U.S. patent application number 12/736903 was filed with the patent office on 2011-04-07 for high strength hot rolled steel sheet for line pipe use excellent in low temperature toughness and ductile fracture arrest performance and method of production of same.
Invention is credited to Hiroshi Abe, Tatsuo Yokoi, Osamu Yoshida.
Application Number | 20110079328 12/736903 |
Document ID | / |
Family ID | 41377195 |
Filed Date | 2011-04-07 |
United States Patent
Application |
20110079328 |
Kind Code |
A1 |
Yokoi; Tatsuo ; et
al. |
April 7, 2011 |
HIGH STRENGTH HOT ROLLED STEEL SHEET FOR LINE PIPE USE EXCELLENT IN
LOW TEMPERATURE TOUGHNESS AND DUCTILE FRACTURE ARREST PERFORMANCE
AND METHOD OF PRODUCTION OF SAME
Abstract
The present invention has as its object the provision of hot
rolled steel sheet (hot coil) for line pipe use in which API5L-X80
standard or better high strength and low temperature toughness and
ductile fracture arrest performance are achieved and a method of
production of the same. For this purpose, the hot rolled steel
sheet of the present invention comprises C, Si, Mn, Al, N, Nb, Ti,
Ca, V, Mo, Cr, Cu, and Ni in predetermined ranges and a balance of
Fe and unavoidable impurities, in which the microstructure is a
continuously cooled transformed structure, in which continuously
cooled transformed structure, precipitates containing Nb have an
average size of 1 to 3 nm and are included dispersed at an average
density of 3 to 30.times.10.sup.22/m.sup.3, granular bainitic
ferrite and/or quasi-polygonal ferrite are included in 50% or more
in terms of fraction, furthermore, precipitates containing Ti
nitrides are included, and they have an average circle equivalent
diameter of 0.1 to 3 .mu.m and include complex oxides including Ca,
Ti, and Al in 50% or more in terms of number.
Inventors: |
Yokoi; Tatsuo; (Tokyo,
JP) ; Abe; Hiroshi; (Tokyo, JP) ; Yoshida;
Osamu; (Tokyo, JP) |
Family ID: |
41377195 |
Appl. No.: |
12/736903 |
Filed: |
May 25, 2009 |
PCT Filed: |
May 25, 2009 |
PCT NO: |
PCT/JP2009/059922 |
371 Date: |
November 18, 2010 |
Current U.S.
Class: |
148/504 ;
148/330; 148/331; 148/332 |
Current CPC
Class: |
C21D 1/18 20130101; C21C
7/0006 20130101; C21D 2211/005 20130101; C22C 38/42 20130101; C22C
38/46 20130101; C22C 38/06 20130101; C21D 9/085 20130101; C21D
8/0226 20130101; C21D 9/08 20130101; C22C 38/002 20130101; C22C
38/58 20130101; C21D 8/0263 20130101; C21D 7/06 20130101; C21D 9/46
20130101; C22C 38/44 20130101; C22C 38/02 20130101; C22C 38/001
20130101; C21D 1/19 20130101; C21C 7/06 20130101; C22C 38/48
20130101; C22C 38/50 20130101; C21D 2211/002 20130101; C21D 8/021
20130101; C21D 2211/004 20130101; C21D 7/04 20130101 |
Class at
Publication: |
148/504 ;
148/332; 148/330; 148/331 |
International
Class: |
C21D 11/00 20060101
C21D011/00; C22C 38/16 20060101 C22C038/16; C22C 38/00 20060101
C22C038/00; C21D 8/02 20060101 C21D008/02 |
Foreign Application Data
Date |
Code |
Application Number |
May 26, 2008 |
JP |
2008-137195 |
Mar 26, 2009 |
JP |
2009-077146 |
Claims
1. High strength hot rolled steel sheet for line pipe use excellent
in low temperature toughness and ductile fracture arrest
performance containing, by mass %, C=0.02 to 0.06%, Si=0.05 to
0.5%, Mn=1 to 2%, P.ltoreq.0.03%, S.ltoreq.0.005%, O=0.0005 to
0.003%, Al=0.005 to 0.03%, N=0.0015 to 0.006%, Nb=0.05 to 0.12%,
Ti=0.005 to 0.02%, Ca=0.0005 to 0.003% and
N-14/48.times.Ti.gtoreq.0% and
Nb-93/14.times.(N-14/48.times.Ti)>0.05%, further containing
V.ltoreq.0.3% (not including 0%), Mo.ltoreq.0.3% (not including
0%), and Cr.ltoreq.0.3% (not including 0%), where
0.2%.ltoreq.V+Mo+Cr.ltoreq.0.65%, containing Cu.ltoreq.0.3% (not
including 0%) and Ni.ltoreq.0.3% (not including 0%), where
0.1%.ltoreq.Cu+Ni.ltoreq.0.5%, and having a balance of Fe and
unavoidable impurities, in which said steel sheet, the
microstructure is a continuously cooled transformed structure, in
which continuously cooled transformed structure, precipitates
containing Nb have an average size of 1 to 3 nm and are included
dispersed at an average density of 3 to 30.times.10.sup.22/m.sup.3,
granular bainitic ferrite .alpha..sub.B and/or quasi-polygonal
ferrite .alpha..sub.q are included in 50% or more in terms of
fraction, furthermore, precipitates containing Ti nitrides are
included, the precipitates containing Ti nitrides have an average
circle equivalent diameter of 0.1 to 3 .mu.m and include complex
oxides including Ca, Ti, and Al in 50% or more in terms of
number.
2. High strength hot rolled steel sheet for line pipe use excellent
in low temperature toughness and ductile fracture arrest
performance as set forth in claim 1, further containing, by mass %,
B=0.0002 to 0.003%.
3. High strength hot rolled steel sheet for line pipe use excellent
in low temperature toughness and ductile fracture arrest
performance as set forth in claim 1, further containing, by mass %,
REM=0.0005 to 0.02%.
4. A method of production of high strength hot rolled steel sheet
for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance comprising preparing molten
steel for obtaining hot rolled steel sheet having the ingredients
as set forth any one of claims 1 to 3 at which time preparing the
molten steel to give a concentration of Siof 0.05 to 0.2% and a
concentration of dissolved oxygen of 0.002 to 0.008%, adding to the
molten steel Ti in a range giving a final content of 0.005 to 0.3%
for deoxidation, then adding Al within 5 minutes to give a final
content of 0.005 to 0.02%, furthermore adding Ca to give a final
content of 0.0005 to 0.003%, then adding the required amounts of
alloy ingredient elements to cause solidification, cooling a
resultant cast slab, heating said cast slab to a temperature range
of an SRT (.degree. C.) calculated by formula (1) to 1260.degree.
C., further holding the slab at said temperature range for 20
minutes or more, then hot rolling by a total reduction rate of a
non-recrystallization temperature range of 65% to 85%, ending the
rolling in a temperature range of 830.degree. C. to 870.degree. C.,
then cooling in a temperature range down to 650.degree. C. by a
cooling rate of 2.degree. C./sec to 50.degree. C./sec and coiling
at 500.degree. C. to 650.degree. C.: SRT(.degree.
C.)=6670/(2.26-log([% Nb].times.[% C]))-273 (1) where [% Nb]and [%
C]show the contents (mass %) of Nb and C in the steel material.
5. A method of production of high strength hot rolled steel sheet
for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance as set forth in claim 4
characterized by cooling before rolling in said
non-recrystallization temperature range.
6. A method of production of high strength hot rolled steel sheet
for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance as set forth in claim 4
characterized by continuously casting said cast slab at which time
lightly rolling it while controlling the amount of reduction so as
to match solidification shrinkage at a final solidification
position of the cast slab.
Description
TECHNICAL FIELD
[0001] The present invention relates to high strength hot rolled
steel sheet for line pipe use excellent in low temperature
toughness and ductile fracture arrest performance and a method of
production of the same.
BACKGROUND ART
[0002] In recent years, the areas being developed for crude oil,
natural gas, and other energy resources have spread to the North
Sea, Siberia, North America, Sakhalin, and other artic regions and,
further, the North Sea, the Gulf of Mexico, the Black Sea, the
Mediterranean, the Indian Ocean, and other deep seas, that is,
areas of harsh natural environments. Further, from the viewpoint of
the emphasis on the global environment, natural gas development has
been increasing. At the same time, from the viewpoint of the
economy of pipeline systems, a reduction in the weight of the steel
materials or higher operating pressures have been sought. To meet
with these changes in the environmental conditions, the
characteristics demanded from line pipe have become both higher and
more diverse. Broadly breaking them down, there are demands for (a)
greater thickness/higher strength, (b) higher toughness, (c)
improved field weldability and accompanying lower carbon
equivalents (Ceq), (d) tougher corrosion resistance, and (e) higher
deformation performance in frozen areas and earthquake and fault
zones. Further, these characteristics are usually demanded in
combination in accordance with the usage environment.
[0003] Furthermore, due to the recent increase in crude oil and
natural gas demand, far off areas for which development had been
abandoned up to now due to lack of profitability and areas of harsh
natural environments have begun to be developed in earnest. The
line pipe used for pipelines for long distance transport of crude
oil and natural gas is being required to be made thicker and higher
in strength to improve the transport efficiency and also is being
strongly required to be made higher in toughness so as to be able
to withstand use in artic areas. Achievement of both these
characteristics is an important technical goal.
[0004] In line pipe in artic zones, fractures are of a concern. The
fractures due to the internal pressure of line pipe may be roughly
divided into brittle fracture and ductile fracture. The arrest of
propagation of the former brittle fracture can be evaluated by a
DWTT (drop weight tear test) (which evaluates the toughness of
steel in low temperature ranges by the ductile fracture rate and
impart absorbed energy at the time of fracture of a test piece by
an impact test machine), while the arrest of propagation of the
latter ductile fracture can be evaluated by the impact absorbed
energy of a Charpy impact test. In particular, in steel pipe for
natural gas pipeline use, the internal pressure is high and the
crack propagation rate is faster than the speed of the pressure
wave after fracture, so there has been an increase in projects
seeking not only low temperature toughness (brittle fracture
resistance), but also high impact absorbed energy from the
viewpoint of prevention of ductile fracture. Achievement of arrest
properties of both brittle fracture and ductile fracture is now
being sought.
[0005] On the other hand, steel pipe for line pipe use may be
classified by production process into seamless steel pipe, UOE
steel pipe, electric resistance welded steel pipe, and spiral steel
pipe. These are selected in accordance with the application, size,
etc. With the exception of seamless steel pipe, in each case, flat
steel sheet or steel strip is shaped into a tube, then welded to
obtain a steel pipe product. Furthermore, these welded steel pipe
can be classified by the type of steel sheet used as material. Hot
rolled steel sheet (hot coil) of a relatively thin sheet thickness
is used by electric resistance welded steel pipe and spiral steel
pipe, while thick-gauge sheet material (sheet) of a thick sheet
thickness is used by UOE steel pipe. For high strength and large
diameter, thick applications, the latter UOE steel pipe is
generally used. However, from the viewpoint of cost and delivery,
electric resistance welded steel pipe and spiral steel pipe using
the former hot rolled steel sheet as a material are advantageous.
Demand for higher strength, larger diameter, and greater thickness
is increasing.
[0006] In UOE steel pipe, the art of production of high strength
steel pipe corresponding to the X120 standard is disclosed (see
NPLT 1). The above art is predicated on use of heavy sheet as a
material. To obtain both high strength and greater thickness,
interrupted direct quench (IDQ), a feature of the sheet production
process, is used to achieve a high cooling rate and low cooling
stop temperature. In particular, to ensure strength, quench
hardening (structural strengthening) is utilized.
[0007] However, the art of IDQ cannot be applied to the hot rolled
steel sheet used as a material for electric resistance welded steel
pipe and spiral steel pipe. Hot rolled steel sheet is produced by a
process including a coiling step. Due to the restrictions in
capacity of coilers, it is difficult to coil a thick material at a
low temperature. Therefore, the low temperature cooling stop
required for quench hardening is impossible. Therefore, securing
strength by quench hardening is difficult.
[0008] On the other hand, PLT 1 discloses, as art for hot rolled
steel sheet achieving high strength, greater thickness, and low
temperature toughness, the art of adding Ca and Si at the time of
refining so as to make the inclusions spherical and, furthermore,
adding the strengthening elements of Nb, Ti, Mo, and Ni and V
having a crystal grain refinement effect and combining low
temperature rolling and low temperature coiling. However, this art
involves a final rolling temperature of 790 to 830.degree. C., that
is, a relatively low temperature, so there is a drop in absorbed
energy due to separation and a rise in rolling load due to low
temperature rolling and consequently problems remain in operational
stability.
[0009] PLT 2 discloses, as art for hot rolled steel sheet
considering field weldability and excellent in both strength and
low temperature toughness, the art of limiting the PCM value to
keep down the rise in hardness of the weld zone and making the
microstructure a bainitic ferrite single phase and, furthermore,
limiting the ratio of precipitation of Nb. However, this art also
substantially requires low temperature rolling for obtaining a fine
structure. There is a drop in absorbed energy due to separation and
a rise in rolling load due to low temperature rolling and
consequently problems remain in operational stability.
[0010] PLT 3 discloses the art of obtaining ultra high strength
steel sheet excellent in high speed ductile fracture
characteristics by making the ferrite area ratio of the
microstructure 1 to 5% or over 5% to 60% and making the density of
(100) of the cross-section rotated 45.degree. from the rolling
surface about the axis of the rolling direction not more than 3.
However, this art is predicated on UOE steel pipe using heavy sheet
as a material. It is not art covering hot rolled steel sheet.
Citation List
Patent Literature
[0011] PLT 1: Japanese Patent Publication (A) No. 2005-503483
[0012] PLT 2: Japanese Patent Publication (A) No. 2004-315957
[0013] PLT 3: Japanese Patent Publication (A) No. 2005-146407
Non-Patent Literature
[0014] NPLT 1: Nippon Steel Technical Report, No. 380, 2004, page
70
SUMMARY OF INVENTION
Technical Problem
[0015] The present invention has as its object the provision of hot
rolled steel sheet (hot coil) for line pipe use which can not only
withstand use in regions where tough fracture resistance is
demanded, but also in which API5L-X80 standard or better high
strength and low temperature toughness and ductile fracture arrest
performance can both be achieved even with a relatively thick sheet
thickness of for example over half an inch (12.7 mm) and a method
enabling that steel sheet to be produced inexpensively and
stably.
Solution to Problem
[0016] The present invention was made to solve the above problem
and has as its gist the following:
[0017] (1) High strength hot rolled steel sheet for line pipe use
excellent in low temperature toughness and ductile fracture arrest
performance containing, by mass%, [0018] C=0.02 to 0.06%, [0019]
Si=0.05 to 0.5%, [0020] Mn=1 to 2%, [0021] P.ltoreq.0.03%, [0022]
S.ltoreq.0.005%, [0023] O=0.0005 to 0.003%, [0024] Al=0.005 to
0.03%, [0025] N=0.0015 to 0.006%, [0026] Nb=0.05 to 0.12%, [0027]
Ti=0.005 to 0.02%, [0028] Ca=0.0005 to 0.003% and [0029]
N-14/48.times.Ti.gtoreq.0% and [0030]
Nb-93/14.times.(N-14/48.times.Ti)>0.05%, [0031] further
containing [0032] V.ltoreq.0.3% (not including 0%), [0033]
Mo.ltoreq.0.3% (not including 0%), and [0034] Cr.ltoreq.0.3% (not
including 0%), where [0035] 0.2%--V+Mo+Cr.ltoreq.0.65%, containing
[0036] Cu.ltoreq.0.3% (not including 0%) and [0037] Ni.ltoreq.0.3%
(not including 0%), where [0038] 0.1%.ltoreq.Cu+Ni.ltoreq.0.5%, and
[0039] having a balance of [0040] Fe and unavoidable impurities,
[0041] wherein in said steel sheet, [0042] the microstructure is a
continuously cooled transformed structure, in which continuously
cooled transformed structure, [0043] precipitates containing Nb
have an average size of 1 to 3 nm and are included dispersed at an
average density of 3 to 30.times.10.sup.22/m.sup.3, [0044] granular
bainitic ferrite a.sub.B and/or quasi-polygonal ferrite
.alpha..sub.q are included in 50% or more in terms of fraction,
[0045] furthermore, precipitates containing Ti nitrides are
included, [0046] the precipitates containing Ti nitrides have an
average circle equivalent diameter of 0.1 to 3 .mu.m and include
complex oxides including Ca, Ti, and Al in 50% or more in terms of
number.
[0047] (2) High strength hot rolled steel sheet for line pipe use
excellent in low temperature toughness and ductile fracture arrest
performance as set forth in (1), further containing, by mass %,
[0048] B=0.0002 to 0.003%.
[0049] (3) High strength hot rolled steel sheet for line pipe use
excellent in low temperature toughness and ductile fracture arrest
performance as set forth in (1) or (2), further containing, by
mass%, [0050] REM=0.0005 to 0.02%.
[0051] (4) A method of production of high strength hot rolled steel
sheet for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance comprising preparing molten
steel for obtaining hot rolled steel sheet having the compositions
as set forth in any one of claims 1 to 3 at which time preparing
the molten steel to give a concentration of Si of 0.05 to 0.2% and
a concentration of dissolved oxygen of 0.002 to 0.008%, adding to
the molten steel Ti in a range giving a final content of 0.005 to
0.3% for deoxidation, then adding Al within 5 minutes to give a
final content of 0.005 to 0.02%, furthermore adding Ca to give a
final content of 0.0005 to 0.003%, then adding the required amounts
of alloy ingredient elements to cause solidification, cooling a
resultant cast slab, heating the cast slab to a temperature range
of an SRT (.degree. C.) calculated by formula (1) to 1260.degree.
C., further holding the slab at the temperature range for 20
minutes or more, then hot rolling by a total reduction rate of a
non-recrystallization temperature range of 65% to 85%, ending the
rolling in a temperature range of 830.degree. C. to 870.degree. C.,
then cooling in a temperature range down to 650.degree. C. by a
cooling rate of 2.degree. C./sec to 50.degree. C./sec and coiling
at 500.degree. C. to 650.degree. C.:
SRT(.degree. C.)=6670/(2.26-log([% Nb].times.[% C]))-273 (1)
[0052] where [% Nb]and [% C]show the contents (mass %) of Nb and C
in the steel material.
[0053] (5) A method of production of high strength hot rolled steel
sheet for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance as set forth in (4)
characterized by cooling before rolling in the
non-recrystallization temperature range.
[0054] (6) A method of production of high strength hot rolled steel
sheet for line pipe use excellent in low temperature toughness and
ductile fracture arrest performance as set forth in (4) or (5)
characterized by continuously casting the cast slab at which time
lightly rolling it while controlling the amount of reduction so as
to match solidification shrinkage at a final solidification
position of the cast slab.
Advantageous Effects of Invention
[0055] By using the hot rolled steel sheet of the present invention
for hot rolled steel sheet for electric resistance welded steel
pipe and spiral steel pipe use in artic areas where tough fracture
resistance properties are demanded, for example, even with a sheet
thickness of over half an inch (12.7 mm), production of API5L-X80
standard or better high strength line pipe becomes possible. Not
only this, but by using the method of production of the present
invention, hot rolled steel sheet for electric resistance welded
steel pipe and spiral steel pipe use can be inexpensively obtained
in large volumes.
BRIEF DESCRIPTION OF DRAWINGS
[0056] FIG. 1 is a view showing the relationship between the size
of the precipitates containing Ti nitrides and the DWTT brittle
fracture unit.
EMBODIMENTS OF THE INVENTION
[0057] The present inventors etc. first investigated the
relationship between the tensile strength and toughness of hot
rolled steel sheet (hot coil) (in particular, the drop in Charpy
absorbed energy (vE.sub.-20) and the temperature at which the
ductile fracture rate in a DWTT becomes 85% temperature
(FATT.sub.85%)) and the microstructure etc. of steel sheet. They
investigated this assuming the API5L-X80 standard. As a result, the
present inventors etc. discovered that if analyzing the
relationship between the Charpy absorbed energy (vE.sub.-20), which
is an indicator of the ductile fracture arrest performance, and the
amount of addition of C, even with substantially the same strength,
the more the amount of addition of C is increased, the more the
Charpy absorbed energy (vE.sub.-20) tends to fall.
[0058] Therefore, they investigated in detail the relationship of
the vE.sub.-20 and microstructure. As a result, a good correlation
was observed between the vE.sub.-20 and the fraction of the
microstructure containing cementite and other coarse carbides such
as pearlite. That is, it was observed that if such a microstructure
increases, the vE.sub.-20 tends to drop. Further, such a
microstructure tends to increase together with an increase in the
amount of addition of C. Conversely, along with a decrease in the
fraction of a microstructure containing cementite and other coarse
carbides, the fraction of the continuously cooled transformed
structure (Zw) relatively increased.
[0059] A "continuously cooled transformed structure (Zw)", as
described in Iron and Steel Institute of Japan, Basic Research
Group, Bainite Investigation and Research Subgroup ed., Recent
Research on Bainite Structure and Transformation Behavior of Low
Carbon Steel (1994, Iron and Steel Institute of Japan), is a
microstructure defined by a microstructure containing polygonal
ferrite or pearlite formed by a diffusion mechanism and a
transformed structure in the intermediate stage of martensite
formed without diffusion by a shear mechanism.
[0060] That is, a continuously cooled transformed structure
[0061] (Zw), as a structure observed under an optical microscope,
as shown in the above reference literature, pages 125 to 127, is
defined as a microstructure mainly comprised of bainitic ferrite
(.alpha..degree..sub.B), granular bainitic ferrite (.alpha..sub.8),
and quasi-polygonal ferrite (.alpha..sub.q) and furthermore
containing small amounts of residual austenite (.gamma..sub.r) and
martensite-austenite (MA). .alpha..sub.q, like polygonal ferrite
(PF), does not reveal its internal structure by etching, but is
acicular in shape and is clearly differentiated from PF. Here, if
the circumferential length of the crystal grain covered is lq and
its circle equivalent diameter is dq, the grains with a ratio of
the same (lq/dq) satisfying lq/dq.gtoreq.3.5 are .alpha..sub.q.
[0062] The "fraction of a microstructure" is defined as the area
fraction of the above continuously cooled transformed structure in
the microstructure.
[0063] This continuously cooled transformed structure is formed
since the Mn, Nb, V, Mo, Cr, Cu, Ni, and other strengthening
elements added for securing strength when reducing the amount of
addition of C cause an improvement in the quenchability. It is
believed that when the microstructure is a continuously cooled
transformed structure, the microstructure does not contain
cementite and other coarse carbides, so the Charpy absorbed energy
(vE.sub.-20), the indicator of the ductile fracture arrest
performance, is improved.
[0064] On the other hand, no clear correlation could be observed
between the temperature in a DWTT test at which the ductile
fracture rate becomes 85%, an indicator of the low temperature
toughness (below, referred to as the "FATT.sub.85%"), and the
amount of addition of C. Further, even if the microstructure was a
continuously cooled transformed structure, the FATT.sub.85% did not
necessarily improve. Therefore, the inventors etc. examined in
detail the fracture planes after DWTT tests, whereupon they found
the trend that good FATT.sub.85%'s were exhibited when the fracture
unit of the cleavage plane of the brittle fracture is finer. In
particular, the trend was shown that if the fracture unit becomes a
circle equivalent diameter of 30 .mu.m or less, the FATT.sub.85%
becomes good.
[0065] Therefore, the inventors etc. studied in detail the
relationship between microstructures forming continuously cooled
transformed structures and the FATT.sub.85% indicator of low
temperature toughness. They thereby found the trend that if the
fraction of the granular bainitic ferrite (.alpha..sub.B) or
quasi-polygonal ferrite (.alpha..sub.q) forming the continuously
cooled transformed structures increases and the fraction becomes
50% or more, the fracture unit becomes a circle equivalent diameter
of 30 .mu.m or less and the FATT.sub.85% becomes good. Conversely,
they found the trend that if the fraction of the bainitic ferrite
(.alpha..degree..sub.B) increases, the fracture unit conversely
coarsens and the FATT.sub.85% deteriorates.
[0066] In general, the bainitic ferrite (.alpha..degree..sub.B)
forming a continuously cooled transformed structure is separated
into a plurality of regions in the grain boundaries separated by
the prior austenite grain boundaries and, furthermore, with crystal
orientations in the same direction. These are called "packets". The
effective crystal grain size, which is directly related to the
fracture unit, corresponds to this packet size. That is, it is
believed that if the austenite grains before transformation are
coarse, the packet size also becomes coarse, the effective crystal
grain size coarsens, the fracture unit coarsens, and the
FATT.sub.85% deteriorates.
[0067] Granular bainitic ferrite (.alpha..sub.B) is a
microstructure obtained by a more diffusive transformation than
bainitic ferrite (.alpha..degree..sub.B) which occurs in a shearing
manner in relatively large units even among the types of diffusive
transformation. Quasi-polygonal ferrite (.alpha..sub.q) is a
microstructure obtained by even further diffusive transformation.
Originally, this is not comprised of packets of a plurality of
separate regions in the grain boundaries separated by the austenite
grain boundaries and with crystal orientations in the same
direction, but is granular bainitic ferrite (.alpha..sub.B) or
quasi-polygonal ferrite (.alpha..sub.q) with the grains after
transformation themselves in numerous orientations, so the
effective crystal grain size, directly related to the fracture
units, corresponds to the grain size of the same. For this reason,
it is believed that the fracture units become finer and the
FATT.sub.85% is improved.
[0068] The inventors etc. engaged in further studies of the steel
ingredients and production processes giving 50% or more fractions
of granular bainitic ferrite (.alpha..sub.8) or quasi-polygonal
ferrite (.alpha..sub.q) of structures forming a continuously cooled
transformed structure.
[0069] To increase the fraction of granular bainitic ferrite
(.alpha..sub.B) or quasi-polygonal ferrite (.alpha..sub.q), it is
effective to increase the austenite crystal grain boundaries
forming the nuclei of transformation of the microstructure, so the
austenite grains before transformation have to be made finer. In
general, to make austenite grains finer, it is effective to add Nb
or other solute drag or pinning elements enhancing the controlled
rolling (TMCP) effect. However, the fracture units and the change
in FATT.sub.85% due to the same were also observed with the same
type of Nb content. Therefore, with addition of Nb or other solute
drag or pinning elements, the austenite grains before
transformation cannot be made sufficiently finer.
[0070] The inventors etc. investigated the microstructures in more
detail, whereupon they found a good correlation between the
fracture units after a DWTT test and the size of precipitates
containing Ti nitrides. They confirmed the trend that if the
average circle equivalent diameter of the size of precipitates
containing Ti nitrides is 0.1 to 3 .mu.m, the fracture unit after a
DWTT test becomes finer and the FATT.sub.85% is clearly
improved.
[0071] Further, they discovered that the size and dispersion
density of precipitates containing Ti nitrides can be controlled by
deoxidation control in the smelting process. That is, they
discovered that only when optimally adjusting the concentration of
Si and the concentration of dissolved oxygen in the molten steel,
adding Ti for deoxidation, then adding Al and further adding Ca in
that order, the dispersion density of the precipitates containing
Ti nitrides becomes 10.sup.1 to 10.sup.3/mm.sup.2 in range and the
FATT.sub.85% becomes good.
[0072] Furthermore, they learned that when optimally controlled in
this way, the precipitates containing Ti nitrides include, in at
least half by number, complex oxides containing Ca, Ti, and Al.
Further, they newly discovered that by the optimum dispersion of
these oxides, which form the nuclei for precipitation of the
precipitates containing Ti nitrides, the precipitation size and
dispersion density of the precipitates containing Ti nitrides are
optimized and the austenite grain size before transformation kept
fine as it is due to suppression of grain growth due to the pinning
effect and that if the fraction of granular bainitic ferrite
(.alpha..sub.8) or quasi-polygonal ferrite (.alpha..sub.q)
transformed from the fine grain austenite becomes 50% or more, the
FATT.sub.85% indicator of low temperature toughness becomes
good.
[0073] This is because if performing such deoxidation control,
complex oxides containing Ca, Ti, and Al form over half of the
total number of oxides. These fine oxides disperse in a high
concentration. The average circle equivalent diameter of the
precipitates containing Ti nitrides precipitating from these
dispersed fine oxides as nucleation sites becomes 0.1 to 3 .mu.m,
so it is believed that the balance between the dispersion density
and size is optimized, the pinning effect is exhibited to the
maximum extent, and the effect of refining the austenite grain size
before transformation becomes maximized. Note that, the complex
oxides are allowed to contain some Mg, Ce, and Zr.
[0074] Next, the reasons for limitation of the chemical composition
of the present invention will be explained. Here, the % for the
compositions means mass %. C is an element necessary for obtaining
the targeted strength (strength required by API5L-X80 standard) and
microstructure. However, if less than 0.02%, the required strength
cannot be obtained, while if adding over 0.06%, a large number of
carbides, which form starting points of fracture, are formed, the
toughness deteriorates, and also the field weldability
significantly deteriorates. Therefore, the amount of addition of C
is made 0.02% to 0.06%. Further, to obtain a homogeneous strength
without regard to the cooling rate in cooling after rolling, not
more than 0.05% is preferable.
[0075] Si has the effect of suppressing the precipitation of
carbides--which form starting points of fracture. For this reason,
at least 0.05% is added. However, if adding over 0.5%, the field
weldability deteriorates. If considering general use from the
viewpoint of field weldability, not more than 0.3% is preferable.
Furthermore, if over 0.15%, tiger stripe-like scale patterns are
liable to be formed and the beauty of the surface impaired, so
preferably the upper limit should be made 0.15%.
[0076] Mn is a solution strengthening element. Further, it has the
effect of broadening the austenite region temperature to the low
temperature side and facilitating the formation of a continuously
cooled transformed structure, one of the constituent requirements
of the microstructure of the present invention, during the cooling
after the end of rolling. To obtain this effect, at least 1% is
added. However, even if adding over 2% of Mn, the effect becomes
saturated, so the upper limit is made 2%. Further, Mn promotes
center segregation in a continuous casting steel slab and causes
the formation of hard phases forming starting points of fracture,
so the content is preferably made not more than 1.8%.
[0077] P is an impurity and preferably is as low in content as
possible. If over 0.03% is contained, this segregates at the center
part of a continuous casting steel slab and causes grain boundary
fracture and remarkably lowers the low temperature toughness, so
the content is made not more than 0.03%. Furthermore, P has a
detrimental effect on pipemaking and field weldability, so if
considering this, the content is preferably made not more than
0.015%.
[0078] S is an impurity. It not only causes cracks at the time of
hot rolling, but also, if too great in content, causes
deterioration of the low temperature toughness. Therefore, the
content is made not more than 0.005%. Furthermore, S segregates
near the center of a continuous casting steel slab, forms elongated
MnS after rolling, and forms starting points for hydrogen induced
cracking. Not only this, "two sheet cracking" and other
pseudo-separation are liable to occur. Therefore, if considering
the sour resistance, the content is preferably not more than
0.001%.
[0079] O is an element required for causing dispersion of a large
number of fine oxides at the time of deoxidation of molten steel,
so at least 0.0005% is added, but if the content is too great, it
will form coarse oxides forming starting points of fracture in the
steel and cause deterioration of the brittle fracture and hydrogen
induced cracking resistance, so the content is made not more than
0.003%. Furthermore, from the viewpoint of the field weldability, a
content of not more than 0.002% is preferable.
[0080] Al is an element required for causing dispersion of a large
number of fine oxides at the time of deoxidation of molten steel.
To obtain this effect, at least 0.005% is added. On the other hand,
if excessively adding this, the effect is lost, so the upper limit
is made 0.03%.
[0081] Nb is one of the most important elements in the present
invention. Nb suppresses the recovery/recrystallization and grain
growth of austenite during rolling or after rolling by the dragging
effect in the solid solution state and/or the pinning effect as a
carbonitride precipitate, makes the effective crystal grain size
finer, and reduces the fracture unit in crack propagation of
brittle fracture, so has the effect of improving the low
temperature toughness. Furthermore, in the coiling process, a
feature of the hot rolled steel sheet production process, it forms
fine carbides and, by the precipitation strengthening of the same,
contributes to the improvement of the strength. In addition, Nb
delays the .gamma./.alpha. transformation and lowers the
transformation temperature and thereby has the effect of stably
making the microstructure after transformation a continuously
cooled transformed structure even at a relatively slow cooling
rate. However, to obtain these effects, at least 0.05% must be
added. On the other hand, if adding over 0.12%, not only do the
effects become saturated, but also formation of a solid solution in
the heating process before hot rolling becomes difficult, coarse
carbonitrides are formed and form starting points of fracture, and
therefore the low temperature toughness and sour resistance are
liable to be degraded.
[0082] Ti is one of the most important elements in the present
invention. Ti starts to precipitate as a nitride at a high
temperature right after solidification of a cast slab obtained by
continuous casting or ingot casting. These precipitates containing
Ti nitrides are stable at a high temperature and will not dissolve
at all even during subsequent slab reheating, so exhibit a pinning
effect, suppress the coarsening of austenite grains during
reheating, refine the microstructure, and thereby improve the low
temperature toughness. Further, Ti has the effect of suppressing
the formation of nuclei for formation of ferrite in .gamma./.alpha.
transformation and promoting the formation of the continuously
cooled transformed structure of one of the requirements of the
present invention. To obtain such an effect, addition of at least
0.005% of Ti is required. On the other hand, even if adding over
0.02%, the effect is saturated. Furthermore, if the amount of
addition of Ti becomes less than the stoichiometric composition
with N(N-14/48.times.Ti<0%), the residual Ti will bond with C
and the finely precipitated TiC is liable to cause deterioration of
the low temperature toughness. Further, Ti is an element required
for causing dispersion of a large number of fine oxides at the time
of deoxidation of the molten steel. Furthermore, using these fine
oxides as nuclei, precipitates containing Ti nitrides finely
crystallize or precipitate, so this also has the effect of reducing
the average circle equivalent diameter of the precipitates
containing Ti nitrides and cause dense dispersion and thereby the
effect of suppressing recovery/recrystallization of austenite
during rolling or after rolling and also suppressing grain growth
of ferrite after coiling.
[0083] Ca is an element required for causing dispersion of a large
number of fine oxides at the time of deoxidation of molten steel.
To obtain that effect, at least 0.0005% is added. On the other
hand, even if adding more than 0.003%, the effect becomes
saturated, so the upper limit is made 0.003%. Further, Ca, in the
same way as REM, is an element which changes the form of
nonmetallic inclusions, which would otherwise form starting points
for fracture and cause deterioration of the sour resistance, to
render them harmless.
[0084] N, as explained above, forms precipitates containing Ti
nitrides, suppresses coarsening of austenite grains during slab
reheating to make the austenite grain size, which is correlated
with the effective crystal grain size in the later controlled
rolling, finer, and makes the microstructure a continuously cooled
transformed structure to thereby improve the low temperature
toughness. However, if the content is less than 0.0015%, that
effect cannot be obtained. On the other hand, if over 0.006% is
contained, with aging, the ductility falls and the shapeability at
the time of pipemaking falls. As explained before, if the N content
becomes less than the stoichiometric composition with Ti
(N-14/48.times.Ti<0%), the residual Ti will bond with C and the
finely precipitating TiC is liable to cause deterioration of the
low temperature toughness. Furthermore, with a stoichiometric
composition of Nb, Ti, and N of
Nb-93/14.times.(N-14/48.times.Ti).ltoreq.0.05%, the amount of fine
precipitates containing Nb formed in the coiling process decreases
and the strength falls. Therefore, N-14/48.times.Ti.gtoreq.0% and
Nb-93/14.times.(N-14/48.times.Ti)>0.05% are defined.
[0085] Next, the reasons for adding V, Mo, Cr, Ni, and Cu will be
explained. The main objective of further adding these elements to
the basic ingredients is to increase the thickness of the sheet
which can be produced and improve the strength, toughness, and
other properties of the base material without detracting from the
superior features of the steel of the present invention. Therefore,
these elements are ones with self-restricted amounts of addition by
nature.
[0086] V forms fine carbonitrides in the coiling process and
contributes to the improvement of the strength by precipitation
strengthening. However, even if adding more than 0.3%, that effect
becomes saturated, so the content was made not more than 0.3% (not
including 0%). Further, if adding 0.04% or more, there is a concern
over reduction of the field weldability, so less than 0.04% is
preferable.
[0087] Mo has the effect of enhancement of the quenchability and
improvement of the strength. Further, Mo, in the copresence of Nb,
has the effect of strongly suppressing the recrystallization of
austenite during controlled rolling, making the austenite structure
finer, and improving the low temperature toughness. However, even
if adding over 0.3%, the effect becomes saturated, so the content
is made not more than 0.3% (not including 0%). Further, if adding
0.1% or more, there is a concern that the ductility will fall and
the shapeability when forming pipe will fall, so less than 0.1% is
preferable.
[0088] Cr has the effect of raising the strength. However, even if
adding over 0.3%, the effect will become saturated, so the content
is made not more than 0.3% (not including 0%). Further, if adding
0.2% or more, there is a concern over reduction of the field
weldability, so less than 0.2% is preferable. Further, if V+Mo+Cr
is less than 0.2%, the targeted strength is not obtained, while
even if adding more than 0.65%, the effect becomes saturated.
Therefore, 0.2%.ltoreq.V+Mo+Cr.ltoreq.0.65% is prescribed.
[0089] Cu has the effect of improvement of the corrosion resistance
and the hydrogen induced cracking resistance. However, even if
adding more than 0.3%, the effect becomes saturated, so the content
is made not more than 0.3% (not including 0%). Further, if adding
0.2% or more, embrittlement cracking is liable to occur at the time
of hot rolling and to become a cause of surface defects, so less
than 0.2% is preferable.
[0090] Ni, compared with Mn or Cr and Mo, forms fewer hard
structures harmful to the low temperature toughness and sour
resistance in the rolled structure (in particular, the center
segregation zone of the slab) and therefore has the effect of
improving the strength without causing deterioration of the low
temperature toughness and field weldability. However, even if
adding over 0.3%, the effect becomes saturated, so the content is
made not more than 0.3% (not including 0%). Further, there is an
effect of prevention of hot embrittlement of Cu, so at least 1/2 of
the amount of the Cu is added as a general rule.
[0091] Further, if Cu+Ni is less than 0.1%, the effect of
improvement of the strength without causing deterioration of the
corrosion resistance, hydrogen induced cracking resistance, low
temperature toughness, and field weldability is not obtained, while
if over 0.5%, the effect becomes saturated. Therefore,
0.1%.ltoreq.Cu+Ni.ltoreq.0.5% is defined.
[0092] B has the effect of improving the quenchability and
facilitating the formation of a continuously cooled transformed
structure. Furthermore, B has the effect of enhancing the effect of
improvement of the quenchability of Mo and of increasing the
quenchability synergistically with the copresence of Nb. Therefore,
this is added as required. However, if less than 0.0002%, this is
not enough for obtaining those effects, while if adding over
0.003%, slab cracking occurs.
[0093] REMs are elements which change the form of nonmetallic
inclusions, which would otherwise form starting points of fracture
and cause deterioration of the sour resistance, to render them
harmless. However, if adding less than 0.0005%, there is no such
effect, while if adding over 0.02%, large amounts of the oxides are
formed resulting in the formation of clusters and coarse inclusions
which cause deterioration of the low temperature toughness of the
weld seams and have a detrimental effect on the field weldability
as well.
[0094] Next, the microstructure of the steel sheet in the present
invention will be explained in detail. To obtain strength of the
steel sheet, the microstructure must have nanometer size
precipitates containing Nb densely dispersed in it. Further, to
improve the absorbed energy, the indicator of the ductile fracture
arrest performance, a microstructure containing cementite and other
coarse carbides must not be included. Furthermore, to improve the
low temperature toughness, the effective crystal grain size must be
reduced. To observe and measure the nanometer size precipitates
containing Nb effective for precipitation strengthening for
obtaining strength of the steel sheet, thin film observation using
a transmission type electron microscope or measurement by the 3D
atom probe method is effective. Therefore, the inventors etc. used
the 3D atom probe method for measurement.
[0095] As a result, in samples given a strength corresponding to
API5L-X80 by precipitation strengthening, the size of the
precipitates containing Nb extended between 0.5 to 5 nm and the
average size was 1 to 3 nm. The measurement results of the
precipitates containing Nb distributed at a density of 1 to
50.times.10.sup.22/m.sup.3 and having an average density of 3 to
30.times.10.sup.22/m.sup.3 were obtained. The average size of the
precipitates containing Nb, if less than 1 nm, is too small and
therefore the precipitation strengthening ability is not
sufficiently manifested, while if over 3 nm, the precipitates are
transitory, the match with the base phase is lost, and the effect
of precipitation strengthening is reduced. If the average density
of the precipitates containing Nb is less than
3.times.10.sup.22/m.sup.3, the density is not sufficient for
precipitation strengthening, while if over
30.times.10.sup.22/m.sup.3, the low temperature toughness
deteriorates. Here, the "average" is the arithmetic average of the
number. These nanosize precipitates are mainly comprised of Nb, but
are allowed to also include the carbonitride-forming Ti, V, Mo, and
Cr.
[0096] Note that, in the 3D atom probe method, an FIB (focused ion
beam) apparatus/FB2000A made by Hitachi Ltd. was used, and a cut
out sample was electrolytically ground to a needle shape by using a
freely shaped scanning beam to make the grain boundary part a
needle point shape. The sample was given contrast at the crystal
grains differing in orientation by the channeling phenomenon of an
SIM (scan electron microscope) and, while observing this, was cut
at a position including a plurality of grain boundaries by an ion
beam. The apparatus used as the 3D atom probe was an OTAP made by
CAMECA. The measurement conditions were a sample position
temperature of about 70K, a probe total voltage of 10 to 15 kV, and
a pulse ratio of 25%. Each sample was measured three times and the
average value used as the representative value.
[0097] Next, to improve the absorbed energy, the indicator of the
ductile fracture arrest performance, it is necessary that no
microstructure containing cementite or other coarse carbides be
included. That is, the continuously cooled transformed structure in
the present invention is a microstructure containing one or more of
.alpha..degree..sub.B, .alpha..sub.B, .alpha..sub.q, .gamma..sub.r,
and MA, but here, since .alpha..degree..sub.B, .alpha..sub.B, and
.alpha..sub.q do not contain cementite or other coarse carbides, if
their fraction is large, an improvement in the absorbed energy
indicator of ductile fracture arrest performance can be expected.
Furthermore, small amounts of .gamma..sub.r and MA may be included,
but the total amount should be not more than 3%.
[0098] To improve the low temperature toughness, to reduce the
effective crystal grain size, it is not enough just that the
microstructure have a continuously cooled transformed structure. It
is necessary that the .alpha..sub.B and/or .alpha..sub.q structures
forming the continuously cooled transformed structure be 50% or
more in fraction in the continuously cooled transformed structure.
If the fraction of these microstructures is 50% or more, the
effective crystal grain size, which is directly related with the
fracture unit considered the main influential factor in cleavage
fracture propagation in brittle fracture, becomes finer and the low
temperature toughness is improved.
[0099] Further, to obtain the above microstructure, the average
circle equivalent diameter of the precipitates containing Ti
nitrides has to be 0.1 to 3 .mu.m and, furthermore, at least half
of them by number have to contain complex oxides containing Ca, Ti,
and Al. That is, to obtain, as a fraction, 50% or more of the
.alpha..sub.B and/or .alpha..sub.q structures forming the
continuously cooled transformed structure, it is important to make
the austenite grain size before transformation finer. For this
reason, the average circle equivalent diameter of the size of the
precipitates containing Ti nitrides has to be 0.1 to 3 .mu.m
(preferably 2 .mu.m or less) and the density has to be 10.sup.1 to
10.sup.3/mm.sup.2.
[0100] To control the average circle equivalent diameter of size
and the density of the precipitates containing Ti nitrides, it is
sufficient that the oxides of Ca, Ti, and Al forming the
precipitation nuclei of these be optimally dispersed. Due to this,
the precipitation size and dispersion density of the precipitates
containing Ti nitrides are optimized, the austenite grain size
before transformation is kept fine due to suppression of grain
growth by the pinning effect, and therefore the austenite can be
made finer. As a result, it is learned that at least half of the
number of the precipitates containing Ti nitrides should contain
complex oxides containing Ca, Ti, and Al. Note that, the complex
oxides are allowed to contain some Mg, Ce, and Zr. Further, here,
the "average" is the arithmetic average of the number.
[0101] Next, the reasons for limitation of the method of production
of the present invention will be explained in detail.
[0102] In the present invention, the process up to the primary
refining by a converter or electric furnace is not particularly
limited. That is, it is sufficient to tap the pig iron from a blast
furnace, then dephosphorize, desulfurize, and otherwise pretreat
the molten pig iron, then refine it by a converter or to melt scrap
or other cold iron sources by an electric furnace etc.
[0103] The secondary refining process after the primary refining is
one of the most important production processes of the present
invention. That is, to obtain the precipitates containing Ti
nitrides of the targeted composition and size, complex oxides
containing Ca, Ti, and Al must be made to finely disperse in the
steel in the deoxidation process. This can first be realized by
successively adding weak deoxidizing elements to strong deoxidizing
elements in the deoxidation process (successive strength
deoxidation).
[0104] "Successive strength deoxidation" is a deoxidation method
which makes use of the phenomenon that by adding strong deoxidizing
elements to molten steel in which weak deoxidizing element oxides
are present, the weak deoxidizing element oxides are reduced and
oxygen is released in a state of a slow feed rate and small
supersaturation degree, whereupon the oxides formed from the added
strong deoxidizing elements become finer. By adding deoxidizing
elements in stages from the weak deoxidizing element Si
successively to Ti and Al and to the strong deoxidizing element Ca,
these effects can be exhibited to the maximum extent. This will be
explained in sequence below.
[0105] First, the amount of Si, which is a weaker deoxidizing
element than even Ti, is adjusted to make the concentration of
dissolved oxygen in equilibrium with the amount of Si 0.002 to
0.008%. If the concentration of the dissolved oxygen is less than
0.002%, finally a sufficient amount of complex oxides containing
Ca, Ti, and Al for reducing the size of the precipitates containing
Ti nitrides cannot be obtained. On the other hand, if over 0.008%,
the complex oxides formed coarsen and the effect of reducing the
size of the precipitates containing Ti nitrides is lost.
[0106] Further, to stably adjust the concentration of dissolved
oxygen at the preceding stage of deoxidation, addition of Si is
necessary. If the concentration of S is less than 0.05%, the
concentration of dissolved oxygen in equilibrium with Si becomes
over 0.008%, while if over 0.2%, the concentration of dissolved
oxygen in equilibrium with Si becomes less than 0.002%. Therefore,
in the preceding stage of deoxidation, the concentration of S is
made 0.05 to 0.2% and the concentration of dissolved oxygen is made
0.002% to 0.008%.
[0107] Next, in the state of this concentration of dissolved
oxygen, Ti is added in a range giving a final content of 0.005 to
0.3% for deoxidation, then immediately Al is added to give a final
content of 0.005 to 0.02%. At this time, the Ti oxides formed would
grow, agglomerate, coarsen, and rise up together with the elapse of
time after charging the Ti, so the Al is immediately charged.
However, if within 5 minutes, the rise of Ti oxides would not be
that significant, so the Al is preferably charged within 5 minutes
from the charging of the Ti. Further, if the amount of Al charged
is one where the final content becomes less than 0.005%, the Ti
oxides will grow, agglomerate, coarsen, and rise up. On the other
hand, if the amount of Al charged is an amount by which the final
content exceeds 0.02%, the Ti oxides will end up being completely
reduced and finally complex oxides containing Ca, Ti, and Al will
not be sufficiently obtained.
[0108] Next, Ca, which is a stronger deoxidizing element than Ti
and Al, is preferably charged within 5 minutes to give a final
content of 0.0005 to 0.003%. However, after this, in accordance
with need, these elements and other alloy ingredient elements
insufficient in amount may be added. Here, if the amount of Ca
charged is an amount giving a final content of less than 0.0005%,
complex oxides containing Ca, Ti, and Al cannot be sufficiently
obtained. On the other hand, if added to become over 0.003%, the
oxides containing Ti and Al will end up being completely reduced to
Ca and the effects will be lost.
[0109] A slab cast by continuous casting or thin slab casting may
be directly charged as is as a high temperature cast slab to the
hot rolling stand. Further, the slab may be cooled to room
temperature, then reheated at a heating furnace, then hot rolled.
However, when performing hot charge rolling (HCR), due to the
.gamma..fwdarw..alpha..fwdarw..gamma. transformation, the cast
structure is destroyed and the austenite grain size at the time of
slab reheating is reduced, so the steel is preferably cooled to
less than the Ar3 transformation point temperature. Furthermore, it
preferably is cooled to less than the Ar1 transformation point
temperature.
[0110] From the viewpoint of the sour resistance, center
segregation is preferably reduced as much as possible. Therefore,
the slab is cast with light rolling in accordance with the
specifications sought.
[0111] Segregation of Mn etc. raises the quenchability of the
segregated part to cause hardening of the structure and, together
with the presence of inclusions, promotes hydrogen induced
cracking.
[0112] To suppress segregation, light rolling at the time of final
solidification in continuous casting is optimum. The light rolling
at the time of final solidification is performed so as to suppress
movement of concentrated molten steel to the unsolidified part at
the center, caused by the movement of concentrated molten steel due
to solidification shrinkage etc., by compensating for the amount of
solidification shrinkage. Light rolling is performed while
controlling the amount of reduction so as to be commensurate with
the solidification shrinkage at the final solidification position
of the cast slab. Due to this, it is possible to reduce center
segregation.
[0113] The specific conditions of the light rolling are a roll
pitch, in the facility at the position corresponding to the end of
solidification where the center solid phase rate becomes 0.3 to
0.7, of 250 to 360 mm and a reduction rate, expressed by the
product of the casting rate (m/min) and rolling set gradient
(mm/m), of 0.7 to 1.1 mm/min in range.
[0114] At the time of hot rolling, the slab reheating temperature
(SRT) is made a temperature calculated by the following formula
(1)
SRT(.degree. C.)=6670/(2.26-log([% Nb].times.[% C]))-273 (1)
[0115] where, [% Nb]and [% C]show the contents (mass %) of Nb and C
in the steel materials. This formula shows the solubilization
temperature of NbC by the NbC solubility product. If less than this
temperature, the coarse precipitates containing Nb formed at the
time of slab production will not sufficiently melt and the effect
of crystal grain refinement caused by suppression of the
recovery/recrystallization and grain growth of austenite by Nb in
the later rolling process and the delay of .gamma./.alpha.
transformation cannot be obtained. Further, not only this, the
effect of the formation of fine carbides and the improvement of
strength by their precipitation strengthening in the coiling
process, a feature of the hot rolled steel sheet production
process, cannot be obtained. However, if heating at less than
1100.degree. C., the amount of scale-off becomes small and there is
a possibility that inclusions at the slab surface can no longer be
removed together with the scale in the subsequent descaling, so the
slab reheating temperature is preferably 1100.degree. C. or
more.
[0116] On the other hand, if over 1260.degree. C., the grain size
of the austenite becomes coarser, the prior austenite grains in the
subsequent controlled rolling coarsen, a granular microstructure
cannot be obtained after transformation, and the effect of
improvement of the FATT.sub.85% due to the effect of refinement of
the effective crystal grain size cannot be expected. More
preferably, the temperature is 1230.degree. C. or so.
[0117] The slab heating time is made at least 20 minutes from
reaching the above temperature so as to enable sufficient melting
of the precipitates containing Nb. If less than 20 minutes, the
coarse precipitates containing Nb formed at the time of slab
production will not sufficiently melt, and the effect of refinement
of the crystal grains due to suppression of
recovery/recrystallization and grain growth of the austenite during
the hot rolling and the delay of .gamma./.alpha. transformation and
the effect of the formation of fine carbides and the improvement of
strength by their precipitation strengthening in the coiling
process cannot be obtained.
[0118] The following hot rolling process usually is comprised of a
rough rolling process performed by several rolling stands including
a reverse rolling stand and a final rolling process performed by
six to seven rolling stands arranged in tandem. In general, the
rough rolling process has the advantages that the number of passes
and the rolling rates at the individual passes can be freely set,
but the time between passes is long and the structure is liable to
recover/recrystallize between the passes. On the other hand, the
final rolling process employs a tandem setup, so the number of
passes becomes the same as the number of rolling stands, but the
time between passes is short and the effects of controlled rolling
can be easily obtained. Therefore, to realize superior low
temperature toughness, the process has to be designed making full
use of the features of these rolling processes in addition to the
steel ingredients.
[0119] Further, for example, in the case of a product thickness
over 20 mm, if the roll gap in the #1 final rolling stand is 55 mm
or less due to restrictions in the facilities, with the final
rolling process alone, the requirement of the present invention,
that is, the condition of the total reduction rate of the
non-recrystallization temperature range being at least 65%, cannot
be satisfied, so controlling rolling in the non-recrystallization
temperature range may also be performed after the rough rolling
process. In the above case, if necessary, it is possible to wait
until the temperature falls to the non-recrystallization
temperature range or to use a cooling apparatus for cooling. The
latter case enables the waiting time to be shortened, so is more
preferable in terms of productivity.
[0120] Furthermore, a sheet bar may be attached between the rough
rolling and final rolling to enable continuous final rolling. At
that time, the coarse bar is coiled up once, stored in a cover
having a heat retaining function if necessary, and then again
unwound and attached.
[0121] In the rough rolling process, the rolling is mainly
performed in the recrystallization temperature range. The reduction
rates in the individual rolling passes are not limited in the
present invention. However, if the reduction rates at the
individual passes of the rough rolling are 10% or less, sufficient
strain required for recrystallization is not introduced, grain
growth occurs due to only grain boundary movement, the grains
coarsen, and the low temperature toughness is liable to
deteriorate, so it is preferable to perform the rolling by
reduction rates over 10% in the respective rolling passes in the
recrystallization temperature range. Similarly, if the reduction
rates at the rolling passes in the recrystallization temperature
range are 25% or more, particularly in the later low temperature
range, dislocation cell walls will be formed due to the repeated
introduction of dislocations and recovery during the rolling and
dynamic recrystallization involving a change from sub-grain to
large angle grain boundaries will occur. In a structure like a
microstructure mainly comprised of such dynamic recrystallization
grains where high dislocation density grains and other grains are
mixed, grain growth occurs in a short time, so relatively coarse
grains are liable to be grown before the non-recrystallization
region rolling, grains are liable to end up being formed by the
later non-recrystallization region rolling, and therefore the low
temperature toughness is liable to deteriorate. Therefore, the
reduction rates in the rolling passes in the recrystallization
temperature range are preferably made less than 25%.
[0122] In the final rolling process, the rolling is performed in
the non-recrystallization temperature range, but when the
temperature at the end of the rough rolling does not reach the
non-recrystallization temperature range, if necessary it is waited
until the temperature falls to the non-recrystallization
temperature range or, if necessary, cooling is performed by a
cooling apparatus between the rough/final rolling stands. In the
latter case, the waiting time can be shortened, so the productivity
is improved. Not only that, the growth of recrystallization grains
is suppressed and the low temperature toughness can be improved.
This is therefore more preferable.
[0123] If the total reduction rate in the non-recrystallization
temperature range is less than 65%, the controlled rolling becomes
insufficient, prior austenite grains coarsen, a granular
microstructure cannot be obtained after transformation, and the
effect of improvement of the FATT.sub.85% due to the effect of
refinement of the effective crystal grain size cannot be expected,
so the total reduction rate in the non-recrystallization
temperature range is made 65% or more. Furthermore, to obtain a
superior low temperature toughness, 70% or more is preferable. On
the other hand, if over 85%, the excessive rolling causes an
increase in the density of the dislocations forming nuclei for
ferrite transformation and causes polygonal ferrite to be mixed in
the microstructure. Further, due to the high temperature ferrite
transformation, the precipitation strengthening of the Nb becomes
transitory and the strength falls. Further, due to crystal
rotation, the anisotropy of the structure after transformation
becomes remarkable, the plastic anisotropy increases, and a drop in
the absorbed energy due to the occurrence of separation is liable
to be invited. Therefore, the total reduction rate in the
non-recrystallization temperature range is made not more than
85%.
[0124] The final rolling end temperature is 830.degree. C. to
870.degree. C. In particular if less than 830.degree. C. at the
center part of sheet thickness, remarkable separation occurs at the
ductile fracture planes and the absorbed energy remarkably falls,
so the final rolling end temperature at the center part of sheet
thickness is made at least 830.degree. C. Further, the sheet
surface temperature is also preferably made at least 830.degree. C.
On the other hand, if 870.degree. C. or more, even if the
precipitates containing Ti nitrides are optimally present in the
steel, recrystallization is liable to cause the austenite grain
size to coarsen and the low temperature toughness to deteriorate.
Further, if performing the final rolling at the low temperature of
the Ar3 transformation point temperature or less, dual-phase
rolling results, the absorbed energy drops due to the occurrence of
separation, and, in the ferrite phase, due to the reduction, the
dislocation density increases, the precipitation strengthening by
Nb becomes transitory, and the strength falls. Further, the worked
ferrite structure falls in ductility.
[0125] Even without particularly limiting the rolling pass schedule
at the different stands in the final rolling, the effects of the
present invention can be obtained, but from the viewpoint of the
precision of sheet shape, the rolling rate at the final stand is
preferable less than 10%.
[0126] Here, the "Ar.sub.3 transformation point temperature" is for
example simply shown in relation to the steel ingredients by the
following formula. That is, Ar.sub.3=910-310.times.% C+25.times.%
Si-80.times.% Mneq
[0127] where, Mneq=Mn+Cr+Cu+Mo+Ni/2+10(Nb-0.02)
[0128] Alternatively, this is the case of addition of
Mneq=Mn+Cr+Cu+Mo+Ni/2+10(Nb-0.02)+1:B.
[0129] After the end of the final rolling, the cooling is started.
The cooling start temperature is not particularly limited, but if
starting the cooling from less than the Ar.sub.3 transformation
point temperature, the microstructure will contain large amounts of
polygonal ferrite and the strength is liable to drop, so the
cooling start temperature is preferably at least the Ar.sub.3
transformation point temperature.
[0130] The cooling rate in the temperature range from the start of
cooling to 650.degree. C. is made 2.degree. C./sec to 50.degree.
C./sec. If this cooling rate is less than 2.degree. C./sec, the
microstructure will contain large amounts of polygonal ferrite and
the strength is liable to drop. On the other hand, with a cooling
rate of over 50.degree. C./sec, heat strain is liable to cause
warping, so the rate is made not more than 50.degree. C./sec.
[0131] Further, when the occurrence of separation at the fracture
plane results in the predetermined absorbed energy not being
obtained, the cooling rate is made at least 15.degree. C./sec.
Furthermore, if 20.degree. C./sec or more, it is possible to
improve the strength without changing the steel ingredients and
without causing deterioration of the low temperature toughness, so
the cooling rate is preferably made at least 20.degree. C./sec.
[0132] The cooling rate in the temperature range from 650.degree.
C. to coiling may be air cooling or a cooling rate corresponding to
the same. However, to obtain the maximum effect of precipitation
strengthening by Nb etc., to prevent the precipitate from
coarsening and thereby becoming transitory, the average cooling
rate from 650.degree. C. to coiling is preferably at least
5.degree. C./sec.
[0133] After cooling, the coiling process, a feature of the hot
rolled steel sheet production process, is effectively utilized. The
cooling stop temperature and coiling temperature are made
temperature ranges of 500.degree. C. to 650.degree. C. If stopping
the cooling at over 650.degree. C. and then coiling, the
precipitates containing Nb will become transitory and precipitation
strengthening will no longer be sufficiently exhibited. Further,
coarse precipitates containing Nb will form and act as starting
points for fracture and therefore the ductile fracture arresting
ability, low temperature toughness, and sour resistance are liable
to be degraded. On the other hand, if ending the cooling at less
than 500.degree. C. and then coiling, the fine precipitates
containing Nb so effective for obtaining the target strength will
not be obtained and the target strength will no longer be able to
be obtained. Therefore, the temperature range for stopping the
cooling and coiling is made 500.degree. C. to 650.degree. C.
EXAMPLES
[0134] Below, examples will be used to explain the present
invention in more detail. Steels of the chemical ingredients shown
in Table 2 were smelted in a converter and secondarily refined by
CAS or RH. The deoxidation was performed by the secondary refining
process. As shown in Table 1, before charging the Ti, the dissolved
oxygen of the molten steel was adjusted by the concentration of S,
then successive deoxidation was performed by Ti, Al, and Ca. These
steels were continuously cast, then directly charged or reheated
and reduced to a sheet thickness of 20.4 mm by rough rolling and
then final rolling, then were cooled at a runout table, then
coiled. The chemical compositions in the tables are shown in mass
%. Further, the N* in Table 2 means the value of
N-14/48.times.Ti.
TABLE-US-00001 TABLE 1 Production conditions Smelting process
Concen- Equilibrium Time until tration dissolved charging Al of S
before oxygen Order of after Ti charging concentration charging Ti,
deoxidation Steel Ti (%) (%) Al, and Ca (min) Remarks A 0.05 0.0037
Ti.fwdarw.Al.fwdarw.Ca 1.0 Inv. ex. B 0.115 0.0036
Ti.fwdarw.Al.fwdarw.Ca 21.0 Comp. ex. C 0.048 0.0083
Ti.fwdarw.Al.fwdarw.Ca 1.0 Comp. ex. D 0.121 0.0032
Al.fwdarw.Ti.fwdarw.Ca -- Comp. ex. E 0.132 0.0030
Ti.fwdarw.Al.fwdarw.Ca 1.0 Inv. ex. F 0.052 0.0077
Ti.fwdarw.Al.fwdarw.Ca 2.0 Inv. ex. G 0.050 0.0074
Ti.fwdarw.Al.fwdarw.Ca 1.5 Inv. ex. H 0.056 0.0068
Ti.fwdarw.Al.fwdarw.Ca 0.6 Inv. ex. I 0.165 0.0024
Ti.fwdarw.Al.fwdarw.Ca 2.0 Inv. ex. J 0.132 0.0029
Ti.fwdarw.Al.fwdarw.Ca 3.0 Inv. ex. K 0.188 0.0022
Ti.fwdarw.Al.fwdarw.Ca 2.5 Inv. ex. L 0.121 0.0030
Ti.fwdarw.Al.fwdarw.Ca 4.5 Inv. ex. M 0.132 0.0031
Ca.fwdarw.Al.fwdarw.Ti -- Comp. ex. N 0.101 0.0029
Ti.fwdarw.Al.fwdarw.Ca 5.0 Inv. ex. O 0.160 0.0022
Ti.fwdarw.Al.fwdarw.Ca 2.1 Inv. ex. P 0.131 0.0028
Ti.fwdarw.Al.fwdarw.Ca 2.9 Inv. ex. Q 0.184 0.0021
Ti.fwdarw.Al.fwdarw.Ca 2.3 Inv. ex. R 0.120 0.0031
Ti.fwdarw.Al.fwdarw.Ca 4.4 Inv. ex.
TABLE-US-00002 TABLE 2 Chemical composition (unit: mass %) Steel C
Si Mn P S O Al N Nb Ti V A 0.045 0.14 1.76 0.009 0.001 0.0019 0.023
0.0038 0.077 0.012 0.039 B 0.046 0.13 1.73 0.011 0.001 0.0018 0.020
0.0038 0.075 0.012 0.038 C 0.047 0.13 1.75 0.008 0.001 0.0017 0.020
0.0042 0.076 0.013 0.036 D 0.045 0.14 1.75 0.010 0.001 0.0018 0.022
0.0039 0.077 0.013 0.039 E 0.071 0.25 1.87 0.008 0.002 0.0017 0.020
0.0037 0.039 0.012 0.000 F 0.059 0.25 1.74 0.002 0.002 0.0019 0.023
0.0034 0.056 0.011 0.070 G 0.029 0.29 1.65 0.003 0.002 0.0017 0.020
0.0043 0.101 0.014 0.032 H 0.066 0.22 1.54 0.009 0.001 0.0022 0.029
0.0033 0.051 0.021 0.030 I 0.067 0.25 1.60 0.010 0.002 0.0021 0.022
0.0038 0.068 0.003 0.055 J 0.016 0.49 1.79 0.028 0.001 0.0011 0.007
0.0037 0.110 0.012 0.080 K 0.050 0.20 1.85 0.010 0.002 0.0022 0.020
0.0041 0.073 0.013 0.050 L 0.044 0.19 1.78 0.011 0.002 0.0022 0.028
0.0054 0.101 0.018 0.01 M 0.049 0.15 1.75 0.007 0.001 0.0016 0.020
0.0035 0.075 0.011 0.040 N 0.054 0.22 1.80 0.009 0.002 0.0016 0.018
0.0044 0.081 0.014 0.100 O 0.055 0.07 1.79 0.008 0.001 0.0020 0.007
0.0038 0.058 0.012 0.01 P 0.058 0.25 1.79 0.002 0.002 0.0023 0.048
0.0036 0.053 0.012 0.077 Q 0.061 0.24 1.70 0.002 0.002 0.0021 0.020
0.0060 0.056 0.018 0.070 R 0.060 0.35 1.21 0.021 0.002 0.0024 0.023
0.0020 0.081 0.006 0.100 Steel Mo Cr Cu Ni V + Mo + Cr Cu + Ni Ca
N** Nb-93/14 .times. N* Others Remarks A 0.09 0.19 0.19 0.27 0.32
0.46 0.0011 0.0003 0.0750 Inv. ex. B 0.10 0.20 0.20 0.28 0.34 0.48
0.0012 0.0003 0.0730 Comp. ex. C 0.09 0.19 0.20 0.26 0.32 0.46
0.0011 0.0004 0.0733 Comp. ex. D 0.08 0.18 0.18 0.29 0.30 0.47
0.0011 0.0001 0.0763 Comp. ex. E 0.00 0.20 0.16 0.15 0.20 0.31
0.0008 0.0002 0.0377 Comp. ex. F 0.26 0.21 0.25 0.24 0.54 0.49
0.0009 0.0002 0.0547 REM: 0.0020% Inv. ex. G 0.24 0.16 0.23 0.22
0.43 0.45 0.0010 0.0002 0.0996 Inv. ex. H 0.11 0.11 0.11 0.13 0.25
0.24 0.0022 -0.0028 0.0698 Comp. ex. I 0.07 0.11 0.09 0.10 0.24
0.19 0.0010 0.0029 0.0486 Comp. ex. J 0.28 0.10 0.28 0.25 0.46 0.53
0.0010 0.0002 0.1087 Comp. ex. K 0.29 0.01 0.18 0.26 0.34 0.44
0.0021 0.0003 0.0710 Inv. ex. L 0.23 0.22 0.00 0.29 0.45 0.29
0.0026 0.0002 0.1000 B: 0.0008% Inv. ex. M 0.10 0.20 0.20 0.50 0.34
0.70 0.0009 0.0003 0.0730 Comp. ex. N 0.01 0.25 0.25 0.13 0.35 0.38
0.0010 0.0003 0.0789 Inv. ex. O 0.30 0.01 0.25 0.25 0.30 0.50
0.0009 0.0003 0.0560 Inv. ex. P 0.24 0.21 0.25 0.25 0.53 0.50
0.0000 0.0001 0.0523 Comp. ex. Q 0.00 0.00 0.00 0.00 0.07 0.00
0.0011 0.0008 0.0510 Comp. ex. R 0.25 0.25 0.24 0.25 0.60 0.49
0.0009 0.0003 0.0793 Inv. ex.
[0135] Details of the production conditions are shown in Table 3.
Here, "composition" indicates the symbols of the slabs shown in
Table 2, "light rolling" indicates the existence of any light
rolling operation at the time of final solidification in continuous
casting, "heating temperature" indicates the actual slab heating
temperature, "solubilization temperature" indicates the temperature
calculated by
SRT(.degree. C.)=6670/(2.26-log([% Nb].times.[% C]))-273
"holding time" indicates the holding time at the actual slab
heating temperature, "cooling between passes" indicates the
existence of any cooling between rolling stands performed for the
purpose of shortening the temperature waiting time occurring before
non-recrystallization temperature range rolling,
"non-recrystallization region total reduction rate" indicates the
total reduction rate of rolling performed in the recrystallization
temperature range, "FT" indicates the final rolling end
temperature, the "Ar3 transformation point temperature" indicates
the calculated Ar3 transformation point temperature, the "cooling
rate to 650.degree. C." indicates the average cooling rate when
passing through a temperature range of the cooling start
temperature to 650.degree. C., and "CT" indicates the coiling
temperature.
TABLE-US-00003 TABLE 3 Production conditions Reduction rates of
passes in Heating Solubilizing Holding recrystallization region
Steel Composi- Light temperature temperature time (%) No. tion
rolling (.degree. C.) (.degree. C.) (min) 1 2 3 4 5 6 7 8 9 10 11 1
A Yes 1180 1140 30 15 12 13 13 13 14 20 22 -- -- -- 2 A No 1080
1140 30 15 12 13 13 13 14 20 22 -- -- -- 3 A No 1280 1140 30 15 12
13 13 13 14 20 22 -- -- -- 4 A No 1180 1140 5 15 12 13 13 13 14 20
22 -- -- -- 5 A Yes 1180 1140 30 15 12 9 10 10 12 12 12 16 13 -- 6
A No 1180 1140 30 15 10 11 11 10 11 11 13 27 -- -- 7 A No 1180 1140
30 15 12 13 13 13 14 20 22 18 18 -- 8 A No 1180 1140 30 15 12 13 13
13 -- -- -- -- -- -- 9 A No 1180 1140 30 15 12 13 13 13 14 20 22 --
-- -- 10 A Yes 1180 1140 30 15 12 13 13 13 14 20 22 -- -- -- 11 A
No 1180 1140 30 15 12 13 13 13 14 20 22 -- -- -- 12 B No 1170 1139
20 15 12 13 13 13 14 20 22 16 13 -- 13 C No 1170 1144 20 15 12 13
13 13 14 20 22 16 13 -- 14 D No 1170 1140 20 15 12 9 10 10 12 12 12
-- -- -- 15 E No 1170 1111 20 15 12 13 13 13 14 20 22 -- -- -- 16 F
No 1170 1134 20 15 12 13 13 13 14 20 22 -- -- -- 17 G No 1230 1119
20 15 12 13 13 13 14 20 22 -- -- -- 18 H No 1200 1136 30 15 12 13
13 13 14 20 22 16 13 -- 19 I No 1200 1177 30 15 12 13 13 13 14 20
22 -- -- -- 20 J No 1200 1057 30 15 12 13 13 13 14 20 22 -- -- --
21 K Yes 1200 1147 30 15 12 13 13 13 14 20 22 -- -- -- 22 L No 1200
1173 30 23 14 15 16 17 20 19 -- -- -- -- 23 M No 1200 1148 30 15 12
13 13 13 14 20 22 -- -- -- 24 N No 1200 1171 30 15 12 13 13 13 14
20 22 16 13 -- 25 O No 1200 1129 30 15 12 13 13 13 14 20 22 16 13
-- 26 P No 1200 1125 30 15 12 13 13 13 14 20 22 16 13 -- 27 Q No
1200 1138 30 15 12 13 13 13 14 20 22 16 13 -- 28 R No 1200 1185 30
15 12 13 13 13 14 20 22 16 13 -- Non- recrystalli- Ar3 zation
transformation Cooling region total point Cooling Steel between
reduction rate FT temperature rate CT No. passes (%) (.degree. C.)
(.degree. C./sec) (.degree. C./sec) (.degree. C.) Remarks 1 Yes 75
850 665 10 600 Inv. ex. 2 No 75 850 665 10 600 Comp. ex. 3 No 75
850 665 10 600 Comp. ex. 4 Yes 75 850 665 10 600 Comp. ex. 5 No 75
850 665 11 600 Inv. ex. 6 No 75 850 665 15 600 Inv. ex. 7 No 62 850
665 10 600 Comp. ex. 8 No 86 850 665 17 600 Comp. ex. 9 No 75 660
665 10 600 Comp. ex. 10 Yes 75 850 665 1 600 Comp. ex. 11 Yes 75
850 665 10 450 Comp. ex. 12 Yes 75 830 665 15 570 Comp. ex. 13 No
75 830 665 15 570 Comp. ex. 14 No 82 830 667 15 570 Comp. ex. 15 No
75 830 695 15 570 Comp. ex. 16 Yes 75 830 663 25 570 Inv. ex. 17
Yes 75 850 652 12 600 Inv. ex. 18 No 75 850 715 13 600 Comp. ex. 19
No 75 850 703 10 600 Comp. ex. 20 No 0 970 639 10 600 Comp. ex. 21
No 75 850 661 10 600 Inv. ex. 22 No 75 850 646 5 600 Inv. ex. 23 No
80 850 655 5 600 Comp. ex. 24 No 75 830 581 30 600 Inv. ex. 25 No
75 830 587 30 600 Inv. ex. 26 No 75 830 583 30 600 Comp. ex. 27 No
75 830 652 30 600 Comp. ex. 28 No 75 830 603 30 600 Inv. ex.
[0136] The grade of the steel sheet obtained in this way is shown
in Table 4. The methods of examination were as shown below. The
microstructure was examined by cutting out a test piece from a
position of 1/4 W or 3/4 W of the sheet width (W) from an end of
the steel sheet in the width direction, polishing the cross-section
in the rolling direction, using a Nital reagent to etch it, then
obtaining a photo of a field at 1/25 of the sheet thickness
observed using an optical microscope at a power of 200 to
500.times.. Further, the "average circle equivalent diameter of the
precipitates containing Ti nitrides" is defined as that obtained by
observing the same sample as the above at a part at 1/45 of the
sheet thickness (t) from the steel sheet surface using an optical
microscope at a power of 1000.times., obtaining values from
photographs of the microstructure of at least 20 fields by an image
processor etc., and taking the average value of the same.
[0137] Further, the ratio of the complex oxides containing Ca, Ti,
and Al forming the nuclei of the precipitates containing Ti
nitrides is defined as the ratio of the precipitates containing Ti
nitrides observed in the above micrographs which contain such
nuclei-forming complex oxides, that is, (number of precipitates
containing Ti nitrides containing nuclei-forming complex
oxides)/(total number of precipitates containing Ti nitrides
observed). Furthermore, the composition of the nuclei-forming
complex oxides was identified by analysis of at least one oxide in
each field and was confirmed by an energy dispersive X-ray
spectroscope (EDS) or electron energy loss spectroscope (EELS)
attached to a scan type electron microscope.
[0138] The tensile test was conducted by cutting out a No.
[0139] 5 test piece described in JIS Z 2201 from the C direction
and following the method of JIS Z 2241. The Charpy impact test was
conducted by cutting out a test piece described in JIS Z 2202 from
the C direction at the center of sheet thickness and following the
method of JIS Z 2242. The DWTT (drop weight tear test) was
conducted by cutting out a test piece of a strip shape of 300
mmL.times.75 mmW.times.thickness (t) mm in the C direction and
pressing it to give it a 5 mm notch. The HIC test was conducted
based on NACETM0284.
[0140] In Table 4, the "microstructure" is the microstructure of
the part at 1/2 t of the sheet thickness from the surface of the
steel sheet. "Zw" is the continuously cooled transformed structure
and is defined as a microstructure including one or more of
.alpha..degree..sub.B, .alpha..sub.B, .alpha..sub.q, .gamma..sub.r,
and MA. "PF" indicates polygonal ferrite, "worked F" indicates
worked ferrite, "P" indicates pearlite, and the
".alpha..sub.B+.alpha..sub.q fraction" indicates the total area
fraction of granular bainitic ferrite (.alpha..sub.8) and
quasi-polygonal ferrite (.alpha..sub.q).
[0141] The "precipitation strengthening particle size" shows the
size of the precipitates containing Nb effective for precipitation
strengthening as measured by the 3D atom probe method. The
"precipitation strengthening particle density" shows the density of
the precipitates containing Nb effective for precipitation
strengthening as measured by the 3D atom probe method. The "average
circle equivalent diameter" shows the average circle equivalent
diameter of precipitates containing Ti nitrides measured by the
above method. The "content ratio" shows the number ratio of the
above precipitates containing Ti nitrides which include complex
oxides forming nuclei. The "composition of complex oxides" show the
results of analysis by EELS, indicated as "G" (good) when the
elements are detected and as "P" (poor) when not. The results of
the "tensile test" show the results of C-direction JIS No. 5 test
pieces. "FATT.sub.85%" shows the test temperature giving a ductile
fracture rate of 85% in a DWTT test. The "absorbed energy
vE.sub.-20.degree.C." shows the absorbed energy obtained in a
Charpy impact test at -20.degree. C. The "fracture unit" shows the
average value of the fracture units obtained by measurement of
fractures for five or more fields by SEM at a power of about
100.times.. Further, the "strength-vE balance" is expressed as the
product of "TS" and the "absorbed energy vE.sub.-20.degree.C.".
Furthermore, "CAR" shows the area ratio of cracks found by the HIC
test.
TABLE-US-00004 TABLE 4 Precipitates containing Ti nitrides
Microstrueture Average Precipitation Precipitation circle
Composition Mechanical properties .alpha..sub.B + .alpha..sub.q
strengthening strengthening equivalent of complex Tensile test
Steel Micro- fraction particle size particle diameter Content
oxides YP TS EI No. structure (%) (nm) density (/m.sup.3) (.mu.m)
(%) Ca Al Ti (MPa) (MPa) (%) 1 Zw 85 1.5 10 .times. 10.sup.22 2 60
.smallcircle. .smallcircle. .smallcircle. 578 708 32 2 Zw 55 5.0 1
.times. 10.sup.18 2 60 .smallcircle. .smallcircle. .smallcircle.
520 644 36 3 Zw 15 1.8 5 .times. 10.sup.22 2 60 .smallcircle.
.smallcircle. .smallcircle. 590 721 31 4 Zw 50 4.5 1 .times.
10.sup.19 2 60 .smallcircle. .smallcircle. .smallcircle. 550 670 34
5 Zw 90 2.0 4 .times. 10.sup.22 2 60 .smallcircle. .smallcircle.
.smallcircle. 583 711 32 6 Zw 80 2.2 3 .times. 10.sup.22 2 60
.smallcircle. .smallcircle. .smallcircle. 571 699 33 7 Zw 20 1.3 20
.times. 10.sup.22 2 60 .smallcircle. .smallcircle. .smallcircle.
592 722 32 8 PF + Zw -- 7.0 5 .times. 10.sup.17 2 60 .smallcircle.
.smallcircle. .smallcircle. 550 674 33 9 Worked F + P -- 6.0 3
.times. 10.sup.17 2 60 .smallcircle. .smallcircle. .smallcircle.
566 693 24 10 PF + P -- 30.0 4 .times. 10.sup.22 2 60 .smallcircle.
.smallcircle. .smallcircle. 548 671 34 11 Zw 60 0.8 50 .times.
10.sup.20 2 60 .smallcircle. .smallcircle. .smallcircle. 481 636 36
12 Zw 55 1.5 15 .times. 10.sup.22 6 25 .smallcircle. .smallcircle.
.smallcircle. 582 710 32 13 Zw 60 1.3 20 .times. 10.sup.22 6 25
.smallcircle. .smallcircle. .smallcircle. 588 715 32 14 Zw 55 1.3
10 .times. 10.sup.22 6 35 .smallcircle. .smallcircle. x 581 707 33
15 Zw + P -- 3.0 3 .times. 10.sup.22 2.5 60 .smallcircle.
.smallcircle. .smallcircle. 530 644 36 16 Zw 75 1.2 5 .times.
10.sup.22 2 65 .smallcircle. .smallcircle. .smallcircle. 612 745 31
17 Zw 90 1.0 30 .times. 10.sup.22 2.5 50 .smallcircle.
.smallcircle. .smallcircle. 604 736 31 18 Zw + P -- 2.5 5 .times.
10.sup.22 2 50 .smallcircle. .smallcircle. .smallcircle. 574 701 33
19 Zw + P -- 1.5 10 .times. 10.sup.22 2 55 .smallcircle.
.smallcircle. .smallcircle. 581 716 32 20 PF -- -- -- 1 50
.smallcircle. .smallcircle. .smallcircle. 520 641 36 21 Zw 65 1.5 4
.times. 10.sup.22 3 90 .smallcircle. .smallcircle. .smallcircle.
564 710 33 22 Zw 55 2.0 15 .times. 10.sup.22 3 55 .smallcircle.
.smallcircle. .smallcircle. 580 692 33 23 Zw 50 2.5 5 .times.
10.sup.22 5 25 x .smallcircle. x 595 722 32 24 Zw 70 1.5 10 .times.
10.sup.22 2 65 .smallcircle. .smallcircle. .smallcircle. 590 713 32
25 Zw 85 1.1 5 .times. 10.sup.22 2 65 .smallcircle. .smallcircle.
.smallcircle. 567 691 33 26 Zw 70 1.2 5 .times. 10.sup.22 5 85
.smallcircle. .smallcircle. .smallcircle. 609 736 31 27 Zw 55 1.8 5
.times. 10.sup.22 6 25 .smallcircle. .smallcircle. .smallcircle.
598 611 33 28 Zw 65 1.3 5 .times. 10.sup.22 2 65 .smallcircle.
.smallcircle. .smallcircle. 593 725 32 Mechanical properties
Toughness evaluation test Absorbed energy Fracture Strength-vE HIC
Steel FATT.sub.85% (vE.sub.-20.degree. C.) unit balance CAR No.
(.degree. C.) (J) (.mu.m) (MPa J) (%) Remarks 1 -45 330 20 233640 0
Inv. ex. 2 -40 260 22 167440 4 Comp. ex. 3 -5 220 48 158620 6 Comp.
ex. 4 -45 250 20 167500 5 Comp. ex. 5 -30 305 25 216855 0 Inv. ex.
6 -25 285 28 199215 3 Inv. ex. 7 0 170 51 122740 3 Comp. ex. 8 -5
155 18 104470 4 Comp. ex. 9 -10 130 21 90090 5 Comp. ex. 10 -35 240
25 161040 1 Comp. ex. 11 -40 250 20 159000 5 Comp. ex. 12 -5 255 60
181050 8 Comp. ex. 13 -5 250 50 178750 4 Comp. ex. 14 0 245 55
173215 6 Comp. ex. 15 -20 190 29 122360 9 Comp. ex. 16 -35 270 24
201150 5 Inv. ex. 17 -20 320 28 235520 5 Inv. ex. 18 -15 150 45
105150 8 Comp. ex. 19 -10 140 50 100240 5 Comp. ex. 20 -40 250 22
160250 6 Comp. ex. 21 -35 280 60 198800 0 Inv. ex. 22 -40 310 85
214520 4 Inv. ex. 23 0 150 55 108300 5 Comp. ex. 24 -20 265 28
188945 6 Inv. ex. 25 -35 310 23 214210 7 Inv. ex. 26 -5 220 48
161920 9 Comp. ex 27 -10 210 51 128310 9 Comp. ex. 28 -30 270 24
195750 4 Inv. ex. PF: polygonal ferrite, P: pearlite, .alpha..sub.B
+ .alpha..sub.q: granular bainitic ferrite (.alpha..sub.B) and
quasi-polygonal ferrite (.alpha..sub.q)
[0142] The steels satisfying the requirements of the present
invention are the 10 steels of the Steel Nos. 1, 5, 6, 16, 17, 21,
22, 24, 25, and 28. These give high strength hot rolled steel
sheets for line pipe use excellent in ductile fracture arrest
performance having tensile strengths corresponding to the X80 grade
as materials before pipemaking characterized by containing
predetermined amounts of steel ingredients, having microstructures
of continuously cooled transformed structures in which precipitates
containing Nb of average sizes of 1 to 3 nm are dispersed at an
average density of 3 to 30.times.10.sup.22/m.sup.3, furthermore
having average circle equivalent diameters of precipitates
containing Ti nitrides contained in steel sheet with an
.alpha..sub.B and/or .alpha..sub.q of a volume fraction of 50% or
more of 0.1 to 3 .mu.m, and, furthermore, having at least half of
these in number contain complex oxides including Ca, Ti, and Al.
Furthermore, Steel Nos. 1, 5, and 21 performed light rolling, so
achieved CAR indicators of the sour resistance of the targeted 3%
or less.
[0143] The other steels are outside the scope of the present
invention for the following reasons. Steel No. 2 has a heating
temperature outside the scope of the present claim 4, so the
average size of the precipitates containing Nb (precipitation
strengthening particle size) and average density (precipitation
strengthening particle density) are outside the scope of claim 1
and a sufficient effect of precipitation strengthening cannot be
obtained, so the strength-vE balance is low.
[0144] Steel No. 3 has a heating temperature outside the scope of
the present claim 4, so the prior austenite grains coarsen, the
desirable continuously cooled transformed structure cannot be
obtained after transformation, and the FATT.sub.85% is a high
temperature.
[0145] Steel No. 4 has a heating holding time outside the scope of
the present claim 4, so a sufficient precipitation strengthening
effect cannot be obtained, so the strength-vE balance is low.
[0146] Steel No. 7 has a total reduction rate of the
non-recrystallization temperature range outside the scope of the
present claim 4, so the prior austenite grains coarsen, the
desirable continuously cooled transformed structure cannot be
obtained after transformation, and the FATT.sub.85% is a high
temperature.
[0147] Steel No. 8 has a total reduction rate of the
recrystallization region outside the scope of the present claim 4,
so the targeted microstructure etc. described in claim 1 cannot be
obtained, and the strength-vE balance is low.
[0148] Steel No. 9 has a final rolling temperature outside the
scope of the present claim 4, so the targeted microstructure etc.
described in claim 1 cannot be obtained, and the strength-vE
balance is low.
[0149] Steel No. 10 has a cooling rate outside the scope of the
present claim 4, so the target microstructure described in claim 1
cannot be obtained, and the strength-vE balance is low.
[0150] Steel No. 11 has a CT outside the scope of the present claim
4, so a sufficient precipitation strengthening effect cannot be
obtained, so the strength-vE balance is low.
[0151] Steel No. 12 has a time in the smelting process until
charging Al after Ti deoxidation outside the scope of the present
claim 4, so the dispersion of the oxides forming the nuclei of the
precipitates containing the Ti nitrides is insufficient, so the
targeted nitride size described in claim 1 becomes over 3 .mu.m and
the FATT.sub.85% is a high temperature.
[0152] Steel No. 13 has an amount of dissolved oxygen before
charging of Ti and an equilibrium amount of dissolved oxygen in the
smelting process outside the scope of the present claim 4, so the
targeted nitride size described in claim 1 becomes over 3 .mu.m and
the FATT.sub.85% is a high temperature.
[0153] Steel No. 14 has an order of charging of successive
deoxidizing elements in the smelting process outside the scope of
the present claim 4, so the targeted nitride size described in
claim 1 becomes over 3 .mu.m and the FATT.sub.85% is a high
temperature.
[0154] Steel No. 15 has a content of C etc. which is outside the
scope of the present claim 1, so the targeted microstructure is not
obtained, and the strength-vE balance is low.
[0155] Steel No. 18 has a content of C etc. which is outside the
scope of the present claim 1, so the targeted microstructure is not
obtained, and the strength-vE balance is low.
[0156] Steel No. 19 has a content of C etc. which is outside the
scope of the present claim 1, so the targeted microstructure is not
obtained, and the strength-vE balance is low.
[0157] Steel No. 20 has a content of C etc. which is outside the
scope of the present claim 1, so the targeted microstructure is not
obtained, and the strength is low.
[0158] Steel No. 23 has an order of charging of successive
deoxidizing elements in the smelting process outside the scope of
the present claim 4, so the targeted nitride size described in
claim 1 becomes over 3 .mu.m and the FATT.sub.85% is a high
temperature.
[0159] Steel No. 26 has a Ca content outside the scope of the
present claim 1, so the targeted nitride size described in claim 1
becomes over 3 .mu.m and the FATT.sub.85% is a high
temperature.
[0160] Steel No. 27 has V, Mo, Cr and Cu, and Ni contents outside
the scope of the present claim 1, so as a material, a tensile
strength corresponding to the X80 grade cannot be obtained.
INDUSTRIAL APPLICABILITY
[0161] By using the hot rolled steel sheet of the present invention
for electric resistance welded steel pipe and spiral steel pipe,
production of line pipe with a high strength of the API5L-X80
standard or more can be produced even with a relatively large sheet
thickness of for example half an inch (12.7 mm) even in artic
regions where tough fracture resistance is demanded. Furthermore,
due to the method of production of the present invention, the hot
rolled steel sheet for electric resistance welded steel pipe and
spiral steel pipe use can be stably produced inexpensively in large
amounts. Therefore, the present invention enables line pipe to be
laid easier under harsh conditions. We are confident that it will
greatly contribute to the construction of pipelines--which is key
to the global distribution of energy.
* * * * *