U.S. patent application number 11/475117 was filed with the patent office on 2006-10-26 for ferritic steel sheet concurrently improved in formability, high-temperature strength, high temperature oxidation resistance, and low temperature toughness.
Invention is credited to Yoshimoto Fujimura, Yoshiaki Hori, Manabu Oku, Takeshi Utsunomiya.
Application Number | 20060237102 11/475117 |
Document ID | / |
Family ID | 32032946 |
Filed Date | 2006-10-26 |
United States Patent
Application |
20060237102 |
Kind Code |
A1 |
Oku; Manabu ; et
al. |
October 26, 2006 |
Ferritic steel sheet concurrently improved in formability,
high-temperature strength, high temperature oxidation resistance,
and low temperature toughness
Abstract
A low-cost ferritic steel sheet possessing not only formability
enabling application to complexly configured automobile exhaust gas
passage components but also high-temperature strength,
high-temperature oxidation resistance and low-temperature toughness
as good as or superior to existing ferritic steels, which ferritic
steel sheet comprises, in mass percent, C: not more than 0.02%, Si:
0.7-1.1%, Mn: not more than 0.8%, Ni: not more than 0.5%, Cr: 8.0
to less than 11.0%, N: not more than 0.02%, Nb: 0.10-0.50%, Ti:
0.07-0.25%, Cu: 0.02-0.5%, B: 0.0005-0.02%, V: 0 (no
addition)-0.20%, one or both of Ca and Mg: 0 (no addition)-0.01% in
total, one or more elements among Y and rare earth elements: 0 (no
addition)-0.20% in total, and the balance of Fe and unavoidable
impurities, and satisfies 3 Cr+40 Si.gtoreq.61, Cr+10 Si.ltoreq.21,
and 420 C-11.5 Si+7 Mn+23 Ni-11.5 Cr-12 Mo+9 Cu-49 Ti-25 (Nb+V)-52
A+470 N+189.ltoreq.70.
Inventors: |
Oku; Manabu; (Shunan-shi,
JP) ; Hori; Yoshiaki; (Kudamatsu-shi, JP) ;
Fujimura; Yoshimoto; (Shunan-shi, JP) ; Utsunomiya;
Takeshi; (Tokyo, JP) |
Correspondence
Address: |
CLARK & BRODY
1090 VERMONT AVENUE, NW
SUITE 250
WASHINGTON
DC
20005
US
|
Family ID: |
32032946 |
Appl. No.: |
11/475117 |
Filed: |
June 27, 2006 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
10670284 |
Sep 26, 2003 |
|
|
|
11475117 |
Jun 27, 2006 |
|
|
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Current U.S.
Class: |
148/325 ;
420/60 |
Current CPC
Class: |
C22C 38/004 20130101;
C22C 38/50 20130101; C22C 38/04 20130101; C22C 38/54 20130101; C22C
38/48 20130101; C22C 38/42 20130101; C22C 38/02 20130101; C22C
38/46 20130101 |
Class at
Publication: |
148/325 ;
420/060 |
International
Class: |
C22C 38/18 20060101
C22C038/18 |
Foreign Application Data
Date |
Code |
Application Number |
Oct 8, 2002 |
JP |
JP2002-294433 |
Sep 11, 2003 |
JP |
JP2003-319733 |
Claims
1-8. (canceled)
9. A ferritic steel sheet concurrently improved in formability,
high-temperature oxidation resistance, high-temperature strength,
and low-temperature toughness comprising a composition, in mass
percent C: not more than 0.02%, Si: 0.7-1.1%, Mn: not more than
0.8%, Ni: not more than 0.5%, Cr: 8.0 to less than 11.0%, N: not
more than 0.02%, Nb: 0.10-0.50%, Ti: 0.07-0.25%, Cu: 0.02-0.5%, B:
0.0005-0.02%, V: 0(no addition)-0.20%, one or both of Ca and Mg: 0
(no addition)-0.01% in total, one or more elements among Y and rare
earth elements : 0 (no addition)-0.20% in total, and the balance of
Fe and unavoidable impurities, the composition satisfying all of
Equations (1)-(3): 3 Cr+40 Si.gtoreq.61 (1) Cr+10 Si.ltoreq.21 (2)
420 C-11.5 Si+7 Mn+23 Ni-11.5 Cr-12 Mo+9 Cu-49 Ti-25 (Nb+V)-52
Al+470 N+189.ltoreq.70 (3) which has a metallic structure obtained
by subjecting a hot rolled steel sheet having the composition to a
partial recrystallization treatment followed by cold rolling and
annealing, wherein the partial recrystallization treatment is
conducted by heating a cooled hot rolled steel sheet in the
temperature range of 850-1000.degree. C. to obtain a structure
which is 10-90 vol. % of recrystallized grains with a remainder
being unrecrystallized grains, and further wherein the annealing
after the cold rolling is conducted to obtain a totally
recrystallized structure.
10. A steel sheet according to claim 9, wherein the content of V is
0.01-0.20%.
11. A steel sheet according to claim 9, wherein the content of one
or both of Ca and Mg is 0.0003-0.01% in total.
12. A steel sheet according to claim 9, wherein the content of one
or more elements among Y and rare earth elements is 0.01-0.20% in
total.
13. A steel sheet according to claim 9, further including Mo: not
more than 0.50% and Al: not more than 0.10%.
14. A steel sheet according to claim 9, which is used as fabricated
into an automobile engine exhaust gas passage component.
15. A steel sheet according to claim 10, which is used as
fabricated into an automobile engine exhaust gas passage
component.
16. A steel sheet according to claim 11, which is used as
fabricated into an automobile engine exhaust gas passage
component.
17. A steel sheet according to claim 12, which is used as
fabricated into an automobile engine exhaust gas passage
component.
18. A steel sheet according to claim 13, which is used as
fabricated into an automobile engine exhaust gas passage
component.
19. A process for producing a ferritic steel sheet concurrently
improved in formability, high-temperature oxidation resistance,
high-temperature strength, and low-temperature toughness
comprising, in mass, percent C: not more than 0.02%, Si: 0.7-1.1%,
Mn: not more than 0.8%, Ni: not more than 0.5%, Cr: 8.0 to less
than 11.0%, N: not more than 0.02%, Nb: 0.10-0.50%, Ti: 0.07-0.25%,
Cu: 0.02-0.5%, B: 0.0005-0.02%, V: 0(no addition)-0.20%, one or
both of Ca and Mg: 0 (no addition)-0.01% in total, one or more
elements among Y and rare earth elements : 0 (no addition)-0.20% in
total, and the balance of Fe and unavoidable impurities, the
composition satisfying all of Equations (1)-(3): 3 Cr+40
Si.gtoreq.61 (1) Cr+10 Si.ltoreq.21 (2) 420 C-11.5 Si+7 Mn+23
Ni-11.5 Cr-12 Mo+9 Cu-49 Ti-25 (Nb+V)-52 Al+470 N+189.ltoreq.70 (3)
characterized in that the process comprises subjecting the hot
rolled steel sheet having said composition to partial
recrystallization treatment followed by cold rolling and annealing,
wherein said partial recrystallization treatment is conducted by
heating the steel sheet cooled after hot-rolling in the temperature
range 850-1000.degree. C. to obtain a structure which is 10-90 vol
% of it being accounted for by recrystallized grains and the
remainder of it being accounted for by un-recrystallized grains,
and said annealing after said cold rolling is conducted to obtain a
structure totally recrystallized.
20. The process according to claim 23, wherein the content of V is
0.01-0.20%.
21. The process according to claim 23, wherein the content of one
or both of Ca and Mg is 0.0003-0.01% in total.
22. The process according to claim 23, wherein the content of one
or more elements among Y and rare earth elements is 0.01-0.20% in
total.
23. A process according to claim 23, further including Mo: not more
than 0.50% and Al: not more than 0.10%.
Description
BACKGROUND OF THE INVENTION
[0001] 1. Field of the Invention
[0002] This invention relates to a ferritic steel sheet
concurrently improved in formability such as deep drawability,
stretch formability and the like, high-temperature strength,
high-temperature oxidation resistance, and low-temperature
toughness, particularly a steel sheet that, being usable in a
800-900.degree. C. high-temperature atmosphere, is suitable for
utilization in automobile engine exhaust gas passage
components.
[0003] 2. Background Art
[0004] Since ferritic stainless steels have a smaller thermal
expansion coefficient than austenitic stainless steels and are
excellent in thermal fatigue property and high-temperature
oxidation property, they are used as heat-resistant materials in
applications where thermal strain is an issue. Typical applications
include automobile engine exhaust gas passage components such as
exhaust manifolds, front pipes, catalyst carrier outer cylinders,
center pipes, mufflers and tailpipes.
[0005] A trend seen in recent automobile engines is to increase
exhaust gas temperature in order to improve exhaust gas
purification efficiency and output. This has increased the need for
high heat resistance (high-temperature strength and
high-temperature oxidation resistance) particularly in the exhaust
manifold, front pipe, outer cylinder of the catalyst carrier and
other components near the engine. A tendency for the shape of
exhaust gas passage components to become more complicated has also
recently emerged. This is especially notable in the exhaust
manifold and outer cylinder of the catalyst carrier, which are
being formed in complex configurations by various methods,
including mechanical pressing, servo pressing, spinning and
hydroforming. It is therefore not sufficient for the materials used
in these components to be good merely in stretching elongation and
bendability. They are now also required to be excellent in
formability such as typically deep drawability and stretch
formability, and the in-plane anisotropy of their workability has
to be small. In addition, owing to the fact that consideration
needs to be given to prevention of ductile fructuer and brittle
cracking during secondary and tertiary machining, they must also be
excellent in low-temperature toughness. Moreover, heat resistance
cannot be sacrificed in the interest of improving formability and
low-temperature toughness, because greater shape complexity
increases the likelihood of thermal fatigue fracture occurring
owing to concentration of thermal strain at a single location
during engine starting and stopping and also boosts vulnerability
to abnormal oxidation caused by the material temperature rising
locally.
[0006] SUH409L and SUS430J1L are known as ferritic stainless steels
having high heat resistance. SUH409L is commonly used as an exhaust
gas passage component material because of its good workability and
low-temperature toughness. However, the level of its heat
resistance makes it unsuitable for applications in which the
material temperature exceeds 800.degree. C. It is also lacks
sufficient deep drawability for application to components with
complicated shapes. SUS430J1L has excellent heat resistance that
makes it usable at 900.degree. C. But it is hard and poor in
formability.
[0007] In light of the foregoing, the following heat-resistant
ferritic steels have been developed.
[0008] Patent Reference 1 among the references listed below teaches
a ferritic heat-resistant stainless steel with a Cr level of
17.0-25.0%. This steel is added with Mo and Cu in combination to
improve high-temperature strength and with Mn to suppress scale
spalling. Degradation of impact value by Mo is overcome to some
degree by combined addition of Cu and Ni. However, the steel's
formability is not adequate to cope with the needs of complexly
shaped exhaust gas passage components. And its high Cr content
makes it disadvantageous from the cost aspect.
[0009] Patent Reference 2 teaches a 13% Cr ferritic stainless steel
that exhibits heat resistance at least as good as an 18% Cr
ferritic stainless steel and also has improved high-temperature
salt corrosion property. In this steel, high-temperature strength
is increased by ensuring the presence of solid solution Nb,
high-temperature oxidation property is improved by liberal addition
of Mn and Si, and NaCl-induced hat corrosion resistance is improved
by the Si. As no particular consideration is given to improvement
of formability and low-temperature toughness, however, the steel
cannot adequately respond to the recent harsh requirements
mentioned earlier.
[0010] Patent Reference 3 teaches a Nb-containing heat-resistant
ferritic stainless steel with a Cr level of 11.0-15.5% that is
aimed at improving high-temperature oxidation resistance and scale
adherence. These properties are markedly improved by strictly
constraining Mn/Sn within the range of 0.7-1.5. This Patent
Reference also teaches improvement of low-temperature toughness and
workability by Cu addition. Regarding workability, for example, it
presents data showing that no cracking occurred in a 180-degree
bending test. In light of the fact that demands regarding the shape
of exhaust gas passage components are becoming more and more
challenging, however, the materials used in these components have
come to require excellent formability that is compatible with
various forming methods (discussed later). Regarding this point,
the steel of Patent Reference 3 pays no attention to deep
drawability and other drawability-related stretch formability and,
as such, cannot be considered capable of responding to today's
severe requirements. In addition, its Cr content of 11.0% or
greater is on the level required in a stainless steel, which is
inconsistent with the desire to reduce cost by lowering the Cr
content in exhaust gas passage components that do not necessarily
require use of stainless steels.
[0011] Patent Reference 4 teaches a ferritic stainless steel for
exhaust manifolds that contains Cr at 11-14%. This is a steel
enhanced in high-temperature strength by positive addition of Si to
a Nb-containing steel. Its high temperature can be considered to be
the same as the steel of Patent Reference 3. As it does not give
consideration to improving formability and low-temperature
toughness beyond the level of the prior art, however, it is
incapable of thoroughly responding to the harsh demands placed on
steels in recent years. It also needs further reduction in Cr
level.
[0012] Patent Reference 5 teaches a ferritic heat-resistant steel
for engine exhaust gas passage components having a Cr content of
8.0-10.0%. This is a steel that achieves better heat resistance
than SUH409L while also reducing cost through low Cr content. This
reference further teaches that Cu is effective for improving both
low-temperature toughness and workability. Regarding workability,
it was for example found to possess ductility on a par with SUH409L
in tensile tests at room temperature. As it is not aimed at
improvement of the in-plane anisotropy of ductility or deep
drawability, however, it leaves unresolved the problem of imparting
formability thoroughly matched to the needs of various forming
methods (discussed later). Nor does it offer a method for
consistently imparting excellent low-temperature toughness. Patent
Reference 5 can therefore not be viewed as sufficiently responding
to the recent severe requirements with respect to exhaust gas
passage components.
[0013] Patent References 6 and 7 teach ferritic steels with a Cr
content of 10 to less than 15% that improve the corrosion
resistance against condensed moisture needed by mufflers and other
low-temperature components and also the high-temperature strength
needed by exhaust manifolds and other high-temperature components.
But they merely assess workability in terms of proof stress, while
offering nothing specific with regard to high-temperature oxidation
resistance. Patent References 6 and 7 are not directed to the goal
of concurrently and consistently improving high-temperature
oxidation resistance and formability with good reproducibility, and
are silent regarding methods for achieving this objective. From the
viewpoint of fabrication into exhaust gas passage components of
various complex shapes, therefore, the steels taught by Patent
References 6 and 7 cannot be considered steels that fully meet all
formability requirements.
[0014] Patent Reference 1 [0015] JP-A-Hei 3(1991)-274245 (p3, upper
right column, line 1-p4, upper right column, line 9)
[0016] Patent Reference 2 [0017] JP-A-Hei 5(1993)-125491
(paragraphs 0012-0016)
[0018] Patent Reference 3 [0019] JP-A-Hei 7(1995)-11394 (paragraphs
0014-0021, 0028, 0029, Table 6, FIG. 1)
[0020] Patent Reference 4 [0021] JP-A-Hei 7(1995)-145453
(paragraphs 0011-0021)
[0022] Patent Reference 5 [0023] JP-A-Hei 10(1998)-147848
(paragraphs 0003-0005, 0014)
[0024] Patent Reference 6 [0025] JP-A-Hei 10(1998)-204590
(paragraphs 0026-0036, 0072)
[0026] Patent Reference 7 [0027] JP-A-Hei 10(1998)-204591
(paragraphs 0028-0037, 0074)
[0028] As explained in the foregoing, steel sheet for automobile
exhaust gas passage components is now being required to contribute
to greater component design freedom by offering excellent
formability that enables fabrication into complicated shapes by a
diversity of forming methods. But this need should best be met
while maintaining high-temperature strength and high-temperature
oxidation resistance at 800-900 .degree. C. on a par with
SUS430J1L, and also ensuring excellent low-temperature toughness.
As can be seen from the aforesaid Patent References, however, no
steel sheet has yet been developed that is concurrently improved to
a high degree in all of formability, high-temperature strength,
high-temperature oxidation resistance, and low-temperature
toughness.
[0029] An object of the present invention is to provide a new
ferritic heat-resistant steel that concurrently offers excellent
formability enabling ready application to complexly configured
automobile exhaust gas passage components, excellent
high-temperature strength and high-temperature oxidation resistance
enabling it to withstanding use at 900.degree. C. and excellent
low-temperature toughness having an energy transition temperature
of minus 50 .degree. C. or lower, and that is lowered in cost by
reducing Cr content to below 11 mass percent.
SUMMARY OF THE INVENTION
[0030] The inventors carried out a study to determine why excellent
formability, high-temperature strength, high-temperature oxidation
resistance, and low-temperature toughness have not yet been
concurrently achieved in a steel sheet. Based on our findings, we
concluded that a major cause is the fact that no means has been
found for concurrently establishing with high stability and good
reproducibility the properties of formability and high-temperature
oxidation resistance among the foregoing properties. Then, in an
ensuing in-depth study, the inventors discovered that when the
austenite balance is adjusted in the manner of Equation (3) set out
below, then, as shown by Equation (2) set out below, a region of Si
and Cr content exists in which both excellent formability and
excellent high-temperature oxidation resistance can be concurrently
established.
[0031] Moreover, in evaluating workability into complexly
configured exhaust gas passage components, the property of deep
drawability among the different aspects of formability cannot be
disregarded. It was found that an effective way to improve the deep
drawability of a heat resistance ferritic steel added with Nb is to
add Ti in combination with the Nb. The inventors further learned
that deep drawability (average plastic strain ratio r.sub.AV) and
the in-plane anisotropy thereof (plastic anisotropy .DELTA.r) can
be improved by partially recrystallizing the hot-rolled sheet.
[0032] It should be noted, however, that Ti addition degrades
low-temperature toughness. It was found that combined addition of
Cu and B more effectively improved the low-temperature toughness
than did addition of Cu alone.
[0033] When the amount of added Cu was progressively increased,
however, an abnormal oxidation-inducing phenomenon was observed to
appear abruptly. Further, an appropriate range of Cu enabling
concurrent improvement of low-temperature toughness and
high-temperature oxidation resistance was discovered.
[0034] The present invention was accomplished based on the
foregoing findings.
[0035] Specifically, the aforesaid object is achieved by a ferritic
steel sheet concurrently improved in formability, high-temperature
oxidation resistance, high-temperature strength, and
low-temperature toughness comprising, in mass percent, C: not more
than 0.02%, Si: 0.7-1.1%, Mn: not more than 0.8%, Ni: not more than
0.5%, Cr: 8.0 to less than 11.0%, N: not more than 0.02%, Nb:
0.10-0.50%, Ti: 0.07-0.25%, Cu: 0.02-0.5%, B: 0.0005-0.02%, V: 0
(no addition)-0.20%, preferably 0.01-0.20%, one or both of Ca and
Mg: 0 (no addition)-0.01% in total, preferably 0.0003-0.01% in
total, one or more elements among Y and rare earth elements: 0 (no
addition)-0.20% in total, preferably 0.01-0.20% in total, and the
balance of Fe and unavoidable impurities, and having a chemical
composition satisfying all of Equations (1)-(3): 3Cr+40Si.gtoreq.61
(1) Cr+10Si.ltoreq.21 (2) 420C-11.5 Si+7 Mn+23 Ni-11.5 Cr-12 Mo+9
Cu-49 Ti-25 (Nb+V)-52 Al+470 N+189.ltoreq.70 (3)
[0036] The steel sheet may further include Mo: not more than 0.50%
and Al: not more than 0.10%.
[0037] Each element symbol in Equations (1)-(3) is replaced by a
value representing the content of the element in mass percent. In
Equation (3), symbols of elements not contained are replaced by
zero.
[0038] In the present invention, the aforesaid steel sheet may have
a metallic structure obtained by cold rolling and annealing a
partially recrystallized hot-rolled sheet. A "partially
recrystallized hot-rolled sheet" as termed here means a hot-rolled
sheet 10-90 vol % of whose structure is accounted for by
recrystallized grains and the remainder of which is accounted for
by un-recrystallized grains. The amount of recrystallized grains
present can be ascertained by observing a cross-section of the
hot-rolled sheet with an optical microscope. By "hot-rolled sheet"
is meant the steel sheet that has been subjected to hot rolling and
may have been subjected to heat treatment after hot rolling but has
not be subjected to cold rolling. The final metallic structure
obtained by conducting cold rolling and annealing is totally
recrystallized.
[0039] In the present invention, further, the aforesaid steel sheet
may have a metallic structure obtained by cold rolling and
annealing a tatally recrystallized hot-rolled sheet. A "tatally
recrystallized hot-rolled sheet" as termed here means a hot-rolled
sheet more than 90 vol % of whose structure is accounted for by
recrystallized grains.
[0040] The steel sheet provided by the present invention is
particularly one used as fabricated into an automobile engine
exhaust gas passage component.
BRIEF EXPLANATION OF THE DRAWINGS
[0041] FIG. 1 is a graph showing how Ti content and difference
between partial and complete recrystallization after hot rolling
affected r value (r.sub.D) at 45 degrees to the rolling direction
in ferritic steels whose basic composition was 10 Cr-0.9 Si-0.3
Nb-0.1 V-0.1 Cu.
[0042] FIG. 2 is a graph showing how Cu content affected energy
transition temperature and amount of oxidation increase after
900.degree. C..times.200 hour heating in the atmosphere in ferritic
steels whose basic composition was 10 Cr-0.9 Si-0.3 Nb-0.1 V-0.001
B.
[0043] FIG. 3 is a graph showing how Cr content and Si content
affected high-temperature oxidation resistance and formability in
ferritic steels whose basic composition was 8 to 14 Cr-0.5 to 1.0
Si-0.3 Nb-0.1 Ti-0.1 V-0.1 Cu.
[0044] FIG. 4 is a graph showing how elongation at 45 degrees to
the rolling direction in a room-temperature tensile test varied
with AM value (AM=420 C-11.5 Si+7 Mn+23 Ni-11.5 Cr-12 Mo+9
Cu-49Ti-25 (Nb+V)-52 Al+470 N+189) in ferritic steels whose basic
composition was 8 to14 Cr-0.5 to 1.0 Si-0.3 Nb-0.1 Ti-0.1 V-0.1 Cu
and that satisfied Equations (1) and (2).
DESCRIPTION OF THE PREFERRED EMBODIMENTS
[0045] FIG. 1 shows how Ti content and difference between partial
and complete recrystallization after hot rolling affected r value
(r.sub.D) at 45 degrees to the rolling direction in ferritic steels
whose basic composition was 10 Cr-0.9 Si-0.3 Nb-0.1 V-0.1 Cu. The
partially recrystallized hot-rolled sheets were 4.0 mm-thick
hot-rolled sheets heat treated at 700-1000.degree. C. for 1 minute
to have a structure 10-90 vol % of which was accounted for by
recrystallized grains. The totally recrystallized hot-rolled sheets
were 4.0 mm-thick hot-rolled sheets heat treated at about
1050.degree. C. for 1 minute. The hot-rolled sheets were cold
rolled to 2.0 mm and totally recrystallized by annealing at
1050.degree. C., whereafter tensile test pieces were cut from them.
As can be seen in FIG. 1, the r.sub.D value rose sharply when Ti
was added to a content of 0.07 mass % or more. Moreover, the
r.sub.D value was improved markedly over the full range of Ti
content by partially recrystallizing the steel sheets after hot
rolling.
[0046] While it is not altogether clear what caused these
improvements, it is likely that Ti, whose carbonitride producing
capability is stronger than that of Nb, fixed C and N, and since
this reduced solid solution C and N, the purity of the matrix was
increased to a high level that promoted development of (111) plane
aggregate texture favorable for improvement of workability during
recrystallization at final annealing. This effect is thought to
manifest itself when the Ti content reaches 0.07 mass % or more. On
the other hand, it is likely that partial recrystallization of the
hot-rolled sheet uniformly produced fine Nb--Ti precipitates that
suppressed development of (100) plane aggregate texture thought to
be detrimental to workability improvement and promoted development
of (111) plane aggregate texture.
[0047] FIG. 2 shows how Cu content affected energy transition
temperature and amount of oxidation increase after 900.degree.
C..times.200 hour heating in the atmosphere in ferritic steels
whose basic composition was 10 Cr-0.9 Si-0.3 Nb-0.1 V-0.001 B. The
specimens used were totally recrystallized steel sheets obtained by
cold rolling partially crystallized 4.0 mm-thick hot-rolled sheets
to a thickness of 2.0 mm and then finally annealing them at
1050.degree. C. The energy transition temperature was ascertained
by a Charpy impact test. No. 5 test pieces (2 mm width) were taken
in accordance with JIS Z 2202 so that the impact direction would be
parallel to the rolling direction, the test was conducted at minus
100 to plus 25.degree. C. in accordance with JIS Z 2242, and the
energy transition temperature was determined from the relationship
between the test temperature and the absorbed energy. Oxidation
amount increase was determined in accordance with JIS Z 2281 by
measuring the weight increase of the test piece when it was
maintained at 900.degree. C. in the atmosphere continuously for 200
hours. As can be seen form FIG. 2, in a ferritic steel containing
an appropriate amount of B, slight addition of Cu at around 0.02
mass % effectively produced an improvement in low-temperature
toughness. However, it was newly discovered that when the amount
added exceeds 0.5 mass %, the oxidation resistance degenerates
rapidly.
[0048] The reasons for the foregoing observations has not yet been
fully ascertained, it is likely regarding low-temperature toughness
that occurrence of twin crystal, a cause of low-temperature
brittleness, was suppressed, while it is likely regarding
occurrence of abnormal oxidation that destabilization of the matrix
phase balance caused by Cr and Si oxidation was aggravated by
Cu.
[0049] FIG. 3 shows how Cr content and Si content affected
high-temperature oxidation resistance and formability in ferritic
steels whose basic composition was 8 to 14 Cr-0.5 to 1.0 Si-0.3
Nb-0.1 Ti-0.1 V-0.1 Cu. The specimens were prepared by the process
explained regarding FIG. 2. The 0.2% proof stress at 45 degrees to
the rolling direction determined in a room-temperature tensile test
was used as an index of formability. When this value exceeded 300
MPa, it was judged that as a material for exhaust gas passage
components the steel was basically lacking in formability capable
of meeting the needs of various forming methods. As can be seen
from FIG. 3, low content of Cr and Si resulted in occurrence of
abnormal oxidation during 900.degree. C..times.100 hour heating in
the atmosphere. On the other hand, formability deteriorated with
increasing Cr and Si content. There was, however, found to be a
Cr--Si combination region in which satisfactory high-temperature
oxidation resistance at 900.degree. C. and satisfactory formability
can both be obtained. The existence of such a region was not known
heretofore. Therefore, despite the development of various ferritic
heat-resistant steels, all were inferior in either high-temperature
oxidation resistance or formability and no steel emerged that could
satisfy both requirements reliably and with good
reproducibility.
[0050] The region in which excellent high-temperature oxidation
resistance and formability can both be obtained is that in which
the blank circle plots are present in FIG. 3. This region is
delimited by Equations (1) and (2): 3 Cr+40 Si.gtoreq.61 (1) Cr+10
Si.gtoreq.21 (2)
[0051] FIG. 4 shows how elongation in a room-temperature tensile
test varied with AM value (AM=420 C-11.5 Si+7 Mn+23 Ni-11.5 Cr-12
Mo+9 Cu-49 Ti-25 (Nb+V)-52 Al+470 N+189) in ferritic steels whose
basic composition was 8 to 14 Cr-0.5 to 1.0 Si-0.3 Nb-0.1 Ti-0.1
V-0.1 Cu and that satisfied Equations (1) and (2). AM value
represents the balance between ferrite phase and austenite phase.
As can be seen from FIG. 4, high ductility was obtained only in the
region of an AM value not more than 70 and degenerated
precipitously when AM exceeded 70. Thus formability and
high-temperature oxidation resistance are concurrently improved
only when Equation (3) is satisfied in addition to Equations (1)
and (2): 420 C-11.5Si+7Mn+23Ni-11.5Cr-12Mo+9Cu-49Ti-25 (Nb+V)-52
Al+470 N+189.ltoreq.70 (3)
[0052] Features defining the present invention will now be
explained.
[0053] C and N are generally effective for improving creep
strength, creep rupture strength and other high-temperature
strength properties. In a ferritic steel, however, low-temperature
toughness is degraded by a high content of C and N. This makes it
necessary to increase the amount of added Nb and Ti in order to
stabilize C and N as carbonitrides. The result is higher cost. On
the other hand, an attempt to markedly lower C and N content makes
steelmaking more onerous, which also increases cost. Through
various studies it was found that in the present invention a
content of up to 0.02 mass % is permissible for both C and N. It
should be noted, however, that when the amount of added Ti and Nb
is appropriately set, particularly good formability and heat
resistance are obtained when the amount of (C+N) is in the range of
0.01-0.02 mass %. The total content of C and N combined is
therefore preferably 0.01-0.02 mass %.
[0054] Si and Cr are both very effective for improving
high-temperature oxidation property but they also harden the steel.
In order to establish both excellent formability and excellent
high-temperature oxidation resistance, the Si and Cr contents need
to be controlled to within the range satisfying Equations (1) and
(2), as explained earlier with reference to FIG. 3. In addition to
being controlled based on these equations, however, upper and lower
limits of Si and Cr content are further defined from the standpoint
of ensuring good corrosion resistance and low-temperature
toughness. To wit, the minimum required level of corrosion
resistance exemplified by SUH409L cannot be achieved when the Si
and Cr contents are too small, whereas the low-temperature
toughness level of the SUH409L steel cannot be realized when their
contents are too high. Si content is therefore defined as 0.7-1.1
mass %. A more preferable range of Si content is 0.8-1.0 mass %. Cr
content is defined as 8.0 to less than 11.0%. A more preferable
range of Cr content is 9.0 to less than 11.0% and a still more
preferable range of Cr content is 9.0 to less than 10.0%.
[0055] Mn hardens the steel and degrades its low-temperature
toughness and formability when added in excess. Particularly in the
composition system of the present invention, excessive addition of
Mn is liable to adversely affect high-temperature oxidation
resistance by causing generation of austenitic phase during hot
use. The upper limit of Mn content is therefore defined as 0.8 mass
%. In the composition system of the present invention, particularly
when excellent scale adherence at the 900.degree. C. level is
required, Mn is preferably added to within the content range of
0.2-0.8 mass %.
[0056] Ni is effective for improving low-temperature toughness but
hardens the steel and degrades its formability when added in
excess. Moreover, in the composition system of the present
invention, excessive addition of Ni is, like excessive addition of
Mn, liable to degrade high-temperature oxidation resistance by
causing generation of austenitic phase during hot use. The upper
limit of Ni content is therefore defined as 0.5 mass %.
[0057] Nb is very effective for improving high-temperature
strength. Since Ti is added in the present invention, substantially
no Nb is fixed to C and N, so that essentially all added Nb can be
considered to work effectively toward enhancing high-temperature
strength. This effect manifests itself at a content of not less
than 0.10 mass %. On the other hand, excessive Nb addition degrades
formability and low-temperature toughness. The Nb content is
therefore defined as 0.10-0.50 mass %. In order to obtain still
higher formability and high-temperature strength, a Nb content in
the range of 0.10-0.40 mass % is preferable.
[0058] Ti fixes C and N and is generally known to improve grain
boundary corrosion resistance. In this invention, however, it is a
very important element for improving formability (particularly deep
drawability). The formability improving effect of Ti appears
conspicuously at content of not less than 0.07 mass % (see FIG. 1).
However, excessive Ti addition degrades toughness and adversely
affects product surface properties. Ti content is therefore defined
as 0.07-0.25 mass %. For obtaining a high level of high-temperature
strength, Ti is preferably added to satisfy Ti.gtoreq.6 (C+N). In
order to obtain a product with surface properties as good as or
better than SUH409L, Ti is preferably added to a content of not
more than 0.20 mass %.
[0059] Mo is effective for increasing high-temperature strength but
makes the steel brittle when present at a high content. In
addition, Mo is very expensive. While adequate heat resistance can
be secured without Mo addition by optimizing the contents of other
constituent elements, Mo addition is advantageous in that it
increases the freedom of composition design. When Mo is
incorporated, its content is preferably not more than 0.50 mass %.
When heat resistance is a bigger concern than cost, Mo can be added
in excess of 0.5 mass % but should not be added in excess of 3.0
mass %, the level beyond which an extreme decline in
low-temperature toughness occurs.
[0060] Cu improves low-temperature toughness. In order to markedly
improve low-temperature toughness to the level required in exhaust
gas passage components, however, it is important to incorporate not
less than 0.02 mass % of Cu in combination with B (discussed
later). When the Cu content exceeds 0.5 mass %, however,
high-temperature oxidation resistance degenerates sharply (see FIG.
2). Cu content is therefore defined as 0.02-0.5 mass % in the
present invention.
[0061] V, like Nb and Ti, is a carbonitride-forming element that is
effective for improving grain boundary corrosion resistance and the
toughness of sites affected by welding heat. Moreover, like Nb, V
contributes to high-temperature strength improvement in the solid
solution state. This effect is particularly pronounced when V is
present together with Nb. In addition, V is thought to be effective
for improving high-temperature oxidation resistance. However, a V
content in excess of 0.20 mass % degrades workability and
low-temperature toughness. When V is added, therefore, its content
must be kept to not more than 0.20 mass %. For thoroughly obtaining
the foregoing effects of V, it is preferably added in the range of
0.01-0.20 mass %.
[0062] Al is highly effective for improving high-temperature
oxidation resistance. However, the composition according to the
present invention is designed to enable excellent high-temperature
oxidation resistance even without incorporation of Al. Excessive Al
addition degrades formability, weldability and low-temperature
toughness. Moreover, deoxidation by Al is not particularly
necessary because the present invention calls for addition of Ti
and Si. When Al is incorporated, it must be added at no more than
0.10 mass %. When adding Al in a case where formability,
weldability and low-temperature toughness are particularly
important, the Al content is preferably restricted to not more than
0.07 mass %.
[0063] B suppresses low-temperature brittleness and secondary work
brittleness in a ferritic steel also containing Nb and Ti. This
effect was found to be pronounced when B is added in combination
with Cu. In order to thoroughly improve low-temperature toughness,
B needs to be added at not less than 0.0005 mass %. On the other
hand, excessive B addition beyond 0.02 mass % leads to generation
of borides that degrade formability and degrade rather than improve
low-temperature toughness. In the present invention, B is
incorporated at 0.0005-0.02 mass % together with Cu at 0.02-0.5
mass %.
[0064] Ca and Mg have strong binding force with S and therefore
reduce the amount of MnS generation to improve corrosion
resistance. In addition, Ca and Mg are elements that in themselves
effectively work to improve high-temperature oxidation resistance.
When importance is attached to corrosion resistance and
high-temperature oxidation resistance, therefore, these elements
can be added as required. However, addition in large amounts
increases inclusions that degrade low-temperature toughness and
formability. When one or both of Ca and Mg are added, therefore,
the combined content thereof needs to be held to not more than 0.01
mass %. In order to bring out the effect of Ca and Mg addition
strongly, the total of the Ca and Mg contents should preferably
made 0.003-0.01 mass %.
[0065] Y and REMs (rare earth elements) such as La and Ce stabilize
the chromium oxide coating that forms on the steel surface and, by
enhancing the adherence between the steel matrix and the oxide
coating, manifestly improve the high-temperature oxidation
resistance of the steel sheet. When high-temperature oxidation
resistance is a major concern, therefore, these elements can be
added as required. However, addition in large amounts not only
degrades formability and low-temperature toughness but also
promotes generation of inclusions that may become starting points
of abnormal oxidation, meaning that high-temperature oxidation
resistance is degraded rather than improved. Therefore, when one or
more elements selected from among Y and rare earth elements are
added, the combined amount thereof must be made not more than 0.20
mass %. For maximizing the effect of Y and REM addition, one or
more elements selected from among these elements *should preferably
be added to a combined total of 0.01-0.20 mass %.
[0066] As additional elements, one or more of Zr, Hf, Ta, W, Re and
Co can be included for their ability to improve high-temperature
strength. Since excessive addition of these elements hardens the
steel, however, they must, when incorporated, be added to a
combined content of not more than 3.0 mass %. The preferable amount
is not more than 0.5 mass % in total.
[0067] The content of P, S, O, Zn, Sn, Pb and other common impurity
elements is preferably reduced to the lowest level possible in
order to ensure good formability and low-temperature toughness.
Specifically, these elements should, at the most lenient, be
restricted to P: not more than 0.04 mass %, S: not more than 0.03
mass %, O: not more than 0.02 mass %, Zn: not more than 0.10 mass
%, Sn: not more than 0.10 mass %, and Pb: not more than 0.10 mass
%. At the actual production site, more severe restrictions can be
imposed in accordance with the product quality desired.
[0068] As explained earlier, Equations (1)-(3) define the
composition range required for concurrent improvement of
formability and high-temperature oxidation resistance. Although no
particular lower limit is specified for the value (AM value) of the
left side of Eqution (3), a steel with a low AM value ordinarily
contains liberal amounts of ferrite generating elements like Si,
Cr, Mo, Ti, Nb, V and Al. When large amounts of these elements are
contained, formability and low-temperature toughness degenerate.
Studies showed that it is preferable to regulate the constituents
so that the AM value is 40 or higher.
[0069] Satisfying the aforesaid chemical composition concurrently
improves formability, high-temperature oxidation resistance,
high-temperature strength and low-temperature toughness.
[0070] In order to further improve the formability realized in this
way, it is highly effective to subject the hot-rolled sheet to
partial recrystallization treatment followed by cold rolling and
annealing. Specifically, the r value, an index of deep drawability,
can be markedly improved by the steps of making a hot-rolled sheet
10-90 vol % of whose structure is accounted for by recrystallized
grains and the remainder of which is accounted for by
un-recrystallized grains, cold rolling the hot-rolled sheet, and
totally recrystallizing it by annealing (see FIG. 1). The steel
sheet having the metallic structure obtained in this manner
possesses formability fully capable of responding to the
increasingly severe shape requirements of today's exhaust gas
passage components.
[0071] Partial recrystallization of the hot-rolled sheet can be
carried out directly during hot rolling process or by heating
conducted between hot rolling and cold rolling.
[0072] Partial recrystallization during hot rolling can, for
instance, be conducted by hot rolling in the temperature range of
950-1250.degree. C., coiling, and cooling in the coiled state.
Optimum conditions can be selected in accordance with the facility
specifications and the hot-rolling pass schedule. Partial
recrystallization by heating after hot rolling can be conducted,
for example, by heating the steel sheet cooled after hot rolling in
the temperature range of 850-1000.degree. C. The heating can be
carried out at any stage before cold rolling.
[0073] The hot-rolled sheet partially recrystallized by one of the
foregoing methods is then totally recrystallized by annealing. The
cold rolling is conducted at a reduction in the range of, for
example, 30-90%. When the steel sheet is to be used in an
automobile exhaust gas passage component, the final sheet thickness
is adjusted to, for example, about 0.4-1.2 mm. The annealing
temperature is preferably in the range of, for instance,
950-1150.degree. C. The ferritic steel sheet obtained is excellent
in formability and low-temperature toughness and these properties
are retained even after fabrication into welded steel tube.
[0074] In case an article having been formed should have a
beautiful surface appearance and make much of good looking outer
surface, it is preferable to use the totally recrystallized
hot-rolled sheet to form the article. The totally recrystallized
hot-rolled sheet can be obtained by subjecting the hot-rolled sheet
into a heat treatment at the temperatures between 950 and
1100.degree. C.
EXAMPLE 1
[0075] Ferritic steels having the chemical compositions shown in
Tables 1 and 2 were made using a high-frequency vacuum melting
furnace and cast into 30-Kg ingots. The ingots were hot-forged and
then hot-rolled into 4.0 mm hot-rolled sheets. The hot rolling was
conducted at a hot-rolling temperature of 700-1250.degree. C. and a
draft (rolling reduction) per pass of about 30%. Each hot-rolled
sheet was water cooled and then held at 900-1000.degree. C. for 1
minute. The cross-sectional metallic structure of the hot-rolled
sheet was observed with an optical microscope. Recrystallized
grains were found to account for 10-90 vol. % of every specimen,
the balance being un-recrystallized structure. It was thus
ascertained that partial recrystallization had been achieved. The
partially recrystallized hot-rolled sheets were cold rolled to a
thickness of 2mm and thereafter totally recrystallized by annealing
for 1 minute at 1050.degree. C. to afford cold-rolled annealed
sheets. Nos. 1-21 in Table 1 are ferritic steels satisfying the
chemical composition defined by the present invention. Nos. 22-31
in Table 2 are comparative steels not meeting compositional
requirements of the present invention. Among the comparative
steels, No. 22 corresponds to SUH409L and No. 23 to SUS430J1L
TABLE-US-00001 TABLE 1 (Mass %) 3Cr + Cr + AM No. C Si Mn Ni Cr Mo
Cu Ti Mb V Al B N Other 40Si 10Si value Inven- 1 0.002 0.93 0.04
0.02 8.06 tr. 0.02 0.14 0.31 0.03 0.08 0.0005 0.003 -- 61.4 17.4
69.3 tion 2 0.009 0.78 0.78 0.09 10.11 0.01 0.10 0.09 0.33 0.05
0.06 0.0010 0.008 -- 61.5 17.9 62.6 steels 3 0.010 1.00 0.20 0.11
9.52 tr. 0.05 0.20 0.20 0.20 0.01 0.0050 0.009 -- 68.6 19.5 60.5 4
0.009 0.79 0.78 0.10 10.45 0.05 0.12 0.14 0.29 0.15 0.01 0.0037
0.007 -- 63.0 18.4 56.7 5 0.009 0.88 0.77 0.10 9.60 tr. 0.11 0.14
0.28 0.04 tr. 0.0026 0.007 -- 64.0 18.4 69.4 6 0.010 0.71 0.08 0.10
10.99 tr. 0.10 0.14 0.34 0.06 0.01 0.0011 0.007 -- 61.4 18.1 48.3 7
0.009 0.80 0.18 0.10 10.27 0.14 0.10 0.14 0.32 0.03 0.03 0.0030
0.007 -- 62.8 18.3 54.4 8 0.009 0.80 0.18 0.10 10.40 tr. 0.48 0.15
0.33 0.07 0.02 0.0023 0.007 -- 63.2 18.4 56.8 9 0.011 0.80 0.18
0.10 10.81 0.46 0.11 0.15 0.33 0.05 tr. 0.0012 0.006 -- 64.4 18.8
45.1 10 0.006 0.91 0.22 0.02 10.46 0.03 0.21 0.24 0.25 0.04 0.01
0.0023 0.007 -- 67.8 19.6 48.1 11 0.005 0.92 0.31 0.12 10.49 0.01
0.15 0.12 0.49 0.04 0.01 0.0017 0.008 -- 68.3 19.7 50.2 12 0.012
0.85 0.44 0.11 10.52 0.01 0.12 0.13 0.11 0.19 0.01 0.0020 0.009 --
65.6 19.0 59.7 13 0.014 0.89 0.21 0.47 10.36 tr. 0.06 0.11 0.25
0.06 0.01 0.0091 0.009 -- 66.7 19.3 68.9 14 0.018 0.95 0.22 0.10
10.38 tr. 0.06 0.09 0.38 0.05 0.01 0.0014 0.003 -- 69.1 19.9 56.4
15 0.002 1.01 0.26 0.11 10.54 0.42 0.10 0.15 0.39 0.03 0.01 0.0016
0.020 -- 72.0 20.6 48.3 16 0.004 1.09 0.19 0.12 10.10 0.10 0.11
0.15 0.28 0.03 0.02 0.0011 0.009 -- 73.9 21.0 54.0 17 0.005 0.98
0.15 0.10 9.88 0.11 0.10 0.14 0.29 0.03 tr. 0.0021 0.008 Ca: 0.004
68.8 19.7 58.0 18 0.008 0.77 0.12 0.10 10.09 tr. 0.09 0.14 0.29
0.02 tr. 0.0037 0.007 Mg: 0.001 61.1 17.8 60.1 19 0.003 0.88 0.04
0.11 8.61 tr. 0.05 0.19 0.33 0.01 tr. 0.0019 0.007 La: 0.01 61.0
17.4 69.9 Ce: 0.04 20 0.007 0.71 0.10 0.11 10.87 tr. 0.10 0.11 0.25
0.02 tr. 0.0015 0.008 Y: 0.03 61.0 18.0 54.5 21 0.003 0.75 0.61
0.12 10.25 tr. 0.11 0.09 0.20 tr. tr. 0.0005 0.002 -- 60.8 17.8
63.3 tr.: Trace
[0076] TABLE-US-00002 TABLE 2 No. C Si Mn Ni Cr Mo Cu Ti Nb
Comparative 22 0.010 0.51* 0.15 0.11 10.92 tr. tr.* 0.24 tr.*
Steels 23 0.009 0.56* 0.16 0.37 18.74* tr. 0.43 tr.* 0.41 24 0.009
0.96 1.07* 0.14 13.58* tr. 0.10 tr.* 0.51 25 0.023* 0.56* 0.80 0.10
11.71* tr. 0.24 0.11 0.59* 26 0.010 0.68* 0.17 0.09 8.73 tr. 0.10
0.10 0.31 27 0.007 0.92 0.21 0.10 10.42 0.06 0.61* 0.14 0.26 28
0.008 0.90 0.15 0.10 10.39 0.01 0.10 0.35* 0.24 29 0.006 1.24* 0.16
0.01 8.39 0.02 0.11 0.17 0.29 30 0.009 0.88 0.23 0.11 10.23 tr.
0.13 0.13 0.31 31 0.010 0.96 0.25 0.09 9.96 tr. 0.09 0.16 0.30
(Mass %) 3Cr + Cr + AM No. V Al B N Other 40Si 10Si value
Comparative 22 tr. 0.08 tr.* 0.008 -- 53.2* 16.0 53.2 Steels 23
0.07 0.01 tr.* 0.014 -- 78.6 24.3* -21.6 24 0.04 0.02 tr.* 0.009 --
79.1 23.2* 26.6 25 0.03 0.01 0.0009 0.004 -- 57.5* 17.3 48.1 26
0.01 tr. 0.0022 0.009 -- 53.4* 15.5 80.5* 27 0.03 0.01 0.0014 0.008
-- 68.1 19.6 59.2 28 0.02 tr. 0.0017 0.008 -- 67.2 19.4 46.8 29
0.02 0.01 0.0021 0.008 -- 74.8 20.8 70.0 30 0.04 0.21* tr.* 0.007
La: 0.01 65.9 19.0 47.6 Ce: 0.04 31 0.03 0.01 0.0152* 0.007 Mg:
0.001 68.3 19.6 58.9 tr.: Trace *Outside invention range
[0077] A test piece cut from each cold-rolled annealed sheet was
subjected to a tensile test, a Charpy impact test, a
high-temperature tensile test, and a high-temperature oxidation
test.
[0078] Formability was evaluated based on the 0.2% strength,
breaking extension and plastic strain ratio determined by the
tensile test. No. 13B test pieces (prescribed by JIS Z 2201) cut
from each steel sheet specimen in directions parallel, 45 degrees
and 90 degrees to the rolling direction were used as the tensile
test pieces. 0.2% strength and breaking extension were determined
by subjecting the test piece taken 45 degrees with respect to the
rolling direction to the tests prescribed by JIS Z 2241. Plastic
strain ratio was determined in accordance with JIS Z 2254 using the
test pieces taken in all three of the aforesaid directions. More
specifically, the plastic strain ratio in each direction was
calculated from the ratio between the lateral strain and thickness
direction strain under application of a 15% uniaxial tensile
pre-strain, and the average plastic strain ratio r.sub.AV and
in-plane plastic anisotropy .DELTA.r were determined in accordance
with the following equations: r.sub.AV=(r.sub.L+2r.sub.D+r.sub.T)/4
.DELTA.r=(r.sub.L-2r.sub.D+r.sub.T)/2 where
[0079] r.sub.L=Plastic strain ratio parallel to rolling
direction
[0080] r.sub.D=Plastic strain ratio 45 degrees to rolling
direction
[0081] r.sub.T=Plastic strain ratio 90 degrees to rolling
direction.
[0082] The Charpy impact test was conducted by the method explained
with reference to FIG. 2. The energy transition temperature was
determined and used as an index of low-temperature toughness.
[0083] The high-temperature tensile test was conducted in
accordance with JIS G 0657 using the tensile test piece taken at 45
degrees. The 0.2% strength at 900.degree. C. was determined and
used as an index of high-temperature strength.
[0084] The high-temperature oxidation test was conducted in
accordance with JIS Z 2281 by determining the amount of oxidation
increase after heating at 900.degree. C. for 200 hours in the
atmosphere. The result was used as an index of high-temperature
oxidation resistance.
[0085] The results of the foregoing tests are shown in Table 3.
TABLE-US-00003 TABLE 3 Elongation at Room-temp. 0.2% room tem. in
Ave. plastic In-plane plastic Energy High-temp. 0.2% Weight gain
after proof stress tensile test strain ratio anisotropy transition
proof stress.sup.1) Oxidation test.sup.2) No. (Mpa) (%)
r.sub..LAMBDA..nu. .DELTA.r temp. (.degree. C.) (Mpa) (mg/cm.sup.2)
Invention 1 275 37 1.7 0.3 -50 14 1.2 Examples 2 280 36 1.6 0.4 -75
15 1.2 3 295 33 1.3 0.2 -50 13 1.1 4 285 35 1.4 0.2 -50 13 1.2 5
280 35 1.4 0.3 -50 13 1.2 6 280 36 1.4 0.3 -75 15 1.3 7 285 35 1.3
0.3 -50 14 1.1 8 285 35 1.4 0.2 -50 14 1.0 9 295 34 1.3 0.2 -50 14
1.1 10 295 33 1.3 0.3 -50 13 0.9 11 290 34 1.4 0.3 -50 16 1.0 12
290 34 1.3 0.4 -50 12 1.0 13 295 33 1.3 0.4 -50 13 1.1 14 295 33
1.3 0.3 -50 14 0.9 15 300 33 1.3 0.4 -50 14 0.7 16 300 33 1.3 0.3
-50 13 0.8 17 295 33 1.3 0.3 -50 13 0.6 18 280 35 1.5 0.3 -50 13
0.6 19 280 35 1.4 0.2 -50 13 0.4 20 280 36 1.6 0.5 -50 12 0.3 21
280 35 1.3 0.5 -50 11 1.4 Comparative 22 265 33 1.1* 0.7* -75 9*
18.0* Examples 23 335* 32* 1.0* 0.7* -50 12 0.8 24 355* 31* 1.1*
0.7* -25* 13 1.0 25 300 35 0.9* 0.5 -25* 15 5.0* 26 410* 29* 1.3
0.0 -50 13 9.0* 27 295 33 1.3 0.4 -25* 12 11.0* 28 290 33 1.3 0.5
-25* 12 1.3 29 300 33 1.3 0.4 -25* 12 1.2 30 290 33 1.3 0.5 -25* 13
0.5 31 295 33 1.3 0.5 -25* 13 0.7 *Object of invention not achieved
.sup.1)900.degree. C. .sup.2)900.degree. C. .times. 200 hr
[0086] As can be seen from Table 3, the steels Nos. 1-21, examples
of the present invention, all had softness (0.2% proof stress)
falling about midway between SUH409L (No. 22) and SUS430J1L (No.
23) and ductility (elongation) similar to SUH409L. They were
superior to SUH409L and SUS430J1L in deep drawability, i.e., in
average plastic strain ratio r.sub.AV and in-plane plastic
anisotropy .DELTA.r. Their low-temperature toughness (energy
transition temperature) performance was also excellent, matching
that of SUH409L. The invention steels were clearly superior to
SUH409L and substantially matched the performance of SUS430J1L in
900.degree. C. heat resistance (high-temperature strength and
high-temperature oxidation resistance). In short, the steels of the
present invention achieved excellent formability while also
thoroughly maintaining high-temperature strength, high-temperature
oxidation resistance and low-temperature toughness.
[0087] In contrast, steel No. 22, a comparative example steel
equivalent to SUH409L, was inferior in heat resistance, and No. 23,
equivalent to SUS430J1L, was hard and insufficient in formability.
Steels Nos. 24 and 25 are types that have actually been used in
automobile engine exhaust gas passage components. However, No. 24
was inferior in formability and low-temperature toughness owing to
the fact that, inter alia, it was not added with Ti and had Si and
Cr contents falling outside the ranges of the present invention,
while No. 25 was poor in formability, low-temperature toughness and
high-temperature oxidation resistance because it was high in C and
Nb and had Si and Cr contents falling outside the ranges of the
present invention. Steel No. 26 was inferior in formability and
high-temperature oxidation resistance because its phase stability
was on the austenitic side. Steels No. 27-31 exhibited deficient
low-temperature toughness because they contained elements harmful
to low-temperature toughness in amounts exceeding the ranges
specified by the present invention.
EXAMPLE 2
[0088] The steels shown in Tables 1 and 2 as from No. 1 to No. 10
and from No. 22 to 26 were hot-rolled and then subjected to the
heat treatment at temperatures between 950 and 1100.degree. C. for
1 minute, thereby to obtain the hot rolled sheet having totally
recrystallized structure. The sheets obtained were cold rolled into
2.0 mm and thereafter totally recrystallized by annealed at
1050.degree. C. for 1 minute to afford cold rolled annealed
sheets.
[0089] As same as Example 1, a piece cut from each cold-rolled
annealed sheet was subjected to the test to evaluate 0.2% strength,
breaking extension, plastic strain ratio and in-plane anisotropy.
Further, in order to evaluate a surface appearance after formed, a
piece cut from each cold-rolled annealed sheet was imposed 20% of
plastic strain in direction parallel to the rolling direction and
subjected to the test to evaluate the surface roughness of the test
piece surface in direction perpendicular to the rolling direction
by using a contact-type surface roughness meter, the surface
roughness being 10 points average roughness Rz in accordance with
JIS B 0660. For comparative, the surface roughness was tested as
same above for the test pieces shown in Table 3 which were derived
from the partially recrystallized hot-rolled sheets and the results
was shown in Table 4 as Comparative value of surface roughness.
TABLE-US-00004 TABLE 4 Elongation at Comparative Room-temp. 0.2%
room-temp. in Ave. plastic In-plane plastic Surface vale of surface
proof stress tensile test strain ratio anisotropy loughness
roughness.sup.1) No. (Mpa) (%) r.sub..LAMBDA..nu. .DELTA.r (.mu.m)
(.mu.m) Invention 1 270 37 1.6 0.5 15 35 Examples 2 275 37 1.5 0.6
20 40 3 290 34 1.3 0.4 15 35 4 285 35 1.4 0.4 10 30 5 280 36 1.3
0.5 10 30 6 275 36 1.3 0.5 15 40 7 280 36 1.3 0.5 25 45 8 285 35
1.4 0.6 20 40 9 295 35 1.2 0.4 15 35 10 290 34 1.2 0.5 15 35
Comparative 22 260 34 1.0* 0.8* 30 50 Examples 23 330 33 1.0* 0.9*
30 45 24 350* 32* 1.0* 0.8* 25 40 25 290 35 0.9* 0.8* 25 45 26 420*
28* 1.2 0.4 25 40 *Object of invention not achieved .sup.1)Test
results for the pieces of partially recrystallized hot-rolled
sheets shown in Table 3.
[0090] In comparison with the invention examples between Table 4
and Table 3, it is seen that the test pieces derived from totally
recrystallized hot-rolled sheets, those of Table 4, tend to have
same or lower average plastic strain ratio and have slightly
increased in-plane anisotropy than those of Table 3 for partially
recrystallized hot rolled sheets. This seems to be based on the
slight lowering of r value in direction 45 degree to the rolling
direction in the case of totally recrystallized hot-rolled sheets.
On the other hand, it is apparent that the surface roughness after
formed is remarkably decreased when the totally recrystallized
hot-rolled sheets are used. This means that by adoption of a
treatment for totally recrystallization in the hot-rolled sheet,
there is provided a steel sheet suitable for use in the article
formed which requires superior fine surface appearance. Comparative
examples in Table 4 are basically below in formability compared
with invention examples.
[0091] Thus the present invention enables concurrent improvement of
formability, high-temperature strength, high-temperature oxidation
resistance and low-temperature toughness in a ferritic
heat-resistant steel sheet. The ferritic steel sheet of the present
invention is particularly notable in that it offers excellent
formability, specifically deep drawability and isotropy thereof,
capable of responding to the needs of a diversity of forming
methods. In this aspect, the steel sheet of the present invention
is endowed with new capabilities not envisioned by conventional
ferritic heat-resistant steel sheets. It also offers
high-temperature strength, high-temperature oxidation resistance
and low-temperature toughness that achieve a performance level
equal to or better than the steel sheets currently used in exhaust
gas passage components. While conventional ferritic steel sheets
have been incapable of concurrently achieving high levels of
formability, high-temperature strength, high-temperature oxidation
resistance and low-temperature toughness, the present invention
concurrently achieves excellent performance on all of these points
at a Cr content of not more than 11%. As such, the present
invention enables application of ferritic heat-resistant steel to
complicatedly shaped exhaust gas passage components, helps to
expand the degree of freedom in designing such components, and
makes a marked contribution to cost reduction.
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