U.S. patent application number 10/385257 was filed with the patent office on 2003-09-25 for high-strength steel pipe of api x65 grade or higher and manufacturing method therefor.
This patent application is currently assigned to JFE STEEL CORPORATION. Invention is credited to Endo, Shigeru, Ishikawa, Nobuyuki, Shinmiya, Toyohisa, Suwa, Minoru.
Application Number | 20030180174 10/385257 |
Document ID | / |
Family ID | 26618660 |
Filed Date | 2003-09-25 |
United States Patent
Application |
20030180174 |
Kind Code |
A1 |
Ishikawa, Nobuyuki ; et
al. |
September 25, 2003 |
High-strength steel pipe of API X65 grade or higher and
manufacturing method therefor
Abstract
The present invention provides a high-strength steel pipe of API
X65 grade or higher consisting essentially of, by mass %, 0.02 to
0.08% of C, 0.01 to 0.5% of Si, 0.5 to 1.8% of Mn, 0.01% or less of
P, 0.002% or less of S, 0.01 to 0.07% of Al, 0.005 to 0.04% of Ti,
0.05 to 0.50% Mo, at least one element selected from 0.005 to 0.05%
of Nb and 0.005 to 0.10% of V, and the balance being Fe, in which
the volume percentage of ferritic phase is 90% or higher, and
complex carbides containing Ti, Mo, and at least one element
selected from Nb and V are precipitated in the ferritic phase. The
high-strength steel pipe in accordance with the present invention
has excellent HIC resistance and good toughness of heat-affected
zone, and can be manufactured stably at a low cost.
Inventors: |
Ishikawa, Nobuyuki;
(Fukuyama, JP) ; Shinmiya, Toyohisa; (Fukuyama,
JP) ; Endo, Shigeru; (Fukuyama, JP) ; Suwa,
Minoru; (Fukuyama, JP) |
Correspondence
Address: |
FRISHAUF, HOLTZ, GOODMAN & CHICK, P.C.
767 THIRD AVENUE
NEW YORK
NY
10017-2023
US
|
Assignee: |
JFE STEEL CORPORATION
|
Family ID: |
26618660 |
Appl. No.: |
10/385257 |
Filed: |
March 10, 2003 |
Related U.S. Patent Documents
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Application
Number |
Filing Date |
Patent Number |
|
|
10385257 |
Mar 10, 2003 |
|
|
|
PCT/JP02/07102 |
Jul 12, 2002 |
|
|
|
Current U.S.
Class: |
420/124 ;
148/593 |
Current CPC
Class: |
C21D 8/02 20130101; C22C
38/04 20130101; C22C 38/12 20130101; C22C 38/14 20130101; C21D 9/08
20130101; C22C 38/02 20130101; C22C 38/002 20130101; Y10S 148/909
20130101; C21D 8/105 20130101; C21D 8/10 20130101; C22C 38/06
20130101 |
Class at
Publication: |
420/124 ;
148/593 |
International
Class: |
C22C 038/14 |
Foreign Application Data
Date |
Code |
Application Number |
Jul 13, 2001 |
JP |
2001-213145 |
Nov 29, 2001 |
JP |
2001-364103 |
Claims
What is claimed is:
1. A high-strength steel pipe of API X65 grade or higher consisting
essentially of, by mass %, 0.02 to 0.08% of C, 0.01 to 0.5% of Si,
0.5 to 1.8% of Mn, 0.01% or less of P, 0.002% or less of S, 0.01 to
0.07% of Al, 0.005 to 0.04% of Ti, 0.05 to 0.50% Mo, at least one
element selected from 0.005 to 0.05% of Nb and 0.005 to 0.10% of V,
and the balance being Fe, in which the volume percentage of
ferritic phase is 90% or higher, and complex carbides containing
Ti, Mo, and at least one element selected from Nb and V are
precipitated in said ferritic phase.
2. The high-strength steel pipe of API X65 grade or higher
according to claim 1, wherein the content of Ti is 0.005 to less
than 0.02%.
3. A high-strength steel pipe of API X65 grade or higher consisting
essentially of, by mass %, 0.02 to 0.08% of C, 0.01 to 0.5% of Si,
0.5 to 1.8% of Mn, 0.01% or less of P, 0.002% or less of S, 0.01 to
0.07% of Al, 0.005 to 0.04% of Ti, at least one element selected
from 0.005 to 0.05% of Nb and 0.005 to 0.10% of V, W and Mo meeting
the condition that the content of (W/2+Mo) is in the range of 0.05
to 0.50% (however, a case where the content of Mo is 0% is
included), and the balance being Fe, in which the volume percentage
of ferritic phase is 90% or higher, and complex carbides containing
Ti, W, Mo, and at least one element selected from Nb and V are
precipitated in said ferritic phase.
4. The high-strength steel pipe of API X65 grade or higher
according to claim 3, wherein the content of Ti is 0.005 to less
than 0.02%.
5. The high-strength steel pipe of API X65 grade or higher
according to claim 1, wherein said steel pipe further contains
0.0005 to 0.0040% of Ca.
6. The high-strength steel pipe of API X65 grade or higher
according to claim 3, wherein said steel pipe further contains
0.0005 to 0.0040% of Ca.
7. The high-strength steel pipe of API X65 grade or higher
according to claim 1, wherein said steel pipe further contains at
least one element selected from 0.5% or less of Cu, 0.5% or less of
Ni, and 0.5% or less of Cr, by mass %.
8. The high-strength steel pipe of API X65 grade or higher
according to claim 3, wherein said steel pipe further contains at
least one element selected from 0.5% or less of Cu, 0.5% or less of
Ni, and 0.5% or less of Cr, by mass %.
9. The high-strength steel pipe of API X65 grade or higher
according to claim 1, wherein the ratio of the C content to the
total content of Mo, Ti, Nb, V and W,
R=(C/12)/[(Mo/96)+(Ti/48)+(Nb/93)+(V/51)+(W/184)], expressed by
mass %, is in the range of 0.5 to 3.0.
10. The high-strength steel pipe of API X65 grade or higher
according to claim 3, wherein the ratio R is in the range of 0.5 to
3.0.
11. The high-strength steel pipe of API X65 grade or higher
according to claim 9, wherein the ratio R is in the range of 0.7 to
2.0.
12. The high-strength steel pipe of API X65 grade or higher
according to claim 10, wherein the ratio R is in the range of 0.7
to 2.0.
13. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 1 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate at a
cooling rate not lower than 2.degree. C./s; coiling said cooled
steel plate at a temperature in the range of 550 to 700.degree. C.;
and forming said coiled steel plate into a steel pipe.
14. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 3 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate at a
cooling rate not lower than 2.degree. C./s; coiling said cooled
steel plate at a temperature in the range of 550 to 700.degree. C.;
and forming said coiled steel plate into a steel pipe.
15. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 1 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate to a
temperature in the range of 600 to 700.degree. C. at a cooling rate
not lower than 2.degree. C./s; cooling said cooled steel plate to
at least 550.degree. C. at a cooling rate not higher than
0.1.degree. C./s; and forming said steel plate into a steel
pipe.
16. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 3 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate to a
temperature in the range of 600 to 700.degree. C. at a cooling rate
not lower than 2.degree. C./s; cooling said cooled steel plate to
at least 550.degree. C. at a cooling rate not higher than
0.1.degree. C./s; and forming said steel plate into a steel
pipe.
17. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 1 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate to a
temperature in the range of 550 to 700.degree. C. at a cooling rate
not lower than 2.degree. C./s; heating said cooled steel plate
immediately after being cooled and keeping it at a temperature in
the range of 550 to 700.degree. C. for three minutes or longer; and
forming said steel plate into a steel pipe.
18. A manufacturing method for a high-strength steel pipe of API
X65 grade or higher, comprising the steps of: heating a steel slab
having chemical composition described in claim 3 to a temperature
in the range of 1000 to 1250.degree. C.; hot rolling said steel
slab at a finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling said steel plate to a
temperature in the range of 550 to 700.degree. C. at a cooling rate
not lower than 2.degree. C./s; heating said cooled steel plate
immediately after being cooled and keeping it at a temperature in
the range of 550 to 700.degree. C. for three minutes or longer; and
forming said steel plate into a steel pipe.
19. The manufacturing method for a high-strength steel pipe of API
X65 grade or higher according to claim 17, where in the heat
treatment for keeping said steel plate at a temperature in the
range of 550 to 700.degree. C. for three minutes or longer is
accomplished by using two or more induction heating apparatuses
provided in series on the same line as rolling equipment and
cooling equipment.
20. The manufacturing method for a high-strength steel pipe of API
X65 grade or higher according to claim 18, where in the heat
treatment for keeping said steel plate at a temperature in the
range of 550 to 700.degree. C. for three minutes or longer is
accomplished by using two or more induction heating apparatuses
provided in series on the same line as rolling equipment and
cooling equipment.
Description
[0001] This application is a continuation application of
International Application PCT/JP02/07102 (not published in English)
filed Jul. 12, 2002.
BACKGROUND OF THE INVENTION
[0002] 1. Field of the Invention
[0003] The present invention relates to a high-strength steel pipe
having a strength of API X65 grade or higher which is used for line
pipes, more particularly, a high-strength steel pipe having
excellent hydrogen-induced cracking resistance (HIC resistance),
and a manufacturing method thereof.
[0004] 2. Description of Related Arts
[0005] A steel pipe for line pipes, which is used for
transportation of crude oil or natural gas containing hydrogen
sulfide, is required to have what we call sour resistance including
HIC resistance and stress corrosion cracking resistance (SCC
resistance) as well as high strength, excellent toughness, and good
weldability. It is said that HIC is caused by an internal pressure
that is produced by a phenomenon that hydrogen ions created by
corrosion reaction are adsorbed on the steel surface, intrude into
steel as atomic hydrogen, and accumulate around nonmetallic
inclusions such as MnS and hard second phases such as martensite in
steel.
[0006] To prevent HIC, Unexamined Japanese Patent Publication No.
54-110119 has disclosed a manufacturing method of line-pipe steels,
in which by adding Ca or Ce in proper amounts relative to the
amount of S, and forming fine spherical inclusions to decrease
stress concentration instead of formation of needle-like MnS
inclusions. Unexamined Japanese Patent Publication No. 61-60866 and
Unexamined Japanese Patent Publication No. 61-165207 have disclosed
a steel in which the formation of island-like martensite that
functions as an origin of cracking in a center segregation region
and hard phases such as martensite or bainite that function as a
propagation path of cracking is restrained by a decrease in amount
of segregation-prone elements (C, Mn, P, etc.), soaking treatment
at a stage of slab heating, accelerated cooling during
transformation at a stage of cooling, etc. Unexamined Japanese
Patent Publication No. 5-9575, Unexamined Japanese Patent
Publication No. 5-271766, and Unexamined Japanese Patent
Publication No. 7-173536 have disclosed a steel plate having a
strength of API X80 grade or higher, in which the shape of
inclusions is controlled by adding Ca to a low-S steel, center
segregation is restrained by lower contents of C and Mn, and high
strength is provided by the addition of Cr, Mn and Ni and
accelerated cooling. All of these methods for preventing HIC are
methods for preventing HIC caused by center segregation.
[0007] However, a steel plate having a strength of API X65 grade or
higher is usually manufactured by accelerated cooling or direct
quenching, so that a near surface region of the steel plate which
receives high cooling rate is more liable to be hardened than the
interior thereof, and hence HIC occurs easily in the near surface
region. Also, microstructure obtained by accelerated cooling
consists of bainite and acicular ferrite having relatively high HIC
sensitivity not only in the near surface region but also in the
interior, so that the above-described method for preventing HIC
caused by center segregation does not suffice. Therefore, in order
to prevent HIC of steel plate completely, measures must be taken
against HIC caused by the microstructure of the near surface region
of steel plate and HIC caused by inclusions such as sulfide or
oxide as well as HIC caused by center segregation.
[0008] On the other hand, Unexamined Japanese Patent Publication
No. 7-216500 has disclosed an API X80 grade HIC-resistant steel
that is composed of ferrite and bainite phases and does not contain
block-like bainite or martensite phases with high HIC-sensitivity.
Unexamined Japanese Patent Publication No. 61-227129 and Unexamined
Japanese Patent Publication No. 7-70697 have disclosed
high-strength steels in which SCC resistance and HIC resistance are
improved by ferritic microstructure and Mo or Ti is added to
utilize precipitation strengthening by carbides.
[0009] However, the microstructure of the high-strength steel
described in Unexamined Japanese Patent Publication No. 7-216500
consists of bainite phases with relatively high HIC sensitivity.
Also, this steel is high in manufacturing cost because the content
of S and Mn is restricted severely and Ca treatment is necessary.
The microstructure of the high-strength steels described in
Unexamined Japanese Patent Publication No. 61-227129 and Unexamined
Japanese Patent Publication No. 7-70697 consists of ductile
ferritic phases, so that the HIC sensitivity is very low, while the
strength is low. In order to obtain higher strength for the steel
described in Unexamined Japanese Patent Publication No. 61-227129,
large amounts of C and Mo are added, cold-rolling is performed
after quench-and-temper, and tempering is performed again to
precipitate a large amount of carbides, resulting in increased
manufacturing cost. The steel described in Unexamined Japanese
Patent Publication No. 7-70697 cannot achieve high strength stably
because Ti is added to obtain high strength by utilizing
precipitation strengthening of TiC at a stage of coiling, but TiC
is liable to be coarsened by the influence of coiling temperature.
Although high strength can be achieved stably by adding large
amount of Ti, the toughness of heat-affected zone deteriorates
significantly when the welding such as electric resistance welding
or submerged arc welding are applied.
SUMMARY OF THE INVENTION
[0010] An object of the present invention is to provide a
high-strength steel pipe of API X65 grade or higher which has
excellent HIC resistance and good toughness after welding, and
which can be manufactured stably at a low cost, and a manufacturing
method thereof.
[0011] The above object can be attained by a high-strength steel
pipe of API X65 grade or higher consisting essentially of, by mass
%, 0.02 to 0.08% of C, 0.01 to 0.5% of Si, 0.5 to 1.8% of Mn, 0.01%
or less of P, 0.002% or less of S, 0.01 to 0.07% of Al, 0.005 to
0.04% of Ti, 0.05 to 0.50% Mo, at least one element selected from
0.005 to 0.05% of Nb and 0.005 to 0.10% of V, and the balance being
Fe, in which the volume percentage of ferritic phase is 90% or
higher, and complex carbides containing Ti, Mo, and at least one
element selected from Nb and V are precipitated in the ferritic
phase.
[0012] This high-strength steel pipe is manufactured, for example,
by a manufacturing method for a high-strength steel pipe of API X65
grade or higher, comprising the steps of heating a steel slab
having chemical composition described above to a temperature in the
range of 1000 to 1250.degree. C.; hot rolling the steel slab at a
finish temperature not lower than the Ar3 transformation
temperature to make a steel plate; cooling the steel plate at a
cooling rate not lower than 2.degree. C./s; coiling the cooled
steel plate at a temperature in the range of 550 to 700.degree. C.;
and forming the coiled steel plate into a steel pipe.
BRIEF DESCRIPTION OF THE DRAWINGS
[0013] FIG. 1 is a diagram showing the relationship between Ti
content and Charpy fracture appearance transition temperature of
heat-affected zone;
[0014] FIG. 2 is a view showing one example of microstructure of a
high-strength steel in accordance with the present invention;
[0015] FIG. 3 is a diagram showing an EDX analysis result of
precipitates;
[0016] FIG. 4 is a view showing one example of a production line
for a steel plate; and
[0017] FIG. 5 is a graph showing one example of heat treatment
using an induction heating apparatus.
DETAILED DESCRIPTION OF THE INVENTION
[0018] The inventors obtained the following findings as a result of
study on HIC resistance and toughness of welded part of a
high-strength steel pipe having a strength of API X65 grade or
higher which is used for line pipes.
[0019] 1) If hard second phases such as bainite, martensite,
pearlite, etc. exist in a ferritic phase, accumulation of hydrogen
and stress concentration are prone to occur at the phase interface,
so that a volume percentage of ferritic phase not lower than 90% is
effective in improving HIC resistance.
[0020] 2) It is well known that Mo and Ti are elements forming
carbides in steel, and the steel is strengthened by precipitation
of MoC or TiC. Carbides precipitated in a ferritic phase by
co-addition of Mo and Ti are represented by (Mo, Ti)C, and these
carbides are complex carbides in which (Mo, Ti) and C are bonded to
each other at an atom ratio of about 1:1. The carbides are very
fine, smaller than 10 nm, because they are stable and have a low
growth rate. Therefore, these complex carbides have a more powerful
strengthening function than the conventional MoC and TiC. Such very
fine carbides exert no influence on HIC.
[0021] 3) In the steels containing Ti, as the Ti content increases,
the toughness of heat-affected zone deteriorates. To prevent this
deterioration, it is effective to add at least one element selected
from Nb and V in addition to Mo and Ti and to precipitate fine
complex carbides containing Mo, Ti, Nb and/or V.
[0022] 4) By the above-described microstructure, both a high
strength of API X65 grade or higher and HIC resistance such that
cracking does not occur in a HIC test in accordance with NACE
Standard TM-02-84 can be achieved. In particular, both a high
strength of API X70 grade or higher and excellent HIC resistance
can be achieved for the first time by the present invention.
[0023] The present invention has been made based on the above
findings. The reason for limiting the content of each element will
be described below.
[0024] C: C is an element for strengthening steel by precipitation,
as carbides. However, if the C content is lower than 0.02%, a
strength of API X65 grade or higher cannot be obtained, and if it
exceeds 0.08%, the HIC resistance and the toughness of welded part
deteriorate. Therefore, the C content should be 0.02 to 0.08%.
[0025] Si: Si is an element necessary for deoxidization of steel.
However, if the Si content is lower than 0.01%, the deoxidization
effect is insufficient, and if it exceeds 0.5%, the weldability and
the toughness deteriorate. Therefore, the Si content should be 0.01
to 0.5%.
[0026] Mn: Mn is an element for strengthening steel and improving
the toughness. However, if the Mn content is lower than 0.5%, its
effect is insufficient, and if it exceeds 1.8%, the weldability and
the HIC resistance deteriorate. Therefore, the Mn content should be
0.5 to 1.8%.
[0027] P: P is an element that deteriorates the weldability and the
HIC resistance. Therefore, the P content should be not higher than
0.01%.
[0028] S: S turns to MnS inclusion in steel and hence deteriorates
the HIC resistance. Therefore, the S content should not be higher
than 0.002%.
[0029] Al: Al is added as a deoxidizer. If the Al content is lower
than 0.01%, the deoxidization effect is not achieved, and if it
exceeds 0.07%, the cleanliness of steel degrades and thus the HIC
resistance deteriorates. Therefore, the Al content should be 0.01
to 0.07%.
[0030] Ti: Ti is an important element in the present invention. If
the Ti content is not lower than 0.005%, Ti forms complex carbides
together with Mo as described above, so that strengthening of steel
is promoted. However, as shown in FIG. 1, if the Ti content exceeds
0.04%, the Charpy fracture appearance transition temperature of
heat-affected zone exceeds -20.degree., and hence the toughness
deteriorates. Therefore, the Ti content should be 0.005 to 0.04%.
Further, if the Ti content is lower than 0.02%, the Charpy fracture
appearance transition temperature of heat-affected zone is not
higher than -40.degree., and hence higher toughness is obtained.
Therefore, the Ti content should preferably be 0.005 to less than
0.02%.
[0031] Mo: As described above, Mo is an important element in the
present invention, like Ti. If the Mo content is not lower than
0.05%, pearlite transformation is restrained at a stage of cooling
after hot rolling, and fine complex carbides are formed together
with Ti so that the strengthening of steel is promoted. However, if
the Mo content exceeds 0.50%, hard phases such as bainite or
martensite are formed, and hence the HIC resistance deteriorates.
Therefore, the Mo content should be 0.05 to 0.50%.
[0032] Nb: Nb improves the toughness by microstructure refining,
and forms complex carbides together with Ti and Mo, contributing to
the strengthening of steel. However, if the Nb content is lower
than 0.005%, its effect is not achieved, and if it exceeds 0.05%,
the toughness of heat-affected zone deteriorates. Therefore, the Nb
content should be 0.005 to 0.05%.
[0033] V: V forms complex carbides together with Ti and Mo, like
Nb, contributing to the strengthening of steel. However, if the V
content is lower than 0.005%, its effect is not achieved, and if it
exceeds 0.1%, the toughness of welded part deteriorates. Therefore,
the Nb content should be 0.005 to 0.1%.
[0034] If at least one element selected from Nb and V is contained,
the strengthening and improvement in toughness of heat-affected
zone are achieved.
[0035] The balance other than the above-described components is Fe.
Also, other elements such as unavoidable impurities may be
contained as far as these elements exert no influence on the
operation and effects of the present invention.
[0036] If the ratio of the number of complex carbides smaller than
10 nm and containing Mo and Ti to the number of all the
precipitates excluding TiN, which contributes less to the
strengthening of steel, is not smaller than 80%, preferably not
smaller than 95%, the strengthening of steel can be promoted.
[0037] FIG. 2 shows one example of a microstructure of the steel in
accordance with the present invention, which is manufactured in a
hot rolling mill for steel sheet (coiling temperature: 650.degree.
C.) using a steel having composition of 0.05% C, 0.15% Si, 1.26%
Mn, 0.11% Mo, 0.018% Ti, 0.039% Nb, and 0.048% V. It can be
verified that many fine precipitates smaller than 10 nm in size are
dispersed. Also, FIG. 3 shows a result of analysis of precipitates
made by an energy dispersion X-ray spectroscopy method (EDX). It
can be seen that the precipitates are complex carbides containing
Ti, Nb, V and Mo.
[0038] Further, W is added in place of Mo or together with Mo so
that the content of (W/2+Mo) is in the range of 0.05 to 0.50%. In
this case as well, fine complex carbides are formed together with
Ti, and hence the strengthening of steel is promoted. If the
content of (W/2+Mo) exceeds 0.50%, hard phases such as bainite or
martensite are formed, deteriorating the HIC resistance.
[0039] Further, if Ca is added, the shape of sulfide inclusions is
controlled, and hence the HIC resistance is improved. However, if
the Ca content is lower than 0.0005%, its effect is insufficient,
and if it exceeds 0.0040%, the cleanliness of steel degrades and
thus the HIC resistance deteriorates. Therefore, the Ca content
should be 0.0005 to 0.0040%.
[0040] Still further, if at least one element selected from Cu, Ni
and Cr is contained in an amount described below, further
strengthening of steel can be achieved.
[0041] Cu: Cu is an effective element for improving the toughness
and increasing the strength. However, if the Cu content exceeds
0.5%, the weldability deteriorates. Therefore, the Cu content
should be not higher than 0.5%.
[0042] Ni: Ni is an effective element for improving the toughness
and increasing the strength. However, if the Ni content exceeds
0.5%, the HIC resistance deteriorates. Therefore, the Ni content
should be not higher than 0.5%.
[0043] Cr: Cr is an effective element for increasing the strength,
like Mn. However, if the Cr content exceeds 0.5%, the weldability
deteriorates. Therefore, the Cr content should be not higher than
0.5%.
[0044] If not only the content of each component but also Ceq
expressed by the following equation (1) is controlled, the
toughness of heat-affected zone is further improved. In particular,
it is preferable that Ceq be not higher than 0.30% for API X65
grade, Ceq be not higher than 0.32% for API X70 grade, and Ceq be
not higher than 0.34% for API X80 grade.
Ceq=C+Mn/6+(Cu+Ni)/15+(Cr+Mo+V)/5 (1)
[0045] Further, if R expressed by the following equation (2) is in
the range of 0.5 to 3.0, thermally stable and very fine complex
carbides can be obtained, so that strengthening of steel and
improvement in toughness of heat-affected zone can be achieved more
stably. In order to obtain a far higher strength, the R should
preferably be 0.7 to 2.0.
R=(C/12)/[(Mo/96)+(Ti/48)+(Nb/93)+(V/51)+(W/184)] (2)
[0046] Next, a manufacturing method for the high-strength steel
pipe in accordance with the present invention will be
described.
[0047] A steel slab having the above-described composition is
heated to a temperature in the range of 1000 to 1250.degree. C.,
and is hot rolled at a finish temperature not lower than the Ar3
transformation temperature. Then the rolled plate is cooled at a
cooling rate not lower than 2.degree. C./s and is coiled at a
temperature in the range of 550 to 700.degree. C., and finally, a
steel pipe is formed. Thereby, a high-strength steel pipe of API
X65 grade or higher which is composed of ferritic phase with a
volume percentage not lower than 90% and complex carbides
containing Ti, Mo, and at least one element selected from Nb and V
which are dispersed in the ferritic phase can be obtained.
[0048] If the heating temperature of slab is lower than
1000.degree. C., the carbides are not resolved sufficiently, so
that a necessary strength cannot be obtained, and if the heating
temperature exceeds 1250.degree. C., the toughness deteriorates.
Therefore, the heating temperature of slab should be 1000 to
1250.degree. C.
[0049] If hot rolling is performed at a finish temperature lower
than the Ar3 transformation temperature, the microstructure becomes
elongated in the rolling direction, and hence the HIC resistance
deteriorates. Therefore, hot rolling should be performed at a
finish temperature not lower than the Ar3 transformation
temperature. To prevent a decrease in toughness due to coarse
microstructure, hot rolling should preferably be performed at a
finish temperature not higher than 950.degree. C.
[0050] After hot rolling, if the rolled plate is cooled at a
cooling rate lower than 2.degree. C./s as in the case of air
cooling or slow cooling, complex carbides begin to precipitate at a
high temperature region and coarsen easily, which inhibits the
strengthening of steel. For this reason, the rolled plate must be
cooled at a cooling rate not lower than 2.degree. C./s. At this
time, if the cooling finish temperature is too high, the
precipitates are coarsened, so that a sufficient strength is not
obtained. Therefore, the cooling finish temperature should
preferably be not lower than the coiling temperature and not higher
than 750.degree. C.
[0051] After being cooled at a cooling rate not lower than
2.degree. C./s, the steel plate must be coiled at a temperature in
the range of 550 to 700.degree. C., preferably in the range of 600
to 660.degree. C., to obtain ferritic phase and fine complex
carbides. If the coiling temperature is lower than 550.degree. C.,
bainitic phase is formed, and hence the HIC resistance
deteriorates. If the coiling temperature exceeds 700.degree. C.,
the complex carbides coarsen, and hence a sufficient strength
cannot be obtained.
[0052] This coiling method for coiling the steel plate at a
temperature in the range of 550 to 700.degree. C. is used when a
steel plate which is a raw material for a steel pipe is
manufactured in a hot rolling mill for steel sheet. In this case,
the steel plate is formed into an electric resistance welded steel
pipe or a spiral steel pipe by the press bent forming method or the
roll forming method.
[0053] When a steel plate which is a raw material for a steel pipe
is manufactured in a hot rolling mill for heavy gauge steel plate,
instead of being coiled at a temperature in the range of 550 to
700.degree. C., it is necessary that the steel plate be cooled to a
temperature in the range of 600 to 700.degree. C. at a cooling rate
not lower than 2.degree. C./s, and then it be slowly cooled at
least to 550.degree. C. at a cooling rate not higher than
0.1.degree. C./s, or the steel plate be cooled to a temperature in
the range of 550 to 700.degree. C., and immediately after that, it
be subjected to heat treatment at temperatures in the range of 550
to 700.degree. C. for three minutes or longer. In this case, the
steel plate is formed into a UOE steel pipe by the UOE forming
method.
[0054] As means for slowly cooling the steel plate at a cooling
rate not higher than 0.1.degree. C./s, there can be used a method
in which steel plates are piled and cooled or a method in which the
steel plate is cooled in a box furnace etc.
[0055] If an induction heating apparatus is provided on a plate
manufacturing line, the heat treatment at temperatures in the range
of 550 to 700.degree. C. for three minutes or longer can be
accomplished without a decrease in the temperature of steel plate
to below 550.degree. C., which does not result in decreased
productivity.
[0056] FIG. 4 shows one example of an equipment layout on a plate
manufacturing line.
[0057] On the manufacturing line 1, a hot rolling mill 3, an
accelerated cooling apparatus 4, an induction heating apparatus 5
and a hot leveler 6 are arranged in order from the upstream side to
the downstream side. After a slab coming out of a heating furnace
is rolled into a steel plate 2 by the hot rolling mill 3, the steel
plate 2 is cooled by the accelerated cooling apparatus 4, and is
subjected to heat treatment by the induction heating apparatus 5.
Then, the steel plate 2 is corrected in shape by the hot leveler 6,
and is sent to a pipe manufacturing process.
[0058] FIG. 5 shows one example of heat treatment using the
induction heating apparatus.
[0059] In this example, the steel plate is kept at temperatures in
the range of 550 to 700.degree. C. by performing two cycles of
heating using the induction heating apparatus. The induction
heating apparatus is turned on and off so that the highest
temperature (Tmax) does not exceed 700.degree. C. and the lowest
temperature (Tmin) is not lower than 550.degree. C., by which the
steel plate is kept at temperatures in the range of 550 to
700.degree. C. for three minutes or longer in total. The induction
heating arises a difference in temperature between the surface
layer and the interior of steel plate. The temperature specified
herein is an average plate temperature when heat transfers from the
surface layer to the interior and becomes even.
EXAMPLE 1
[0060] Electric resistance welded steel pipes Nos. 1 to 29 with an
outside diameter of 508.0 mm and a wall thickness of 12.7 mm were
manufactured, using the steels A to O having chemical composition
given in Table 1 and hot rolled under conditions given in Table 2
in a hot rolling mill for steel sheet. Also, UOE steel pipes Nos.
30 to 35 with an outside diameter of 914.4 mm and a wall thickness
of 19.1 mm and with an outside diameter of 1219.2 mm and a wall
thickness of 25.4 mm were manufactured, using steel plates which
were produced under conditions given in Table 3 in a hot rolling
mill for steel plate. The steel plates were piled and slowly cooled
to room temperature from a certain temperature. The mean cooling
rate from the start of slow cooling to 550.degree. C. is
additionally shown in Table 3. Also, the UOE steel pipes given in
Table 3 were expanded by 1.2% after they were seam welded by
submerged arc welding.
[0061] The microstructure of steel pipe was observed using an
optical microscope and a transmission electron microscope (TEM).
The composition of precipitates was analyzed by an energy
dispersion X-ray spectroscopy method (EDX).
[0062] Also, a full-thickness tensile test piece in accordance with
API standard was cut out in the circumference direction to conduct
a tensile test, by which yield strength and tensile strength were
measured. Considering variations due to manufacturing conditions,
the steel pipe having a tensile strength not lower than 550 MPa was
regarded as meeting the standard of API X65 grade, the steel pipe
having a tensile strength not lower than 590 MPa was regarded as
meeting the standard of API X70 grade, and the steel pipe having a
tensile strength not lower than 680 MPa was regarded as meeting the
standard of API X80 grade.
[0063] Further, HIC resistance and toughness of heat-affected zone
(HAZ) were measured. For HIC resistance, a HIC test of dipping time
of 96 hours in accordance with NACE Standard TM-02-84 was
conducted, and the case where cracking was not recognized was
indicated by .largecircle., and the case where cracking occurred
was indicated by. For HAZ toughness, a 2-mm V notch Charpy test
piece was taken in the circumference direction in the electric
resistance welded portion or the seam welded portion to measure
fracture appearance transition temperature (vTrs). At this time,
the V notch was formed in the center of electric resistance welded
portion for steel pipes Nos. 1 to 29 and in the bond portion
(fusion line) at the position of t/2 (t is plate thickness) for
steel pipes Nos. 30 to 35.
[0064] The test results are given in Tables 2 and 3.
[0065] All of steel pipes Nos. 1 to 18 in accordance with the
present invention were of X65 grade or higher, and had excellent
HIC resistance and HAZ toughness. The microstructure of those steel
pipes was substantially a ferritic phase, in which fine carbides
with a particle diameter smaller than 10 nm which contained Ti, Mo,
and at least one element selected from Nb and V were dispersed.
Steel pipes Nos. 3, 4, 5, 10, 11, 12, 17 and 18 using B, C, F and I
steels in which the Ti content is lower than 0.005 to 0.02%
exhibited higher HAZ toughness. Also, steel pipes Nos. 1 to 15
using A to G steels in which the ratio of the C content to the
total content of Mo, Ti, Nb, V and W was in the range of 0.7 to 2.0
had a higher strength than steel pipes Nos. 16 to 18 using H and I
steels.
[0066] For steel pipes Nos. 19 to 23 as comparative examples, the
microstructure thereof was not substantially a ferritic phase
because the manufacturing method was outside the range of the
present invention, and fine carbides containing Ti, Mo, and at
least one element selected from Nb and V were not precipitated, so
that a sufficient strength was not obtained and cracking was
observed in the HIC test. For steel pipe No. 19, a sufficient
amount of solute carbon could not be secured because of low heating
temperature, and a sufficient strength could not be obtained
because of lack in carbides precipitated at the coiling time. For
steel pipe No. 20, since the rolling finish temperature was low,
the microstructure became elongated in the rolling direction, and
hence the HIC resistance deteriorated. For steel pipe No. 21, since
the cooling rate after rolling was low, carbides began to
precipitate from a high temperature region and were coarsened, so
that the strength was decreased. For steel pipe No. 22, since the
coiling temperature was high, carbides were coarsened, so that a
sufficient strength was not obtained. For steel pipe No. 23, since
the coiling temperature was low, the structure contained bainitic
phase, so that the HIC resistance deteriorated.
[0067] Also, steel pipes Nos. 24 to 29 as comparative examples had
problems of insufficient strength, occurrence of cracking in HIC
test, and deteriorated HAZ toughness because the chemical
composition was outside the range of the present invention. For
steel pipes Nos. 24 and 25, since the content of Mo or Ti was low,
sufficient precipitation strengthening was not achieved, so that
the strength was low. For steel pipe No. 26, since the Ti content
was too high, the microstructure was coarsened by welding heat, so
that the HAZ toughness deteriorated. For steel pipe No. 27, since
the C content was low, sufficient precipitation strengthening was
not achieved, so that the strength was low. For steel pipe No. 28,
since the C content was too high, bainitic phase was formed, and
hence the HIC resistance deteriorated. For steel pipe No. 29, since
the S content was too high, many sulfide inclusions were formed, so
that the HIC resistance deteriorated.
[0068] All of steel pipes Nos. 30 to 33 in accordance with the
present invention had a tensile strength of 580 MPa or higher, and
also had high HIC resistance and HAZ toughness. The structure of
steel pipe was substantially a ferritic phase, in which fine
carbides with a particle diameter smaller than 10 nm which
contained Ti, Mo, and at least one element selected from Nb and V
were dispersed.
[0069] For steel pipe No. 34 as comparative examples, since the
cooling rate was high at the time of slow cooling, and the
microstructure contained bainitic phase, the HIC resistance
deteriorated. Also, for steel pipe No. 35, since the chemical
composition was outside the range of the present invention and the
Ti content was high, the HAZ toughness deteriorated.
1TABLE 1 C/ (Mo + Ti + Nb + Steel V + Re- Type C Si Mn P S Mo Ti Nb
V Al Cu Ni Cr Ca W Ceq Zr)* marks A 0.045 0.18 1.15 0.008 0.0008
0.13 0.022 0.046 0.032 0.26 1.62 Ex- B 0.053 0.25 1.23 0.005 0.0004
0.21 0.016 0.069 0.041 0.31 1.14 ample C 0.051 0.15 1.26 0.008
0.0007 0.11 0.018 0.039 0.048 0.042 0.29 1.47 D 0.061 0.30 1.16
0.010 0.0009 0.32 0.034 0.028 0.025 0.32 1.17 E 0.042 0.26 1.09
0.009 0.0013 0.14 0.028 0.052 0.030 0.12 0.18 0.0015 0.18 0.28 0.86
F 0.047 0.14 1.20 0.002 0.0008 0.12 0.013 0.038 0.048 0.034 0.12
0.0021 0.30 1.36 G 0.050 0.21 1.25 0.007 0.0010 0.28 0.035 0.021
0.028 0.032 0.0023 0.32 0.94 H 0.032 0.25 1.06 0.008 0.0012 0.24
0.022 0.030 0.035 0.022 0.0015 0.26 0.67 I 0.060 0.30 1.22 0.005
0.0009 0.06 0.008 0.028 0.035 0.025 0.0032 0.28 2.80 J 0.055 0.23
1.24 0.006 0.0010 0.02 0.019 0.032 0.023 0.033 0.27 3.27 Com- K
0.047 0.32 1.33 0.008 0.0013 0.14 0.002 0.048 0.041 0.026 0.30 1.39
para- L 0.049 0.22 1.58 0.011 0.0015 0.16 0.051 0.009 0.045 0.027
0.35 1.10 tive M 0.012 0.17 1.16 0.005 0.0008 0.23 0.008 0.034
0.034 0.24 0.31 0.29 0.31 ex- N 0.093 0.29 1.15 0.003 0.0009 0.12
0.008 0.030 0.052 0.039 0.0024 0.32 2.80 ample O 0.050 0.28 1.36
0.004 0.0023 0.21 0.013 0.028 0.048 0.025 0.16 0.22 0.0018 0.35
1.12 Unit: mass %, *; at % Underline indicates outside the range of
this invention
[0070]
2TABLE 2 Cooling Rolling rate Treatment Toughness Outside Wall
Heating finish (after method Coiling Yeild Tensile of welded Steel
diameter thickness temp. temp. rolling) after temp. strength
strength HIC vTrs part re- No. type (mm) (mm) (.degree. C.)
(.degree. C.) (.degree. C./s) cooling (.degree. C.) (MPa) (MPa)
Grade resistance (.degree. C.) marks 1 A 508.0 12.7 1150 910 20
Coiling 650 505 587 X65 .largecircle. -50 2 A 508.0 12.7 1150 870
20 Coiling 620 502 592 X65 .largecircle. -45 3 B 508.0 12.7 1200
900 20 Coiling 635 552 642 X70 .largecircle. -57 4 C 508.0 12.7
1200 900 20 Coiling 650 558 658 X70 .largecircle. -64 5 C 508.0
12.7 1200 900 20 Coiling 615 542 635 X70 .largecircle. -60 6 D
508.0 12.7 1200 880 20 Coiling 590 588 706 X80 .largecircle. -47 7
D 508.0 12.7 1150 880 20 Coiling 620 608 726 X80 .largecircle. -40
8 D 508.0 12.7 1150 880 20 Coiling 680 608 713 X80 .largecircle.
-43 9 E 508.0 12.7 1150 880 20 Coiling 635 548 638 X70
.largecircle. -48 10 F 508.0 12.7 1200 900 20 Coiling 665 578 669
X70 .largecircle. -65 11 F 508.0 12.7 1200 900 8 Coiling 650 548
652 X70 .largecircle. -62 12 F 508.0 12.7 1150 900 20 Coiling 650
570 665 X70 .largecircle. -72 13 G 508.0 12.7 1150 880 20 Coiling
635 630 738 X80 .largecircle. -31 14 G 508.0 12.7 1150 880 20
Coiling 600 614 726 X80 .largecircle. -37 15 G 508.0 12.7 1050 870
8 Coiling 600 591 682 X80 .largecircle. -35 16 H 508.0 12.7 1150
870 20 Coiling 630 481 556 X65 .largecircle. -27 17 I 508.0 12.7
1050 870 20 Coiling 640 458 551 X65 .largecircle. -60 18 I 508.0
12.7 1200 800 20 Coiling 620 473 553 X65 .largecircle. -57 19 A
508.0 12.7 950 740 20 Coiling 585 415 512 X52 X -53 20 A 508.0 12.7
1200 730 20 Coiling 625 483 564 X65 X -57 21 A 508.0 12.7 1200 910
1 Coiling 630 446 520 X52 .largecircle. -50 22 A 508.0 12.7 1200
910 20 Coiling 725 392 458 X52 .largecircle. -45 23 B 508.0 12.7
1150 880 20 Coiling 520 538 633 X70 X -60 24 J 508.0 12.7 1150 900
20 Coiling 650 419 483 X52 .largecircle. -60 25 K 508.0 12.7 1150
900 20 Coiling 635 416 483 X52 .largecircle. -33 26 L 508.0 12.7
1150 900 20 Coiling 640 630 736 X80 .largecircle. 3 27 M 508.0 12.7
1150 900 20 Coiling 640 395 492 X52 .largecircle. -65 28 N 508.0
12.7 1200 900 20 Coiling 635 540 625 X70 X -27 29 0 508.0 12.7 1200
900 20 Coiling 635 537 658 X70 X -40 Underline indicates outside
the range of this invention
[0071]
3TABLE 3 Cooling Cooling Cool- rate Outside Wall Rolling Rate ing
Treatment (at the Toughness di- thick- Heating finish (after stop
method slow Yeild Tensile of welded Steel ameter ness temp. temp.
rolling) temp. after cooling) strength strength HIC vTrs part Re-
No. type (mm) (mm) (.degree. C.) (.degree. C.) (.degree. C.)
(.degree. C.) cooling (.degree. C/s). (MPa) (MPa) grade resistance
(.degree. C.) marks 30 A 1219.2 25.4 1150 900 22 640 Slow 0.04 485
596 X65 .largecircle. -45 Ex- cooling ample 31 F 914.4 19.1 1150
880 30 660 Slow 0.08 520 650 X70 .largecircle. -48 cooling 32 F
1219.2 25.4 1200 900 22 655 Slow 0.04 542 663 X70 .largecircle. -55
cooling 33 G 1219.2 25.4 1200 900 22 670 Slow 0.04 585 710 X80
.largecircle. -42 cooling 34 D 1219.2 25.4 1200 900 22 635 Slow 1
448 537 X65 X -58 Com- cooling para- 35 L 914.4 19.1 1200 900 30
650 Slow 0.08 574 706 X80 .largecircle. 2 tive cooling ex- ample
Underline indicates outside the range of this invention
EXAMPLE 2
[0072] Steel plates were manufactured under the conditions given in
Table 5 in a hot rolling mill for a steel plate by making slabs
from steels a to i having chemical composition given in Table 4 by
the continuous casting method. After being hot rolled, the rolled
steel plates were immediately cooled by using a water-cooled inline
accelerated cooling apparatus, and were subjected to heat treatment
by using three inline induction heating apparatuses provided in
series on the manufacturing line or a gas-fired furnace. In Table
5, each temperature is an average plate temperature, and the
maximum and minimum temperatures are the above-described highest
and lowest temperatures at the time of heat treatment. Also, the
number of cycles means the number of cycles of heating performed by
using the induction heating apparatuses to keep the steel plate at
temperatures in the range of 550 to 700.degree. C. for three
minutes or longer. In the case of gas firing, the steel plate was
kept at a fixed temperature.
[0073] As in example 1, UOE steel pipes Nos. 36 to 51 with an
outside diameter of 914.4 mm and a wall thickness of 19.1 mm and
with an outside diameter of 1219.2 mm and a wall thickness of 25.4
mm were manufactured, and the microstructure, yield strength,
tensile strength, HIC resistance, and HAZ toughness were
measured.
[0074] The measurement results are given in Table 5.
[0075] All of steel pipes Nos. 36 to 43, which were examples of the
present invention, had a tensile strength not lower than 600 MPa,
and also had high HIC resistance and HAZ toughness. The
microstructure of steel pipe was substantially a ferrite phase, in
which fine carbides with a particle diameter smaller than 10 nm
which contained at least one element selected from Ti, Mo, and Nb
and V were dispersed.
[0076] For steel pipes Nos. 44 to 48, which were comparative
examples, the manufacturing method thereof was outside the range of
the present invention, and for steel pipes Nos. 49 to 51, the
chemical composition thereof was outside the range of the present
invention. Therefore, for these steel pipes, the microstructure
thereof was not substantially a ferrite phase, and fine carbides
containing at least one element selected from Ti, Mo, and Nb and V
were not precipitated, so that there caused a problem in that a
sufficient strength was not obtained and cracking occurred in the
HIC test.
[0077] Even if heat treatment was accomplished by either the
induction heating apparatus or the gas-fired furnace, no difference
in result was recognized.
4TABLE 4 C/(Mo + Steel Ti + Nb + re- type C Si Mn P S Mo Ti Nb V Al
Cu Ni Cr Ca W Ceq V + Zr)* marks a 0.050 0.19 1.23 0.006 0.0010
0.14 0.018 0.014 0.046 0.028 0.29 1.44 Ex- b 0.042 0.21 1.30 0.008
0.0005 0.20 0.035 0.035 0.032 0.30 1.10 ample c 0.039 0.26 1.45
0.010 0.0008 0.16 0.019 0.049 0.051 0.035 0.0025 0.32 0.90 d 0.052
0.27 1.55 0.010 0.0010 0.22 0.038 0.027 0.042 0.036 0.21 0.36 0.81
e 0.063 0.32 1.31 0.002 0.0008 0.37 0.028 0.024 0.045 0.026 0.0021
0.36 0.94 f 0.045 0.21 1.26 0.008 0.0006 0.24 0.012 0.035 0.030
0.033 0.16 0.12 0.0024 0.33 1.01 g 0.045 0.24 1.19 0.006 0.0009
0.06 0.081 0.024 0.027 0.26 1.46 Com- h 0.055 0.16 1.28 0.007
0.0006 0.15 0.002 0.023 0.045 0.035 0.0022 0.31 1.67 para- i 0.049
0.33 1.21 0.009 0.0015 0.02 0.018 0.015 0.026 0.25 5.47 tive Unit:
mass %. *: at % Underline indicates outside the range of this
invention
[0078]
5TABLE 5 Reheating Rolling Cooling Max. Min. HIC Toughness Outside
Wall Heating finish Cooling stop temp. temp. No. of Yield Tensile
re- of welded Steel diameter thickness temp. temp. rate temp. Tmax
Tmin Time cycles strength strength sist- part vTrs re- No. type
(mm) (mm) (.degree. C.) (.degree. C.) (.degree. C./s) (.degree. C.)
Method (.degree. C.) (.degree. C.) (min) (cycle) (MPa) (MPa) Grade
ance (.degree. C.) marks 36 a 914.4 19.1 1200 880 30 580 Induction
650 600 4.2 3 502 597 X70 .largecircle. -65 Ex- heating furnace
ample 37 b 914.4 19.1 1150 920 30 640 Gas-fired 660 -- 7.3 1 532
654 X70 .largecircle. -43 furnace 38 b 914.4 19.1 1200 900 30 590
Induction 630 580 5.5 4 515 633 X70 .largecircle. -48 heating
furnace 39 c 1219.2 25.4 1150 850 22 570 Induction 650 630 3.2 2
508 612 X70 .largecircle. -57 heating furnace 40 c 1219.2 25.4 1150
930 22 680 Induction 680 620 9.2 5 528 646 X70 .largecircle. -60
heating furnace 41 d 914.4 19.1 1200 920 30 600 Gas-fired 650 --
6.0 1 604 731 X80 .largecircle. -42 furnace 42 e 914.4 19.1 1200
880 30 620 Induction 630 600 5.1 3 648 795 X80 .largecircle. -44
heating furnace 43 f 914.4 19.1 1100 900 30 630 Gas-fired 650 --
4.5 1 545 655 X70 .largecircle. -70 heating furnace 44 b 914.4 19.1
1200 850 1 660 Induction 620 570 4.2 3 451 548 X60 .largecircle.
-40 Com- heating furnace para- 45 b 914.4 19.1 1200 850 30 750
Gas-fired 660 -- 3.3 1 417 526 X60 .largecircle. -44 tive furnace
ex- 46 c 1219.2 25.4 1200 900 55 580 Gas-fired 580 -- 1.5 1 427 518
X60 X -56 ample furnace 47 c 1219.2 25.4 1200 900 55 650 Gas-fired
750 -- 6.1 1 411 497 X52 .largecircle. -61 furnace 48 c 1219.2 25.4
1100 900 60 400 Induction 460 420 -- 2 488 583 X65 X -55 heating
furnace 49 g 914.4 19.1 1150 850 45 650 Induction 650 600 4.3 2 634
742 X80 .largecircle. 18 heating furnace 50 h 914.4 19.1 1150 850
45 650 Induction 650 600 9.1 4 425 513 X60 .largecircle. +89
heating furnace 51 i 914.4 19.1 1150 850 45 650 Gas-fired 650 --
14.5 1 418 501 X52 .largecircle. -58 furnace Underline indicates
outside the range of this invention
* * * * *