing the hardenability by suppressing
the formation of cementite and can largely contribute to the
increase of the strength by largely promoting the formation of the
low-temperature transformation phase in the cooling stage after the
soaking at the annealing step. Further, the recrystallization of
worked austenite is suppressed in the hot rolling step to promote
the transformation from non-recrystallized austenite to ferrite and
develop {113}<110>, and the Young's modulus can be increased
at the subsequent cold rolling and annealing steps. In order to
obtain such an effect, Cr is preferable to be included in an amount
of not less than 0.1%. On the other hand, when the Cr content
exceeds 1.0%, the above effect is saturated and the alloy cost
increases, so that Cr is preferable to be included in an amount of
not more than 1.0%. Moreover, when the thin steel sheet of the
invention is used as a galvanized steel sheet, the oxide of Cr
produced on the surface induces the non-plating, so that Cr is
preferable to be included in an amount of not more than 0.5%.
[0085] Ni: 0.1-1.0%
[0086] Ni is an element stabilizing austenite to enhance the
hardenability, and can largely contribute to the increase of the
strength by largely promoting the formation of the low-temperature
transformation phase in the cooling stage after the soaking at the
annealing step. Further, Ni as an austenite stabilizing element
lowers Ac.sub.1 transformation point in the temperature rising
stage at the annealing step after the cold rolling to promote the
transformation from the non-recrystallized ferrite to austenite,
and develops the orientation useful for the increase of the Young's
modulus with respect to the orientation of the low-temperature
transformation phase produced in the cooling stage after the
soaking, whereby the lowering of the Young's modulus accompanied
with the formation of the low-temperature transformation phase can
be suppressed. Since Ni is an element suppressing the
recrystallization of worked austenite in the hot rolling and
stabilizing austenite, when Ar.sub.3 transformation point is
lowered to conduct the rolling just above Ar.sub.3, it is possible
to conduct the rolling at a lower temperature region to further
suppress the recrystallization of worked austenite, and also the
transformation from the non-recrystallized austenite to ferrite is
promoted to develop {113}<110>, whereby the Young's modulus
can be increased at the subsequent cold rolling and annealing
steps. In case of adding Cu, the surface defect is induced by
cracking accompanied with the lowering of the hot ductility in the
hot rolling, but the occurrence of the surface defect can be
controlled by composite addition of Ni. In order to obtain such an
action, Ni is preferable to be included in an amount of not less
than 0.1%.
[0087] On the other hand, when the Ni content exceeds 1.0%,
Ac.sub.3 transformation point is extremely lowered in the
temperature rising stage at the annealing step after the cold
rolling and the recrystallization of ferrite phase at the two-phase
region is difficult, and hence it is required to raise the
temperature up to austenite single phase region above Ac.sub.3
transformation point. As a result, ferrite of orientation obtained
by the recrystallization of worked ferrite and useful for the
increase of the Young's modulus can not be developed to bring about
the decrease of the Young's modulus. And also, the alloy cost
increases. Therefore, Ni is preferable to be included in an amount
of not more than 1.0%.
[0088] Mo: 0.1-1.0%
[0089] Mo is an element enhancing the hardenability by making small
the mobility of the interface, and can largely contribute to the
increase of the strength by largely promoting the formation of the
low-temperature transformation phase in the cooling stage at the
annealing step after the cold rolling. Further, the
recrystallization of worked austenite can be suppressed, and the
transformation from the non-recrystallized austenite to ferrite is
promoted to develop {113}<110> and the Young's modulus can be
increased at the subsequent cold rolling and annealing steps. In
order to obtain such an action, Mo is preferable to be included in
an amount of not less than 0.1%. On the other hand, when the Mo
content exceeds 1.0%, the above effect is saturated and the alloy
cost increases, so that Mo is preferable to be included in an
amount of not more than 1.0%.
[0090] B: 0.0005-0.0030%
[0091] B is an element suppressing the transformation from
austenite phase to ferrite phase to enhance the hardenability, and
can largely contribute to the increase of the strength by largely
promoting the formation of the low-temperature transformation phase
in the cooling stage at the annealing step after the cold rolling.
Further, the recrystallization of worked austenite can be
suppressed, and the transformation from the non-recrystallized
austenite to ferrite is promoted to develop {113}<110> and
the Young's modulus can be increased at the subsequent cold rolling
and annealing steps. In order to obtain such an effect, B is
preferable to be included in an amount of not less than 0.0005%. On
the other hand, when the B content exceeds 0.0030%, the above
effect is saturated, so that B is preferable to be included in an
amount of not more than 0.0030%.
[0092] Cu: 0.1-2.0%
[0093] Cu is an element enhancing the hardenability, and can
largely contribute to the increase of the strength by largely
promoting the formation of the low-temperature transformation phase
in the cooling stage at the annealing step after the cold rolling.
In order to obtain such an effect, Cu is preferable to be included
in an amount of not less than 0.1%. On the other hand, when the Cu
content exceeds 2.0%, the hot ductility is lowered and the surface
defect accompanied with the cracking in the hot rolling is induced
and the hardening effect by Cu is saturated, so that Cu is
preferable to be included in an amount of not more than 2.0%.
[0094] The reason on the limitation of the texture according to the
invention will be described below.
[0095] In the thin steel sheet of the invention, it is required to
have a texture comprising a ferrite phase as a main phase and
having a martensite phase at an area ratio of not less than 1%.
[0096] The term "ferrite phase as a main phase" used herein means
that the area ratio of the ferrite phase is not less than 50%.
[0097] Since the ferrite phase is less in the strain, useful for
the increase of the Young's modulus, excellent in the ductility and
good in the workability, the texture is required to be the ferrite
phase as a main phase.
[0098] Also, in order to render the tensile strength of the steel
sheet into not less than 590 MPa, it is required that the
low-temperature transformation phase as a hard phase is formed in a
portion other than the ferrite phase as a main phase or a so-called
second phase to provide a composite phase. At this moment, the
feature that a hard martensite phase among the low-temperature
transformation phases is particularly existent in the texture is
advantageous because the fraction of the second phase for obtaining
the target tensile strength level is made small and the fraction of
ferrite phase is made large, whereby the increase of the Young's
modulus is attained and further the workability can be improved.
For this end, the martensite phase is required to be not less than
1% as an area ratio to the whole of the texture. In order to obtain
the strength of lot less than 700 MPa, the area ratio of the
martensite phase is preferable to be not less than 16%.
[0099] The texture of the steel sheet according to the invention is
preferable to be a texture comprising ferrite phase and martensite
phase, but there is no problem that phases other than the ferrite
phase and martensite phase such as bainite phase, residual
austenite phase, pearlite phase, cementite phase and the like are
existent at the area ratio of not more than 10%, preferably not
more than 5%. That is, the sum of area ratios of ferrite phase and
martensite phase is preferably not less than 90%, more preferably
not less than 95%.
[0100] Next, the reason on the production conditions limited for
obtaining the high-stiffness high-strength thin steel sheet
according to the invention and preferable production conditions
will be explained.
[0101] The composition of the starting material of steel used in
the production method of the invention is the same as the
composition of the aforementioned steel sheet, so that the
description of the reason on the limitation of the starting
material of steel is omitted.
[0102] The thin steel sheet according to the invention can be
produced by successively conducting a hot rolling step of
subjecting the starting material of steel having the same
composition as the composition of the steel sheet to a hot rolling
to obtain a hot rolled sheet, a cold rolling step of subjecting the
hot rolled sheet after pickling to a cold rolling to obtain a cold
rolled sheet, and an annealing step of attaining the
recrystallization and composite texture in the cold rolled
sheet.
[0103] (Hot Rolling Step)
[0104] Finish rolling: total rolling reduction below 950.degree. C.
is not less than 30%, and the rolling is terminated at
Ar.sub.3-900.degree. C.
[0105] In the final rolling at the hot rolling step, the rolling is
conducted just above Ar.sub.3 transformation point to develop a
non-recrystallized austenite texture having a crystal orientation
of {112}<111>, and the {112}<111> non-recrystallized
austenite can be transformed to ferrite in the subsequent cooling
stage to develop ferrite orientation of {113}<110>. This
orientation advantageously acts to the improvement of the Young's
modulus in the formation of the texture at the subsequent cold
rolling and annealing steps. In order to obtain such an action, it
is required that the total rolling reduction below 950.degree. C.
(total rolling reduction) is not less than 30%, more preferably the
total rolling reduction below 900.degree. C. is not less than 30%,
and the finish rolling is terminated at a temperature region of
Ar.sub.3-900.degree. C., preferably Ar.sub.3-850.degree. C.
[0106] Coiling temperature: not higher than 650.degree. C.
[0107] When the coiling temperature after the finish rolling
exceeds 650.degree. C., the carbonitride of Nb is coarsened and the
effect of suppressing the recrystallization of ferrite becomes
small in the temperature rising stage at the annealing step after
the cold rolling and it is difficult to transform the
non-recrystallized ferrite into austenite. As a result, the
orientation of the low-temperature transformation phase transformed
in the cooling stage after the soaking can not be controlled, and
the Young's modulus is largely lowered by the low-temperature
transformation phase having such a strain. Therefore, the coiling
temperature after the finish rolling is required to be not higher
than 650.degree. C.
[0108] Moreover, when the coiling temperature is too low, a great
amount of the hard low-temperature transformation phase is produced
and the subsequent cold rolling becomes difficult, so that it is
preferable to be not lower than 400.degree. C.
[0109] (Cold Rolling Step)
[0110] Cold rolling is carried out at a rolling reduction of not
less than 50% after the pickling.
[0111] After the hot rolling step, the pickling is carried out for
removing scale formed on the surface of the steel sheet. The
pickling may be conducted according to the usual manner.
Thereafter, the cold rolling is conducted. By the cold rolling at a
rolling reduction of not less than 50% can be turned the
orientation of {113}<110> developed on the hot rolled steel
sheet to an orientation of {112}<110> effective for the
increase of the Young's modulus. Thus, as the orientation of
{112}<110> is developed by the cold rolling, the orientation
of {112}<110> in ferrite is enhanced in the texture after the
subsequent annealing step and further the orientation of
{112}<110> is developed in the low-temperature transformation
phase, whereby the Young's modulus can be increased. In order to
obtain such an effect, the rolling reduction in the cold rolling is
required to be not less than 50%.
[0112] (Annealing Step)
[0113] Temperature rising rate from 500.degree. C. to soaking
temperature: 1-40.degree. C./s, Soaking temperature:
780-900.degree. C.
[0114] The temperature rising rate at the annealing step is an
important process condition in the invention. In the course of
raising the temperature to a soaking temperature of two-phase
region or a soaking temperature of 780-900.degree. C. at the
annealing step, the recrystallization of ferrite having an
orientation of {112}<110> is promoted, while a part of
ferrite grains having an orientation of {112}<110> is arrived
to a two-phase region at a non-recrystallized state, whereby the
transformation from the non-recrystallized ferrite having an
orientation of {112}<110> can be promoted. Therefore, the
Young's modulus can be increased by promoting the growth of ferrite
grains having an orientation of {112}<110> when austenite is
transformed into ferrite in the cooling after the soaking. Further,
when the strength is increased by producing the low-temperature
transformation phase, austenite phase transformed from ferrite
having an orientation of {112}<110> is re-transformed in the
cooling, so that {112}<110> can be also developed with
respect to the crystal orientation of the low-temperature
transformation phase. By developing {112}<110> of ferrite
phase is increased the Young's modulus, while {112}<110> is
particularly developed in the orientation of the low-temperature
transformation phase largely influencing the lowering of the
Young's modulus, whereby the lowering of the Young's modulus
accompanied with the formation of the low-temperature
transformation phase can be suppressed while forming the
low-temperature transformation phase. When austenite is transformed
from the non-recrystallized ferrite while promoting the
recrystallization of ferrite in the temperature rising stage, an
average temperature rising rate largely exerting on the
recrystallization behavior from 500.degree. C. to 780-900.degree.
C. as a soaking temperature is required to be 1-40.degree. C./s,
preferably 1-30.degree. C./s.
[0115] In this case, the reason why the soaking temperature is
780-900.degree. C. is due to the fact that when it is lower than
780.degree. C., the recrystallization is not completed, while when
it exceeds 900.degree. C., the fraction of austenite becomes large
and ferrite having an orientation of {112}<110> reduces or
disappears. Moreover, the soaking time is not particularly limited,
but it is preferable to be not less than 30 seconds for forming
austenite, while it is preferable to be not more than about 300
seconds because the production efficiency is deteriorated as the
time is too long.
[0116] Cooling rate to 500.degree. C. after soaking: not less than
5.degree. C./s
[0117] In the cooling stage after the soaking, it is required to
form the low-temperature transformation phase containing martensite
for increasing the strength. Therefore, an average cooling rate to
500.degree. C. after the soaking is required to be not less than
5.degree. C./s.
[0118] In the invention, steel having a chemical composition in
accordance with the target strength level is first melted. As the
melting method can be properly applied a usual converter process,
an electric furnace process and the like. The molten steel is cast
into a slab, which is subjected to a hot rolling as it is or after
the cooling and heating. After the finish rolling under the
aforementioned finish conditions in the hot rolling, the steel
sheet is coiled at the afore-mentioned coiling temperature and then
subjected to usual pickling and cold rolling. As to the annealing,
the temperature is raised under the aforementioned condition, and
in the cooling after the soaking, the cooling rate can be increased
within a range of obtaining a target low-temperature transformation
phase. Thereafter, the cold rolled steel sheet may be subjected to
an overaging treatment, or may be passed through a hot dip zinc in
case of producing as a galvanized steel sheet, or further in case
of producing as an alloyed galvanized steel sheet, a re-heating may
be conducted up to a temperature above 500.degree. C. for the
alloying treatment.
EXAMPLES
[0119] The following examples are given in illustration of the
invention and are not intended as limitations thereof.
[0120] At first, a steel A having a chemical composition shown in
Table 1 is melted in a vacuum melting furnace of a laboratory and
cooled to room temperature to prepare a steel ingot (steel raw
material).
TABLE-US-00001 TABLE 1 Kind Chemical composition of X Y steel C Si
Mn P S Al N Nb value value Remarks A 0.04 0.2 2.5 0.02 0.001 0.03
0.002 0.08 0.03 0.011 Acceptable example
[0121] Thereafter, the hot rolling, pickling, cold rolling and
annealing are successively conducted in the laboratory. The basic
production conditions are as follows. After the steel ingot is
heated at 1250.degree. C. for 1 hour, the hot rolling is conducted
under conditions that the total rolling reduction below 900.degree.
C., i.e. total rolling reduction ratio below 900.degree. C. is 40%
and the final rolling temperature (corresponding to a final
temperature of finish rolling) is 830.degree. C. to obtain a hot
rolled sheet having a thickness of 4.0 mm. Thereafter, the coiling
condition (corresponding to a coiling temperature of 600.degree.
C.) is simulated by leaving the hot rolled sheet up to 600.degree.
C. and keeping in a furnace of 600.degree. C. for 1 hour and then
cooling in the furnace. The thus obtained hot rolled sheet is
pickled and cold-rolled at a rolling reduction of 60% to a
thickness of 1.6 mm. Then, the temperature of the cold rolled sheet
is raised at 10.degree. C./s on average up to 500.degree. C. and
further from 500.degree. C. to a soaking temperature of 820.degree.
C. at 5.degree. C./s on average. Next, the soaking is carried out
at 820.degree. C. for 180 seconds, and thereafter the cooling is
carried out at an average cooling rate of 10.degree. C./s up to
500.degree. C., and further the temperature of 500.degree. C. is
kept for 80 seconds, and then the sheet is cooled in air. Moreover,
Ar.sub.3 transformation point of this steel under the above
production conditions is 730.degree. C.
[0122] In this experiment, the following conditions are further
individually changed under the above production conditions as a
basic condition. That is, the experiment is carried out under the
basic condition except for the individual changed conditions that
the total rolling reduction below 950.degree. C. or total rolling
reduction below 900.degree. C. is 20-65% and the final temperature
of the hot finish rolling is 710-920.degree. C. and the coiling
temperature is 500-670.degree. C. and the rolling reduction of the
cold rolling is 40-75% (thickness: 2.4-1.0 mm) and the average
temperature rising rate from 500.degree. C. to the soaking
temperature (820.degree. C.) in the annealing is 0.5-45.degree.
C./s.
[0123] From the sample after the annealing is cut out a test
specimen of 10 mm.times.120 mm in a direction perpendicular to the
rolling direction as a longitudinal direction, which is finished to
a thickness of 0.8 mm by a mechanical polishing and a chemical
polishing for removing strain, and thereafter a resonance frequency
of the sample is measured by using a lateral vibration type
internal friction measuring device to calculate a Young's modulus
therefrom. With respect to the sheet subjected to a temper rolling
of 0.5%, a tensile test specimen of JIS No. 5 is cut out in the
direction perpendicular to the rolling direction and subjected to a
tensile test. Further, the sectional texture is observed by a
scanning type electron microscope (SEM) after the corrosion with
Nital to judge the kind of the texture, while three photographs are
shot at a visual region of 30 .mu.m.times.30 .mu.m and then area
ratios of ferrite phase and martensite phase are measured by an
image processing to determine an average value of each phase as an
area ratio (fraction) of each phase.
[0124] As a result, the values of the mechanical characteristics
under the basic condition in the experiment according to the
production method of the invention are Young's modulus E: 245 GPa,
TS: 800 MPa, E1: 20%, fraction of ferrite phase: 70% and fraction
of martensite phase: 25%, from which it is clear that the thin
steel sheet has an excellent balance of strength-ductility and a
high Young's modulus. Moreover, the remainder of the texture other
than ferrite phase and martensite phase is either of bainite phase,
residual austenite phase, pearlite phase and cementite phase.
[0125] Then, the relationship between the production conditions and
Young's modulus is explained based on the above test results with
reference to the drawings. Even in any experimental conditions, the
tensile strength is 750-850 MPa, and the fraction of ferrite phase
is 80-60%, the fraction of martensite phase is 17-40%, and the
remainder of the texture of the second phase other than martensite
phase is either of bainite phase, residual austenite phase,
pearlite phase and cementite phase.
[0126] In FIG. 1 is shown influences of the total rolling reduction
below 950.degree. C. and the total rolling reduction below
900.degree. C. upon Young's modulus, respectively. When the total
rolling reduction below 950.degree. C. is not less than 30% being
the acceptable range of the invention, the Young's modulus
indicates an excellent value of not less than 225 GPa, and further
when the total rolling reduction below 900.degree. C. is not less
than 30%, the Young's modulus indicates a more excellent value of
not less than 240 GPa.
[0127] In FIG. 2 is shown an influence of the final temperature of
the hot finish rolling upon the Young's modulus. When the final
temperature is Ar.sub.3-900.degree. C. being the acceptable range
of the invention, the Young's modulus indicates an excellent value
of not less than 225 GPa, and further when the final temperature is
Ar.sub.3-850.degree. C., the Young's modulus indicates a more
excellent value of not less than 240 GPa.
[0128] In FIG. 3 is shown an influence of the coiling temperature
upon the Young's modulus. When the coiling temperature is not
higher than 650.degree. C. being the acceptable range of the
invention, the Young's modulus indicates an excellent value of not
less than 225 GPa.
[0129] In FIG. 4 is shown an influence of the rolling reduction of
the cold rolling upon the Young's modulus. When the rolling
reduction is not less than 50% being the acceptable range of the
invention, the Young's modulus indicates an excellent value of not
less than 225 GPa.
[0130] In FIG. 5 is shown an influence of the average temperature
rising rate from 500.degree. C. to the soaking temperature of
820.degree. C. in the annealing upon the Young's modulus. When the
temperature rising rate is 1-40.degree. C./s being the acceptable
range of the invention, the Young's modulus indicates an excellent
value of not less than 225 GPa, and further when the temperature
rising rate is 1-30.degree. C./s, the Young's modulus indicates a
more excellent value of not less than 240 GPa.
[0131] Furthermore, steels B-Z and AA-BF having a chemical
composition as shown in Tables 2 and 3 are melted in a vacuum
melting furnace of a laboratory and then successively subjected to
the hot rolling, pickling, cold rolling and annealing under the
above basic condition, respectively. In Tables 4 and 5 are shown
characteristics obtained by the aforementioned tests. Moreover, the
Ar.sub.3 transformation point in the steels B-Z and AA-BF under the
above production conditions is 650-760.degree. C. Also, the
residual texture other than ferrite phase and martensite phase in
the tables is either of bainite phase, residual austenite phase,
pearlite phase and cementite phase.
TABLE-US-00002 TABLE 2 Kind Chemical composition (mass %) of other
X Y steel C Si Mn P S Al N Nb components value N* value Remarks B
0.02 0.2 2.5 0.02 0.001 0.03 0.002 0.07 -- 0.01 -- 0.009 Acceptable
Steel C 0.02 0.2 2.5 0.02 0.001 0.03 0.002 0.14 -- 0.00 -- 0.020
Comparative Steel D 0.06 0.2 2.5 0.02 0.001 0.03 0.002 0.12 -- 0.05
-- 0.017 Acceptable Steel E 0.07 0.2 2.5 0.02 0.001 0.03 0.002 0.08
-- 0.06 -- 0.011 Acceptable Steel F 0.04 0.2 2.5 0.02 0.001 0.03
0.002 0.25 -- 0.01 -- 0.036 Acceptable Steel G 0.06 0.2 2.5 0.02
0.001 0.03 0.002 0.35 -- 0.02 -- 0.051 Acceptable Steel H 0.05 0.2
2.5 0.02 0.001 0.03 0.002 0.05 -- 0.05 -- 0.006 Acceptable Steel I
0.05 0.2 2.5 0.02 0.001 0.03 0.002 0.04 -- 0.05 -- 0.005 Acceptable
Steel J 0.11 0.2 2.5 0.02 0.001 0.03 0.002 0.30 -- 0.07 -- 0.044
Comparative Steel K 0.04 0.2 1.4 0.02 0.001 0.03 0.002 0.08 -- 0.03
-- 0.011 Comparative Steel L 0.04 0.2 1.5 0.02 0.001 0.03 0.002
0.08 -- 0.03 -- 0.011 Acceptable Steel M 0.04 0.2 2.0 0.02 0.001
0.03 0.002 0.08 -- 0.03 -- 0.011 Acceptable Steel N 0.04 0.2 3.5
0.02 0.001 0.03 0.002 0.08 -- 0.03 -- 0.011 Acceptable Steel O 0.04
0.2 3.7 0.02 0.001 0.03 0.002 0.08 -- 0.03 -- 0.011 Acceptable
Steel P 0.02 0.01 2.5 0.01 0.001 0.03 0.002 0.07 -- 0.01 -- 0.009
Acceptable Steel Q 0.02 1.5 2.5 0.01 0.001 0.03 0.002 0.07 -- 0.01
-- 0.009 Acceptable Steel R 0.02 0.2 2.5 0.01 0.001 0.5 0.002 0.07
-- 0.01 -- 0.009 Acceptable Steel S 0.02 0.2 2.5 0.01 0.001 1.0
0.002 0.07 -- 0.01 -- 0.009 Acceptable Steel T 0.02 0.2 2.5 0.01
0.001 1.5 0.002 0.07 -- 0.01 -- 0.009 Acceptable Steel U 0.02 1.5
2.5 0.01 0.001 1.0 0.002 0.07 -- 0.01 -- 0.009 Acceptable Steel V
0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Ti: 0.01 0.03 0.000 0.011
Acceptable Steel W 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Ti: 0.05
0.02 0.000 0.011 Acceptable Steel X 0.07 0.2 2.5 0.02 0.001 0.03
0.002 0.08 Ti: 0.18 0.02 0.000 0.011 Acceptable Steel Y 0.04 0.2
2.5 0.02 0.001 0.03 0.002 0.08 V: 0.05 0.02 0.002 0.011 Acceptable
Steel Z 0.08 0.2 2.5 0.02 0.001 0.03 0.002 0.08 V: 0.20 0.02 0.002
0.011 Acceptable Steel Note) In case of adding no Ti or V, X value
= C + (12/14) .times. N - (12/92.9) .times. Nb In case of adding Ti
or V, X value = C + (12/14) .times. N* - (12/92.9) .times. Nb -
(12/47.9) .times. Ti* - (12/50.9) .times. V Y value = (14/92.9)
.times. (Nb - 0.01) provided that N* = N - (14/47.9) .times. Ti at
N - (14/47.9) .times. Ti > 0, N* = 0 at N - (14/47.9) .times. Ti
.ltoreq. 0, Ti* = Ti - (47.9/14) .times. N - (47.9/32.1) .times. S
at Ti - (47.9/14) .times. N - (47.9/32.1) .times. S > 0, Ti* = 0
at Ti - (47.9/14) .times. N - (47.9/32.1) .times. S .ltoreq. 0.
TABLE-US-00003 TABLE 3 Kind Chemical composition (mass %) of other
X Y steel C Si Mn P S Al N Nb components value N* value Remarks AA
0.07 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Ti: 0.10, 0.01 0.000 0.011
Acceptable V: 0.10 Steel AB 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08
Cr: 0.1 0.03 -- 0.011 Acceptable Steel AC 0.04 0.2 2.5 0.02 0.001
0.03 0.002 0.08 Cr: 1.0 0.03 -- 0.011 Acceptable Steel AD 0.04 0.2
2.5 0.02 0.001 0.03 0.002 0.08 Ni: 0.2 0.03 -- 0.011 Acceptable
Steel AE 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Ni: 1.0 0.03 --
0.011 Acceptable Steel AF 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08
Mo: 0.2 0.03 -- 0.011 Acceptable Steel AG 0.04 0.2 2.5 0.02 0.001
0.03 0.002 0.08 Mo: 1.0 0.03 -- 0.011 Acceptable Steel AH 0.04 0.2
2.5 0.02 0.001 0.03 0.002 0.08 Cu: 0.3 0.03 -- 0.011 Acceptable
Steel AI 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Cu: 2.0 0.03 --
0.011 Acceptable Steel AJ 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08
B: 0.0010 0.03 -- 0.011 Acceptable Steel AK 0.04 0.2 2.5 0.02 0.001
0.03 0.002 0.08 B: 0.0030 0.03 -- 0.011 Acceptable Steel AL 0.04
0.2 2.5 0.02 0.001 0.03 0.002 0.08 Cr: 0.1, 0.03 -- 0.011
Acceptable Ni: 0.1 Steel AM 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08
Cr: 0.1, 0.03 -- 0.011 Acceptable Mo: 0.1 Steel AN 0.04 0.2 2.5
0.02 0.001 0.03 0.002 0.08 Cr: 0.1, 0.03 -- 0.011 Acceptable B:
0.0010 Steel AO 0.04 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Cr: 0.1,
0.03 -- 0.011 Acceptable Ni: 0.1, Steel Mo: 0.1, Cu: 0.1, B: 0.0010
AP 0.06 0.2 2.5 0.02 0.001 0.03 0.002 0.08 Ti: 0.1, 0.01 0.000
0.011 Acceptable V: 0.05, Steel Cr: 0.2, B: 0.0010 AQ 0.04 0.2 2.5
0.02 0.001 0.03 0.002 0.08 Ti: 0.05, 0.01 0.000 0.011 Acceptable V:
0.02, Steel Cr: 0.2, Ni: 0.2, Mo: 0.1, Cu: 0.1, B.0005 AR 0.13 0.01
2.0 0.02 0.001 0.03 0.003 0.06 Ti: 0.15, 0.06 0.000 0.008
Acceptable V: 0.10 Steel AS 0.15 0.1 2.1 0.03 0.002 0.02 0.002 0.03
Ti: 0.24, 0.06 0.000 0.003 Acceptable V: 0.10 Steel AT 0.16 0.1 2.2
0.01 0.001 0.03 0.001 0.09 Ti: 0.31 0.07 0.000 0.012 Comparative
Steel AX 0.06 0.2 3.9 0.01 0.001 0.03 0.002 0.02 Ti: 0.02 0.05
0.000 0.002 Acceptable Steel AY 0.07 0.01 4.2 0.02 0.002 0.05 0.001
0.05 Ti: 0.01 0.06 0.000 0.006 Comparative Steel AZ 0.05 0.01 2.3
0.03 0.001 0.04 0.002 -- -- 0.05 -- -0.002 Comparative Steel BA
0.06 0.01 1.8 0.01 0.001 0.03 0.001 0.01 -- 0.06 -- 0.000
Comparative Steel BB 0.04 0.01 1.5 0.01 0.001 0.02 0.003 0.02 Ti:
0.02 0.04 0.000 0.002 Acceptable Steel BC 0.08 0.1 1.9 0.01 0.001
0.01 0.002 0.05 Ti: 0.25 0.01 0.000 0.006 Acceptable Steel BD 0.14
0.1 1.8 0.02 0.002 0.02 0.002 0.05 Ti: 0.45 0.02 0.000 0.006
Acceptable Steel BE 0.10 0.1 1.9 0.01 0.001 0.03 0.002 0.05 V: 0.20
0.05 0.000 0.006 Acceptable Steel BF 0.12 0.01 1.8 0.02 0.002 0.04
0.001 0.05 V: 0.40 0.02 0.000 0.006 Acceptable Steel Note) In case
of adding no Ti or V, X value = C + (12/14) .times. N - (12/92.9)
.times. Nb In case of adding Ti or V, X value = C + (12/14) .times.
N* - (12/92.9) .times. Nb - (12/47.9) .times. Ti* - (12/50.9)
.times. V Y value = (14/92.9) .times. (Nb - 0.01) provided that N*
= N - (14/47.9) .times. Ti at N - (14/47.9) .times. Ti > 0, N* =
0 at N - (14/47.9) .times. Ti .ltoreq. 0, Ti* = Ti - (47.9/14)
.times. N - (47.9/32.1) .times. S at Ti - (47.9/14) .times. N -
(47.9/32.1) .times. S > 0, Ti* = 0 at Ti - (47.9/14) .times. N -
(47.9/32.1) .times. S .ltoreq. 0.
TABLE-US-00004 TABLE 4 Steel texture Fraction Fraction of of
Mechanical Kind ferrite martensite properties of phase phase TS E1
E steel (%) (%) (MPa) (%) (GPa) Remarks B 95 4 600 30 252 Invention
Example C 100 0 530 32 255 Comparative Example D 50 46 990 15 240
Invention Example E 50 55 1060 12 235 Invention Example F 75 20 860
18 243 Invention Example G 80 19 890 16 242 Invention Example H 70
26 750 21 240 Invention Example I 70 25 750 22 235 Invention
Example J 30 68 1180 10 220 Comparative Example K 90 8 570 30 231
Comparative Example L 85 12 590 29 241 Invention Example M 80 17
650 28 242 Invention Example N 60 35 860 17 242 Invention Example O
50 50 890 16 235 Invention Example P 98 1 590 30 253 Invention
Example Q 90 7 630 30 248 Invention Example R 94 3 620 29 242
Invention Example S 94 3 630 29 241 Invention Example T 93 4 640 28
240 Invention Example U 92 4 650 27 240 Invention Example V 70 25
810 20 246 Invention Example w 75 23 780 21 247 Invention Example X
73 24 810 19 245 Invention Example Y 72 22 800 20 246 Invention
Example Z 68 28 890 15 243 Invention Example
TABLE-US-00005 TABLE 5 Steel texture Fraction Fraction of of
Mechanical Kind ferrite martensite properties of phase phase TS E1
E steel (%) (%) (MPa) (%) (GPa) Remarks AA 85 13 780 20 248
Invention Example AB 65 30 810 19 245 Invention Example AC 60 36
850 17 242 Invention Example AD 64 30 810 19 245 Invention Example
AE 58 37 860 17 241 Invention Example AF 65 31 820 18 243 Invention
Example AG 59 37 870 17 241 Invention Example AH 67 29 810 20 243
Invention Example AI 60 33 840 17 242 Invention Example AJ 60 34
850 17 243 Invention Example AK 50 43 900 15 241 Invention Example
AL 63 34 820 18 242 Invention Example AM 61 34 820 18 241 Invention
Example AN 59 37 860 17 242 Invention Example AO 57 38 870 16 240
Invention Example AP 82 15 760 22 248 Invention Example AQ 84 13
800 20 247 Invention Example AR 70 25 900 15 235 Invention Example
AS 68 30 950 13 226 Invention Example AT 45 55 920 14 213
Comparative Example AX 60 35 850 16 225 Invention Example AY 40 60
1000 13 208 Comparative Example AZ 75 20 760 21 210 Comparative
Example BA 73 21 770 20 215 Comparative Example BB 80 18 700 24 226
Invention Example BC 70 28 820 17 245 Invention Example BD 73 25
920 14 240 Invention Example BE 80 20 800 20 241 Invention Example
BF 70 28 890 17 243 Invention Example
[0132] In the steel C, the C content (X-value) not fixed as a
carbonitride is as small as 0.00%, and the ferrite phase is 100%,
and the fraction of the second phase is 0%, and TS is smaller than
the acceptable range of the invention. In the steel J, the X-value
is as high as 0.07%, and the Young's modulus is smaller than the
acceptable range of the invention. In the steel K, the Mn content
is as low as 1.4%, and TS is smaller than the acceptable range of
the invention. In the steel AT, the C content is as high as 0.16%,
and the X-value is as high as 0.07, and the Young's modulus is
smaller than the acceptable range of the invention. In the steel
AZ, the Mn content is as large as 4.2%, and the Young's modulus is
smaller than the acceptable range of the invention. In the steel
AZ, Nb is not contained, while in the steel BA, the Mb content is
as small as 0.01%, so that the Young's modulus is smaller than the
acceptable range of the invention.
[0133] With respect to the other steels, all items are within the
acceptable range of the invention, and TS and Young's modulus
satisfy the acceptable range of the invention.
INDUSTRIAL APPLICABILITY
[0134] According to the invention, it is possible to provide
high-stiffness high-strength thin steel sheets having a tensile
strength of not less than 590 MPa and a Young's modulus of not less
than 225 GPa.
* * * * *