Technical Field
[0001] The present invention relates to steel plates having excellent cryogenic toughness
(toughness at cryogenic temperatures). Specifically, the present invention relates
to a steel plate having good toughness [in particular, toughness in a crosswise direction
(width direction; C-direction)] at cryogenic temperatures of -196°C or lower even
when having a reduced Ni content of about 5.0% to about 7.5%. Hereinafter descriptions
will be made while centering on steel plates for liquefied natural gas (LNG), in which
the steel plates are exposed to the cryogenic temperatures. Such steel plates are
represented by those for use in storage tanks and transport ships. It should be noted,
however, that the present invention is not limited to steel plates of this type and
is applicable to all steel plates for use in applications where the steel plates are
exposed to cryogenic temperatures of -196°C or lower.
Background Art
[0002] LNG tank-use steel plates are to be used in liquefied natural gas (LNG) storage tanks
and require not only high strengths but also such high toughness as to endure a cryogenic
temperature of - -196°C. In general, steels are known to have better hardness-toughness
balance especially at low temperatures by the addition of Ni. For this reason, steel
plates having a Ni content of about 9% (9% Ni steel plates) have been used as steel
plates for the use. However, owing to increasing Ni cost in recent years, there are
developed steel plates having excellent cryogenic toughness even though containing
Ni in a lower content of less than 9%.
[0003] Typically, Nonpatent Literature 1 describes the effects of a heat treatment of 6%
Ni steel in the ferrite-austenite two-phase region on low-temperature toughness. Specifically,
a heat treatment (L treatment; lamellarizing) in the ferrite-austenite two-phase region
(between A
c1 and A
c3) is added prior to a tempering treatment of a conventional process. This allows the
formation of a large amount of finely dispersed retained austenite, where the retained
austenite is stable even against impact load at cryogenic temperatures. Accordingly,
the resulting 6% Ni steel can ensure cryogenic toughness at -196°C of equal to or
better than the cryogenic toughness of a 9% Ni steel that has undergone regular quenching
and tempering treatments. The 6% Ni steel has excellent cryogenic toughness in the
L-direction, but, unfortunately, tends to generally have inferior cryogenic toughness
in a crosswise direction (C-direction) as compared with the cryogenic toughness in
a rolling direction (longitudinal direction; L-direction). In addition, the literature
fails to describe percent brittle fracture.
[0004] Techniques as in Nonpatent Literature 1 are also described in Patent Literature 1
and Patent Literature 2. Of these, Patent Literature 1 describes a method. This method
uses a steel having a Ni content of 4.0% to 10% and having austenite grain size and
other factors controlled within predetermined ranges. In the method, the steel is
subjected to a specific treatment one or more times and tempered at a temperature
equal to or lower than the A
c1 transformation temperature. In the treatment, the steel is hot-rolled, heated to
a temperature between A
c1 and A
c3, and then cooled. This treatment corresponds to the L treatment in Nonpatent Literature
1. Patent Literature 2 describes another method. This method employs a steel having
a Ni content of 4.0% to 10% and having a particle size of AlN particles before hot
rolling of 1 pm or less. The method subjects the steel to heat treatments (L treatment
and subsequent tempering treatment) as in Patent Literature 1. The values of impact
energy at -196°C (vE
-196) described in these methods are probably those in the L-direction, and the values
as toughness in the C-direction are not found therein. In addition, the methods fail
to consider the strength and to describe the percent brittle fracture.
[0005] Nonpatent Literature 2 describes development of a 6% Ni steel for LNG tanks, in which
the L treatment (two-phase region quenching treatment) and a thermal-mechanical control
process (TMCP) are employed in combination. This literature describes that the resulting
steel has a satisfactory value of toughness in the rolling direction (L-direction),
but fails to describe the toughness value in the crosswise direction (C-direction).
[0006] In contrast, Patent Literature 3 describes a steel plate for cryogenic temperature
use, where the steel plate has a reduced Ni content, and a method for producing the
steel plate. The steel plate employs a Ni steel plate having a Ni content of greater
than 5.0% to less than 8.0% and having a yield strength of 590 MPa or more at room
temperature. Even in use environments, the steel plate has excellent safety against
fracture equivalent to that of 9% Ni steels. According to the technique disclosed
in Patent Literature 3, a steel ingot is heated at a low temperature for a short time
in a heating process, and the heated steel ingot in a rolling process is subjected
to rough rolling to such a reduction that the steel ingot upon the completion of rough
rolling has a thickness 3 to 8 times as much as the thickness of the product (thickness
of the steel plate after finish rolling). The technique is based on a finding that
a steel plate can have better safety against fracture (namely, can have higher toughness
in a low-temperature environment) when the steel plate is allowed to surely have a
higher yield point in an environment at low temperatures (cryogenic temperatures),
where the resulting steel plate is used in such cryogenic temperature environment.
In the working examples in the literature, the workpiece (slab) is rolled from a slab
thickness of 300 mm down to a finish thickness of 50 mm or less (mostly to a finish
thickness of less than 50 mm). The resulting steel plate, as surely having a relatively
high rolling reduction, has both retained austenite in a certain fraction and a finely
dispersed matrix phase and achieves cryogenic toughness at a level equal to the 9%
Ni steels. However, the steel plate disclosed in Patent Literature 3 has a tensile
strength TS at room temperature of at largest 741 MPa.
[0007] Patent Literature 3 mentions the absorbed energy in the C-direction, but fails to
describe the percent brittle fracture. In addition, the steel plate disclosed in Patent
Literature 3 has a tensile strength TS at room temperature of at largest about 741
MPa.
Citation List
Patent Literature
Nonpatent Literature
Summary of Invention
Technical Problem
[0010] As is described above, on Ni steels having a Ni content of about 5.0% to about 7.5%,
techniques to give Ni steels having excellent cryogenic toughness at -196°C have been
proposed, but sufficient investigations on cryogenic toughness in the C-direction
has not yet been made. Such steels, when allowed to have higher strengths, are advantageous
typically in that they can be designed with larger margin. However, there has been
provided no technique relating to a steel plate having high strengths and excellent
cryogenic toughness.
[0011] In addition, none of the above-mentioned literature makes investigations on percent
brittle fracture. The percent brittle fracture refers to a percentage of brittle fracture
occurring upon the application of a load in a Charpy impact test. In a region where
brittle fracture occurs, energy to be absorbed by the steel until the fracture occurs
is remarkably lowered, and this causes the fracture to easily proceed. To prevent
this and to provide better cryogenic toughness, the steel should very importantly
have not only a better general Charpy impact value (vE
-196), but also a percent brittle fracture of 10% or less. However, there has not yet
been proposed a technique relating to a high-strength steel plate having a high base
metal (steel) strength and also having a percent brittle fracture meeting the condition.
[0012] The present invention has been made in consideration of these circumstances. It is
an object of the present invention to provide a high-strength steel plate that includes
a Ni steel having a Ni content of about 5.0% to about 7.5%, has excellent cryogenic
toughness (in particular, cryogenic toughness in the C-direction) at -196°C, and achieves
a percent brittle fracture of equal to or less than 10%. It is another object of the
present invention to provide a method for producing the steel plate.
Solution to Problem
[0013] The present invention has achieved the objects and provides, in one embodiment (first
embodiment), a steel plate having excellent cryogenic toughness. The steel plate contains,
in mass percent, C in a content of 0.02% to 0.10%, Si in a content of 0.40% or less
(excluding 0%), Mn in a content of 0.50% to 2.0%, P in a content of 0.007% or less
(excluding 0%), S in a content of 0.007% or less (excluding 0%), Al in a content of
0.005% to 0.050%, Ni in a content of 5.0% to 7.5%, N in a content of 0.010% or less
(excluding 0%), and at least one element selected from the group consisting of Cr
in a content of 1.20% or less (excluding 0%) and Mo in a content of 1.0% or less (excluding
0%) with the remainder consisting of iron and inevitable impurities. The steel plate
has a Di value as specified by Formula (1) of 2.5 or more. Formula (1) is defined
by steel chemical compositions and is expressed as follows:
where [C], [Si], [Mn], [Cu], [Ni], [Cr], [Mo], and [V] are contents (in mass percent)
respectively of C, Si, Mn, Cu, Ni, Cr, Mo, and V in the steel. The steel plate includes
a retained austenite phase (retained γ) existing at -196°C in a volume fraction of
2.0% to 12.0%. The steel plate has a retained austenite stabilization parameter as
specified by Formula (2) of 3.1 or more. Formula (2) is defined by chemical compositions
contained in the retained austenite and is expressed as follows:
where <C>, <Mn>, <Al>, <Cu>, <Ni>, <Cr>, <Mo>, and <V> are contents (in mass percent)
respectively of C, Mn, Al, Cu, Ni, Cr, Mo, and V in the retained austenite existing
at -196°C.
[0014] In a preferred embodiment according to the first embodiment of the present invention,
the steel plate may have a retained austenite phase volume fraction-retained austenite
stabilization parameter as specified by Formula (3) of 40 or less. Formula (3) is
defined by the retained austenite phase volume fraction and the retained austenite
stabilization parameter and is expressed as follows:
[0015] In a preferred embodiment of the present invention, the element contents, Di value,
and retained austenite volume fraction of the steel plate may be controlled within
narrower ranges (more specified ranges), and the Mn content in the retained austenite
may be controlled instead of the retained austenite stabilization parameter. This
steel plate can have still higher base metal strengths and can offer satisfactory
cryogenic toughness.
[0016] Specifically, the present invention provides, in another embodiment (second embodiment),
a steel plate having excellent cryogenic toughness. The steel plate contains, in mass
percent, C in a content of 0.02% to 0.10%, Si in a content of 0.40% or less (excluding
0%), Mn in a content of 0.6% to 2.0%, P in a content of 0.007% or less (excluding
0%), S in a content of 0.007% or less (excluding 0%), Al in a content of 0.005% to
0.050%, Ni in a content of 5.0% to 7.5%, N in a content of 0.010% or less (excluding
0%), Mo in a content of 0.30% to 1.0%, and Cr in a content of 1.20% or less (excluding
0%) with the remainder consisting of iron and inevitable impurities. The steel plate
has a Di value as specified by Formula (1) of greater than 5.0. The steel plate includes
a retained austenite phase (retained γ) existing at -196°C in a volume fraction of
2.0% to 5.0%. The retained austenite phase (retained γ) existing at -196°C has a Mn
content of 1.05% or more. In addition, the steel has Mn and Ni contents (in mass percent)
meeting a condition specified by Formula (4):
where [Mn] and [Ni] are contents (in mass percent) respectively of Mn and Ni in the
steel.
[0017] In a preferred embodiment of the present invention, the steel plate may further contain
Cu in a content of 1.0% or less (excluding 0%).
[0018] In a preferred embodiment of the present invention, the steel plate may further contain
at least one selected from the group consisting of Ti in a content of 0.025% or less
(excluding 0%), Nb in a content of 0.100% or less (excluding 0%), and V in a content
of 0.50% or less (excluding 0%).
[0019] In a preferred embodiment of the present invention, the steel plate may further contain
B in a content of 0.0050% or less (excluding 0%).
[0020] In a preferred embodiment of the present invention, the steel plate may further contain
at least one selected from the group consisting of Ca in a content of 0.0030% or less
(excluding 0%) and at least one rare-earth element (REM) in a content of 0.0050% or
less (excluding 0%).
[0021] In a preferred embodiment of the present invention, the steel plate may further contain
Zr in a content of 0.005% or less (excluding 0%).
[0022] The present invention also achieves the objects and further provides, in yet another
embodiment (third embodiment), a method for producing the steel plate according to
the first embodiment of the present invention. The method includes the steps a) and
b). In the step a), a steel plate is formed from a steel and subjected to a heat treatment
(L treatment) in a ferrite-austenite two-phase region (between A
c1 and A
c3). The steel has such controlled chemical compositions and the L treatment is performed
at such a controlled temperature (L treatment temperature) that an L parameter as
specified by Formula (5) is 0.25 to 0.45 and a λ
L parameter as specified by Formula (6) is 7 or less. The step b) is performed after
the L treatment, in which the steel plate is water-cooled down to room temperature
and subjected to a tempering treatment (T treatment) at a temperature equal to or
lower than the A
c1 temperature for 10 to 60 minutes. Formula (5) is defined by the L treatment temperature
and the A
c1 and A
c3 temperatures in the steel. Formula (6) is defined by the L parameter and the steel
chemical compositions. Formulae (5) and (6) are expressed as follows:
where [Mn], [Cr], and [Mo] are contents (in mass percent) respectively of Mn, Cr,
and Mo in the steel.
[0023] In addition, the present invention achieves the objects and provides, in still another
embodiment (fourth embodiment), a method for producing the steel plate according to
the second embodiment of the present invention. In the method, the L treatment temperature
and the steel chemical compositions are adjusted so that an L parameter as specified
by Formula (5) is 0.6 to1.1 and a λ
L parameter as specified by Formula (6) is 0 or less.
Advantageous Effects of Invention
[0024] The present invention can provide high-strength steel plates each including a Ni
steel having a Ni content of about 5.0% to about 7.5%. The steel plates have high
base metal strengths, specifically, have a tensile strength TS of greater than 741
MPa and a yield strength YS of greater than 590 MPa, and preferably have a tensile
strength TS of equal to or greater than 830 MPa and a yield strength YS of equal to
or greater than 690 MPa. Even having such high base metal strengths, the steel plates
have excellent cryogenic toughness (in particular, cryogenic toughness in the C-direction)
at temperatures of -196°C or lower and have a percent brittle fracture of equal to
or less than 10% at -196°C, and preferably have a percent brittle fracture of equal
to or less than 50% at -233°C.
Description of Embodiments
[0025] The present inventors made investigations so as to provide a steel plate that has
a Ni content of 7.5% or less and, when subjected to a Charpy impact test in the C-direction,
has a percent brittle fracture of 10% or less at -196°C, a tensile strength TS of
greater than 741 MPa, and a yield strength YS of greater than 590 MPa.
[0026] In particular, the present inventors made investigations while paying attention to
following points.
[0027] Initially, a production method according to the present invention was designed to
attain cryogenic toughness at the same level as compared with a 9% Ni steel even without
strictly controlling rolling and cooling conditions after T treatment as in the techniques
disclosed in Patent Literatures 1 and 3. Specifically, the chemical compositions of
a material steel were designed in consideration of the case where such rolling reduction
as in Patent Literature 3 is not obtained. Rolling in the present invention was designed
on the assumption that the rolling reduction at a temperature of 830°C or higher is
controlled to about 50% or less, and the rolling reduction at a temperature of 700°C
or higher is controlled to about 85% or less; and that cooling after the tempering
treatment (T treatment) is performed not by water cooling, but by air cooling. The
rolling reduction (%) was calculated as: 100×[(Thickness before rolling)-(Thickness
after rolling)]/(Thickness before rolling).
[0028] The steel (plate) was designed to have cryogenic toughness as evaluated in the C-direction,
in which direction toughness is tend to be hardly ensured as compared with the L-direction.
In addition, the cryogenic toughness was to be evaluated not by absorbed energy, but
by percent brittle fracture so as to surely provide toughness at a certain level.
The steel plate herein was also designed to have a tensile strength (TS) of greater
than 741 MPa. This is because a higher tensile strength TS within specifications is
better in consideration of safety in designing of pressure vessels to be used at cryogenic
temperatures.
[0029] Specifically, the present inventors made intensive investigations so as to provide
a steel plate that is produced under the production conditions, has a percent brittle
fracture at -196°C of equal to or less than 10% in the C-direction in a Charpy impact
test, a tensile strength TS of greater than 741 MPa, and a yield strength YS of greater
than 590 MPa.
[0030] As a result, the present inventors have found the controls of retained austenite
morphology [control conditions (A) and (B)] and λ
L parameter [control condition (C)] as follows; and have found that the controls allow
the steel plate to include stable retained austenite at a certain level and to have
excellent cryogenic toughness, where the stable retained austenite does not transform
into martensite, but plastically deforms during the Charpy impact test. The controls
are expressed as follows:
- (A) The Di value is controlled by an appropriate balance among steel chemical compositions,
as specified by Formula (1). This control is performed so as to surely provide stable
retained austenite (i.e., to provide better stability of retained austenite), because
such stable retained austenite does not transform into martensite, but plastically
deforms during the application of impact at a cryogenic temperature, and thus contributes
to better toughness.
- (B) Steel chemical compositions and an L treatment temperature are balanced by the
L parameter as specified by Formula (5). The L treatment temperature refers to a temperature
of a heat treatment (L treatment) in a ferrite-austenite two-phase region (between
Ac1 and Ac3). After the L treatment, the steel plate is water-cooled down to room temperature,
subjected to a tempering treatment (T treatment) under predetermined conditions, and
then air-cooled. Thus, the volume fraction of retained austenite (retained γ) existing
at -196°C is controlled within the range of 2.0% to 12.0%; and the retained austenite
stabilization parameter as specified by Formula (2) is controlled to 3.1 or more,
where the retained austenite stabilization parameter is defined by chemical compositions
in the retained austenite existing at -196°C. Preferably, the retained austenite volume
fraction-retained austenite stabilization parameter as specified by Formula (3) is
controlled to 40 or less, where the retained austenite volume fraction-retained austenite
stabilization parameter is defined by the retained austenite volume fraction and the
retained austenite stabilization parameter.
- (C) The λL parameter is controlled as specified by Formula (5), where the λL parameter is defined by chemical compositions (Mn, Cr, and Mo) and the L treatment
temperature.
[0031] Specifically, the steel plate according to the embodiment of the present invention
contains, in mass percent, C in a content of 0.02% to 0.10%, Si in a content of 0.40%
or less (excluding 0%), Mn in a content of 0.50% to 2.0%, P in a content of 0.007%
or less (excluding 0%), S in a content of 0.007% or less (excluding 0%), Al in a content
of 0.005% to 0.050%, Ni in a content of 5.0% to 7.5%, N in a content of 0.010% or
less (excluding 0%), and at least one element selected from the group consisting of
Cr in a content of 1.20% or less (excluding 0%) and Mo in a content of 1.0% or less
(excluding 0%), with the remainder consisting of iron and inevitable impurities. The
steel plate has a Di value as specified by Formula (1) of 2.5 or more, a volume fraction
of retained austenite phase (retained γ) existing at -196°C of 2.0% to 12.0%, and
a retained austenite stabilization parameter as specified by Formula (2) of 3.1 or
more. Formula (1) is defined by the steel chemical compositions. Formula (2) is defined
by chemical compositions contained in the retained austenite existing at -196°C. Formulae
(1) and (2) are expressed as follows:
where [C], [Si], [Mn], [Cu], [Ni], [Cr], [Mo], and [V] are contents (in mass percent)
respectively of C, Si, Mn, Cu, Ni, Cr, Mo, and V in the steel,
where <C>, <Mn>, <Al>, <Cu>, <Ni>, <Cr>, <Mo>, and <V> are contents (in mass percent)
respectively of C, Mn, Al, Cu, Ni, Cr, Mo, and V in the retained austenite existing
at -196°C.
1. Steel Chemical Compositions
[0032] Initially, the steel chemical compositions will be described.
Carbon (C): 0.02% to 0.10%
[0033] Carbon (C) is essential to ensure strengths and retained austenite at certain levels.
To have such activities effectively, the carbon content is controlled in its lower
limit to be 0.02% or more, preferably 0.03% or more, and more preferably 0.04% or
more. However, carbon, if added in excess, may cause the steel plate to have excessively
high strengths and to have inferior cryogenic toughness contrarily. To prevent this,
the carbon content is controlled in its upper limit to be 0.10% or less, preferably
0.08% or less, and more preferably 0.06% or less.
Silicon (Si): 0.40% or less (excluding 0%)
[0034] Silicon (Si) is useful as a deoxidizer. However, silicon, if added in excess, may
accelerate the formation of hard martensite islands and may cause the steel plate
to have inferior cryogenic toughness. To prevent this, the Si content is controlled
in its upper limit to be 0.40% or less, preferably 0.35% or less, and more preferably
0.20% or less.
Manganese (Mn): 0.50% to 2.0%
[0035] Manganese (Mn) stabilizes austenite (y) and contributes to a larger amount of retained
austenite. To have such activities effectively, the Mn content is controlled in its
lower limit to be 0.50% or more, preferably 0.6% or more, and more preferably 0.7%
or more. However, Mn, if added in excess, may cause temper embrittlement and cause
the steel plate to fail to surely have desired cryogenic toughness. To prevent this,
the Mn content is controlled in its upper limit to be 2.0% or less, preferably 1.5%
or less, and more preferably 1.3% or less.
Phosphorus (P): 0.007% or less (excluding 0%)
[0036] Phosphorus (P) is an impurity element and causes grain boundary fracture. To ensure
desired cryogenic toughness, the phosphorus content is controlled in its upper limit
to be 0.007% or less, and preferably 0.005% or less. The phosphorus content is preferably
minimized, but it is difficult to industrially control the phosphorus content to be
0%.
Sulfur (S): 0.007% or less (excluding 0%)
[0037] Sulfur (S) is an impurity element and causes grain boundary fracture, as with phosphorus.
To ensure desired cryogenic toughness, the sulfur content is controlled in its upper
limit to be 0.007% or less. As will be indicated in after-mentioned experimental examples,
a steel has an increasing percent brittle fracture with an increasing sulfur content
and fails to achieve desired cryogenic toughness (percent brittle fracture at -196°C
of equal to or less than 10%). The sulfur content is preferably controlled in its
upper limit to be 0.005% or less. The sulfur content is preferably minimized, but
it is difficult to industrially control the sulfur content to be 0%.
Aluminum (Al): 0.005% to 0.050%
[0038] Aluminum (Al) accelerates desulfurization and fixes nitrogen. The steel plate, if
having an insufficient Al content, may have increased contents typically of solute
sulfur and solute nitrogen to thereby have inferior cryogenic toughness. To prevent
this, the Al content is controlled in its lower limit to be 0.005% or more, preferably
0.010% or more, and more preferably 0.015% or more. However, Al, if added in excess,
may cause oxides, nitrides, and other particles to coarsen and may also cause the
steel plate to have inferior cryogenic toughness. To prevent this, the Al content
is controlled in its upper limit to be 0.050% or less, preferably 0.045% or less,
and more preferably 0.04% or less.
Nickel (Ni): 5.0% to 7.5%
[0039] Nickel (Ni) is essential to allow the steel plate to surely include retained austenite
(retained γ) that is useful for better cryogenic toughness. To have such activities
effectively, the Ni content is controlled in its lower limit to be 5.0% or more, preferably
5.2% or more, and more preferably 5.4% or more. However, Ni, if added in excess, may
cause increased cost of the raw material. To prevent this, the Ni content is controlled
in its upper limit to be 7.5% or less, preferably 7.0% or less, more preferably 6.5%
or less, furthermore preferably 6.2% or less, and still more preferably 6.0% or less.
Nitrogen (N): 0.010% or less (excluding 0%)
[0040] Nitrogen (N) causes strain aging and causes the steel plate to have inferior cryogenic
toughness. To prevent this, the nitrogen content is controlled in its upper limit
to be 0.010% or less, preferably 0.006% or less, and more preferably 0.004% or less.
[0041] At least one element selected from the group consisting of chromium (Cr) in a content
of 1.20% or less (excluding 0%) and molybdenum (Mo) in a content of 1.0% or less (excluding
0%)
[0042] Chromium (Cr) and molybdenum (Mo) each contribute to higher strengths. Each of these
elements may be added alone or in combination. To have the activities effectively,
the Cr content is controlled to be 0.05% or more, and the Mo content is controlled
to be 0.01% or more. However, each of the elements, if added in excess, may cause
the steel plate to have excessively high strengths and to fail to ensure desired cryogenic
toughness. To prevent this, the Cr content is controlled in its upper limit to be
1.20% or less, preferably 1.1% or less, more preferably 0.9% or less, and furthermore
preferably 0.5% or less; and the Mo content is controlled in its upper limit to be
1.0% or less, preferably 0.8% or less, and more preferably 0.6% or less.
[0043] The steel plate according to the present invention contains the chemical compositions
as basic compositions with the remainder consisting of iron and inevitable impurities.
[0044] The steel plate according to the present invention may further contain one or more
of following selective chemical compositions so as to further have one or more properties.
Copper (Cu): 1.0% or less (excluding 0%)
[0045] Copper (Cu) stabilizes austenite and contributes to an increased amount of retained
austenite. To have such activities effectively, Cu is preferably contained in a content
of 0.05% or more. However, Cu, if added in excess, may cause the steel plate to have
excessively high strengths and to fail to effectively have desired cryogenic toughness.
To prevent this, the Cu content is preferably controlled in its upper limit to be
1.0% or less, more preferably 0.8% or less, and furthermore preferably 0.7% or less.
[0046] At least one element selected from the group consisting of titanium (Ti) in a content
of 0.025% or less (excluding 0%), niobium (Nb) in a content of 0.100% or less (excluding
0%), and vanadium (V) in a content of 0.50% or less (excluding 0%)
[0047] Titanium (Ti), niobium (Nb), and vanadium (V) each precipitate as carbonitrides and
allows the steel to have higher strengths. Each of these elements may be added alone
or in combination. To have the activities effectively, the Ti, Nb, and V contents
are each preferably controlled to be 0.005% or more. However, each of these elements,
if added in excess, may cause the steel plate to have excessively high strengths and
to fail to ensure desired cryogenic toughness. To prevent this, the Ti content is
preferably controlled in its upper limit to be 0.025% or less, more preferably 0.018%
or less, and furthermore preferably 0.015% or less. Likewise, the Nb content is preferably
controlled in its upper limit to be 0.100% or less, more preferably 0.05% or less,
and furthermore preferably 0.02% or less. The V content is preferably controlled in
its upper limit to be 0.50% or less, more preferably 0.3% or less, and furthermore
preferably 0.2% or less.
Boron (B): 0.0050% or less (excluding 0%)
[0048] Boron (B) element allows the steel to have better hardenability and to thereby have
higher strengths. To have the activities effectively, the boron content is preferably
controlled to be 0.0005% or more. However, boron, if added in excess, may cause the
steel plate to have excessively high strengths and to fail to ensure desired cryogenic
toughness. To prevent this, the boron content is preferably controlled in its upper
limit to be 0.0050% or less, more preferably 0.0030% or less, and furthermore preferably
0.0020% or less.
[0049] At least one element selected from the group consisting of calcium (Ca) in a content
of 0.0030% or less (excluding 0%) and at least one rare-earth element (REM) in a content
of 0.0050% or less (excluding 0%)
[0050] Calcium (Ca) and rare-earth elements (REM) fix solute sulfur and make sulfides harmless.
Each of these elements may be added alone or in combination. Each of these elements,
if present in an insufficient content, may cause the steel plate to have a higher
solute sulfur content and to have inferior toughness. To prevent this, the Ca content
is preferably controlled to be 0.0005% or more; and the REM content is preferably
controlled to be 0.0005% or more. The term "REM content" refers to the content of
one rare-earth element when the one rate-earth element alone as selected from the
rare-earth elements mentioned below is contained; and refers to the total content
of two or more rate-earth elements when the two or more rate-earth elements are contained
Hereinafter this is true for the REM contents. Each of Ca and REM, if added in excess,
may cause particles such as sulfides, oxides, and nitrides to coarsen and may also
cause the steel plate to have inferior toughness. To prevent this, the Ca content
is preferably controlled in its upper limit to be 0.0030% or less, and more preferably
0.0025% or less. Likewise, the REM content is preferably controlled in its upper limit
to be 0.0050% or less, and more preferably 0.0040% or less.
[0051] As used herein the term "REM (rare-earth element)" refers to the group of elements
including lanthanoid elements as well as Sc (scandium) and Y (yttrium). The lanthanoid
elements are fifteen elements from La with atomic number 57 to Lu with atomic number
71 in the periodic table of elements. The steel may contain each of these elements
alone or in combination. Among the rare-earth elements, Ce and La are preferred The
REM may be added in a form not limited Typically, the REM may be added in the form
of a misch metal or in the form of a single element such as Ce or La. The misch metal
mainly contains Ce and La and may for example contain about 70% of Ce and about 20%
to about 30% of La.
Zirconium (Zr): 0.005% or less (excluding 0%)
[0052] Zirconium (Zr) fixes nitrogen. The steel plate, if having an insufficient Zr content,
may have an increased solute nitrogen content and have inferior toughness. To prevent
this, the Zr content is preferably controlled to be 0.0005% or more. However, Zr,
if added in excess, may cause particles such as oxides and nitrides to coarsen and
may cause the steel plate to have inferior toughness. To prevent this, the Zr content
is preferably controlled in its upper limit to be 0.005% or less, and more preferably
0.0040% or less.
[0053] The steel chemical compositions in the embodiment of the present invention have been
described above.
2. Retained Austenite Phase (Retained γ) Volume Fraction
[0054] In addition, the steel plate according to the present invention contains a retained
austenite phase existing at -196°C in a volume fraction of 2.0% to 12.0% (preferably
4.0% to 12.0%).
[0055] For better cryogenic toughness, it is effective to ensure retained austenite in a
certain amount. This is because the retained austenite readily plastically deforms
during an impact test at a cryogenic temperature. To obtain desired cryogenic toughness,
the volume fraction of the retained austenite phase is controlled to 2.0% or more
of the total microstructure (all phases) existing at -196°C. However, the retained
austenite is relatively soft as compared with the matrix phase and, if present in
an excessively large amount, may cause the steel plate to fail to have a yield strength
YS at a predetermined level. To prevent this, the retained austenite phase volume
fraction is controlled in its upper limit to be 12.0% or less. The retained austenite
phase volume fraction is preferably 4.0% or more, and more preferably 6.0% or more
in its lower limit; and is preferably 11.5% or less, and more preferably 11.0% or
less in its upper limit.
[0056] Of phases existing at -196°C, the retained austenite phase is importantly controlled
in its volume fraction in the steel plate according to the present invention. The
other phases than the retained austenite are not limited and may be those generally
present in steel plates. The other phases than retained austenite are exemplified
by bainite, martensite, cementite, and other carbides.
3. Di value
[0057] The steel plate according to the present invention also has a Di value as specified
by Formula (1) of 2.5 or more. Formula (1) is defined by the steel chemical compositions
and is expressed as follows:
where [C], [Si], [Mn], [Cu], [Ni], [Cr], [Mo], and [V] are contents (in mass percent)
respectively of C, Si, Mn, Cu, Ni, Cr, Mo, and V in the steel.
[0058] Formula (1) relates to the Di value indicating hardenability and is described as
Grossmann's equation in Trans. Metall. Soc. AIME, 150 (1942), p. 227. With increasing
amounts of the alloy elements defining Formula (1) as the Di value, the steel plate
is more readily hardened (has a higher Di value) and more readily include a finely
dispersed microstructure. In addition, with an increasing Di value, the steel plate
has higher strengths and more surely has strengths at desired levels. After investigations,
the present inventors have found that there is a correlation between the Di value
and the microstructure size after rolling; and that the Di value may be controlled
to 2.5 or more so as to obtain a finely dispersed microstructure after rolling and
to obtain high strengths at desired levels. Specifically, the Di value is found to
be a parameter as a useful index so as to obtain a finely dispersed rolling microstructure
even with a low rolling reduction in a non-crystallization region. This leads to the
formation of retained austenite in a sufficient volume fraction as a result of the
subsequent heat treatment and thereby ensures stable retained austenite that is useful
for better cryogenic toughness. The Di value is also an effective parameter so as
to ensure good properties even when a reduced process load is applied. The process
load herein is reduced by mitigating production conditions described in Patent Literature
3, such as a lowered rolling reduction at low temperatures in a non-recrystallization
region, and time restriction until cooling start.
[0059] To have such activities effectively, the Di value is controlled to be 2.5 or more.
The steel plate, if having a Di value of less than 2.5, may fail to obtain a finely
dispersed microstructure after rolling and fail to obtain retained austenite in a
predetermined amount. In addition, this steel plate may fail to control the after-mentioned
retained austenite stabilization parameter and retained austenite volume fraction-retained
austenite stabilization parameter at predetermined levels and may thereby fail to
include stable retained austenite phase and to ensure desired cryogenic toughness.
The Di value is preferably 3.0 or more. In contrast, the upper limit of the Di value
is not limited from the above-mentioned viewpoints, but is preferably about 5.0 or
less. This is preferred from the viewpoints typically of cost and in consideration
that the current strength specification range for LNG tank steels is 830 MPa or less.
4. Retained Austenite Stabilization Parameter
[0060] The steel plate according to the present invention also has a retained austenite
stabilization parameter as specified by Formula (2) controlled to be 3.1 or more.
This allows the steel plate to have desired cryogenic toughness. Formula (2) is expressed
as follows:
where <C>, <Mn>, <Al>, <Cu>, <Ni>, <Cr>, <Mo>, and <V> are contents (in mass percent)
respectively of C, Mn, Al, Cu, Ni, Cr, Mo, and V in the retained austenite existing
at -196°C.
[0061] For better cryogenic toughness, it is effective to ensure stable retained austenite
that does not transform into martensite, but plastically deforms during the impact
test, as is described above. This may be achieved probably by ensuring retained austenite
in a sufficient volume fraction before the impact test and by stabilizing (increasing
the stability of) the retained austenite so as not to transform into martensite, but
to plastically deform even upon impact receipt. The present inventors specified the
retained austenite volume fraction within the range from the former viewpoint. In
addition, the present inventors made experiments also from the latter viewpoint. As
a result, they have found that the stability of retained austenite existing at -196°C
is determined by chemical compositions in the retained austenite existing at -196°C;
and that, for better stability, it is effective to control the parameter specified
by Formula (2). A steel, if having a reduced Ni content of 7.5% or less as in the
present invention, generally has inferior hardenability, thereby includes a coarsened
microstructure after rolling, and fails to ensure the retained austenite volume fraction
obtained after the heat treatment, or fails to obtain a Di value at a certain level.
According to the present invention, however, these parameters or conditions can be
appropriately controlled by appropriately controlling the retained austenite stabilization
parameter, where the parameter is determined in consideration of balance among chemical
compositions in the retained austenite. The retained austenite stabilization parameter
is derived with reference to an equation relating to the Ms point (martensite transformation
start temperature).
[0062] To ensure desired cryogenic toughness, the retained austenite stabilization parameter
is controlled in its lower limit to be 3.1 or more, preferably 3.3 or more, more preferably
3.5 or more, and furthermore preferably 3.7 or more. The retained austenite stabilization
parameter is not critical in its upper limit from the viewpoint of providing better
cryogenic toughness.
5. Retained Austenite Volume Fraction-Retained Austenite Stabilization Parameter
[0063] The steel plate according to the present invention preferably has a retained austenite
volume fraction-retained austenite stabilization parameter as specified by Formula
(3) as controlled to be 40 or less so as to ensure still better cryogenic toughness.
Formula (3) is expressed as follows:
[0064] As specified by Formula (3), the parameter is defined by the retained austenite volume
fraction and the retained austenite stabilization parameter. The present inventors
have conceived that the distribution of retained austenite significantly affects the
improvement of cryogenic toughness and have defined the parameter as above, because
the retained austenite plastically deforms during the impact test at a cryogenic temperature
and effectively contributes to better toughness. Specifically, in a steel having a
high retained austenite volume fraction and a high retained austenite stabilization
parameter, retained austenite grains are present at small intervals and are finely
dispersed, and these grains do not transform into martensite, but plastically deform
even at cryogenic temperature. Thus, the steel plate has good cryogenic toughness.
[0065] The retained austenite volume fraction-retained austenite stabilization parameter
is preferably 35 or less, and more preferably 30 or less. From the viewpoint of better
cryogenic toughness, the parameter is preferably minimized. The parameter is not critical
in its lower limit in relation to the cryogenic toughness, but is preferably about
10 or more in consideration of the chemical compositions specified in the present
invention.
[0066] As is demonstrated in after-mentioned Experimental Example 2, the steel plate, when
being controlled to have a retained austenite volume fraction-retained austenite stabilization
parameter within a more appropriate range, can have a percent brittle fracture at
a good level of 50% or less even at a temperature of -233°C, lower than the temperature
of -196°C. Specifically, when the retained austenite volume fraction-retained austenite
stabilization parameter is minimized in its upper limit (about 30 or less), the steel
plate can have a percent brittle fracture at -233°C as reduced to 50% or less.
[0067] The above-described configuration according to the present invention can give a steel
plate that has a percent brittle fracture at -196°C of equal to or less than 10% in
the C-direction in a Charpy impact test and has a tensile strength TS of greater than
741 MPa and a yield strength YS of greater than 590 MPa. The present inventors made
further investigations so as to give a steel plate that has high base metal strengths
and still has satisfactory cryogenic toughness.
[0068] Specifically, the present inventors made investigations in order to provide a steel
plate that has a percent brittle fracture at -196°C of 10% or less in the C-direction
in a Charpy impact test, a tensile strength TS of greater than 830 MPa, and a yield
strength YS of greater than 690 MPa. As a result, the present inventors have found
that steel plate can have higher base metal strengths as above and still have satisfactory
cryogenic toughness when the element contents, Di value, and retained austenite volume
fraction in the steel plate are controlled within more specific ranges, and the Mn
content in the retained austenite is controlled instead of the retained austenite
stabilization parameter.
[0069] The element contents, Di value, and retained austenite volume fraction in the steel
plate are controlled within the more specific ranges as follows.
- (a) The Mn content is controlled in its lower limit to be 0.6%.
- (b) Both Cr and Mo are added as essential elements, and the Mo content is controlled
in its lower limit to be 0.30%.
- (c) The Di value is controlled to be greater than 5.0.
- (d) The volume fraction of retained austenite existing at -196°C is controlled in
its upper limit to be 5.0%.
- (e) The balance between Mn and Ni contents in the steel as specified by Formula (4)
is appropriately controlled Formula (4) is expressed as follows:
where [Mn] and [Ni] are contents (in mass percent) respectively of Mn and Ni in the
steel.
In addition to the conditions, the steel plate, when meeting the condition (f) below,
can have a percent brittle fracture at -196°C of 10% or less, a tensile strength TS
of greater than 830 MPa, and a yield strength YS of greater than 690 MPa even without
controlling the retained austenite stabilization parameter. The condition (f) is expressed
as follows:
- (f) The Mn content in the retained austenite existing at -196°C is controlled to be
1.05% or more.
[0070] The steel plate according to the first embodiment of the present invention has a
percent brittle fracture at -196°C of equal to or less than 10% in the C-direction
in a Charpy impact (absorption) test, a tensile strength TS of greater than 741 MPa,
and a yield strength YS of greater than 590 MPa. The steel plate (according to the
second embodiment) meeting the conditions (a) to (e) (preferably further meeting the
condition (f)) has a configuration different from the steel plate according to the
first embodiment. The conditions (a) to (f) will be described below.
[0071]
- (a) The Mn content lower limit is controlled to be 0.6%.
The Mn content in the steel is controlled to be 0.6% or more so as to have still higher
strengths of a tensile strength TS of greater than 830 MPa and a yield strength YS
of greater than 690 MPa and to have a Mn content in the retained austenite at a certain
level. The Mn content is preferably 0.7% or more in its lower limit.
- (b) Both Cr and Mo are essentially contained, and the Mo content lower limit is controlled
to be 0.30%.
- (c) The Di value is controlled to be greater than 5.0.
To have still higher strengths of a tensile strength TS of greater than 830 MPa and
a yield strength YS of greater than 690 MPa, both Cr and Mo are essentially added,
and Mo is controlled to be contained in a content of 0.30% or more. In addition, the
Di value is controlled to be greater than 5.0.
- (d) The volume fraction of retained austenite existing at -196°C is controlled in
its upper limit to be 5.0%.
The retained austenite phase volume fraction is preferably higher from the viewpoint
of better cryogenic toughness. However, the retained austenite is relatively soft
as compared with the matrix phase and, if contained in an excessively high volume
fraction, may cause the steel plate to fail to surely have a yield strength YS and
a tensile strength TS at the predetermined levels. To prevent this, the retained austenite
volume fraction is controlled in its upper limit to be 5.0%, and preferably 4.8%.
The retained austenite phase volume fraction is preferably controlled in its lower
limit to be 3.0%, and more preferably 3.5%.
- (e) The balance between Mn and Ni contents in the steel as specified by Formula (4)
is appropriately controlled, where Formula (4) is expressed as follows:
[0072] The steel plate, when meeting this condition, can have still better stability of
retained austenite. Hereinafter the condition as specified by Formula (4) is also
referred to as "Ni-Mn balance in the steel" or simply referred to as "Ni-Mn balance".
[0073] The story that led up to Formula (4) will be schematically illustrated below. The
present inventors intended to ensure satisfactory balance between strengths and toughness
at cryogenic temperatures while controlling or reducing the Ni content to be 7.5%
or less. The present inventors found that, for this purpose, it is important to effectively
use Mn among steel chemical compositions, because Mn acts as an austenite-stabilizing
element; and that the balance between Mn and Ni is also important because Ni is contained
in a relatively high content among the steel chemical compositions. Based on these
considerations, the present inventors made investigations on steel design guidelines
so as to have better stability of retained austenite. Specifically, the present inventors
made intensive investigations as including the Di value and a Ms temperature (martensite-start
temperature) in consideration that how the reduction of Ni content affects the hardenability,
how alloy elements are concentrated (enriched) upon the L treatment, and how martensite-austenite
(MA) constituents formed upon impact are refined (reduced in size). As a result, the
present inventors found that the size of MA constituents formed upon impact has a
correlation with the size of the phase as rolled and has a correlation with the Ni
and Mn contents in the steel. Based on the findings, the present inventors made further
investigations and have specified Formula (4) as an index for Ni-Mn balance in the
steel so as to ensure desired strength-toughness balance at cryogenic temperatures.
[0074]
(f) The Mn content in the retained austenite existing at -196°C is controlled to be
1.05% or more.
[0075] The steel plate, when meeting this condition, can have better stability of the retained
austenite and achieve excellent strength-toughness balance at cryogenic temperatures.
[0076] The Mn content in the retained austenite existing at -196°C is preferably 1.40% or
more, and more preferably 1.75% or more. The upper limit of the preferred Mn content
in the retained austenite is not critical in relation with the activities, but is
preferably about 2.50% or less in consideration typically of the Mn content range
in the steel.
[0077] In a preferred embodiment, at least one of the factors (i) the retained austenite
volume fraction, (ii) the Mn content in the retained austenite, and (iii) the λ
L parameter is controlled within a more appropriate range. The steel plate according
to the second embodiment can have a percent brittle fracture at a good level of 50%
or less even at -233°C, lower than -196°C. This was demonstrated in after-mentioned
Experimental Example 4. Specifically, (i) the retained austenite volume fraction is
controlled to be about 3.5% to about 4.8%, (ii) the Mn content in the retained austenite
is controlled to be about 1.40% to about 2.5%, and/or (iii) the λ
L parameter is controlled to be about -10 or less. At least one of these controls allows
the steel to have better toughness at -233°C. In another preferred embodiment, at
least two of the conditions (i) to (iii) are controlled within the ranges, and/or
(i) the Mn content in the retained austenite is further controlled to be 1.75% to
2.50%. The steel plate according to this embodiment can have still better toughness
at -233°C.
[0078] The steel plates according to the embodiments of the present invention have been
illustrated above.
[0079] Next, methods for producing the steel plates according to the embodiments of the
present invention will be illustrated. In an embodiment (third embodiment), the method
for producing the steel plate according to the first embodiment of the present invention
includes the steps a) and b). In the step a), a steel plate is formed from a steel
and is subjected to a heat treatment (L treatment) in a ferrite-austenite two-phase
region (between A
c1 and A
c3). The steel has such chemical compositions, and the L treatment is performed at such
a temperature (L treatment temperature) that an L parameter as specified by Formula
(5) is 0.25 to 0.45 and a λ
L parameter as specified by Formula (6) is 7 or less. Formula (5) is defined by the
L treatment temperature and the A
c1 and A
c3 temperatures in the steel. Formula (6) is defined by the L parameter and the steel
chemical compositions. The step b) is performed after the L treatment. In the step
b), the steel plate is water-cooled down to room temperature and subjected to a tempering
treatment (T treatment) at a temperature equal to or lower than the A
c1 temperature for 10 to 60 minutes. Formulae (5) and (6) are expressed as follows:
where [Mn], [Cr], and [Mo] are contents (in mass percent) respectively of Mn, Cr,
and Mo in the steel.
[0080] The individual steps will be illustrated in detail below.
[0081] The production method according to the third embodiment of the present invention
appropriately controls a rolling process and a subsequent tempering treatment (T treatment)
so as to produce the steel plate meeting the conditions. A steel making process herein
is not limited and can be performed by a generally employed procedure.
[0082] Steps (processes) including the rolling process and subsequent processes that feature
the present invention will be sequentially illustrated in detail below.
[0083] Initially, in a preferred embodiment, the heating temperature is controlled to be
about 900°C to about 1100°C, the finish rolling temperature (FRT) is controlled to
be about 700°C to about 900°C, and the start cooling temperature (SCT) is controlled
to be about 650°C to about 800°C. In this process, the start cooling temperature (SCT)
is preferably controlled within the range within 60 seconds after the finish rolling.
This gives a finely dispersed microstructure after rolling and subsequent cooling,
where the finely dispersed microstructure usefully contributes to better toughness.
[0084] Next, the workpiece is cooled in a temperature range from 800°C to 500°C at an average
cooling rate of about 10°C/s or more. In particular, the average cooling rate in the
temperature range is controlled herein so as to give a finely dispersed microstructure
after cooling. The average cooling rate is not critical in its upper limit.
[0085] The cooling herein is preferably performed at an average cooling rate of about 10°C/s
or more at least in the temperature range. The cooling at the average cooling rate
is preferably stopped at a temperature of 200°C or lower. This can reduce the amount
of untransformed austenite and can give a finely dispersed and homogeneous microstructure.
[0086] After the hot rolling, the workpiece is heated to and held at a temperature in a
ferrite (α)-austenite (y) two-phase region between the A
c1 and A
c3 temperatures and then water-cooled. This process is referred to as an "L treatment",
and the heating and holding temperature in this process is referred to as an "L treatment
temperature". According to the embodiment of the present invention, the L treatment
temperature and chemical compositions in the steel are appropriately controlled so
that the L parameter specified by Formula (5) and the λ
L parameter specified by Formula (6) fall within the predetermined ranges. This control
is performed so as to control the retained austenite volume fraction and the retained
austenite stabilization parameter (preferably retained austenite volume fraction-retained
austenite stabilization parameter) within the ranges specified herein.
[0087] Initially, the L treatment temperature after the hot rolling is preferably controlled
to be within the range of A
c1 to (A
c1+A
c3)/2. This allows alloy elements such as Ni to concentrate in formed austenite phase,
part of which becomes a metastable retained austenite phase that metastably exists
at room temperature. The L treatment, if performed at a temperature of lower than
the A
c1 temperature or higher than [(A
c1+A
c3)/2], may eventually fail to allow the steel plate to have a sufficient volume fraction
of retained austenite existing at -196°C and/or sufficient retained austenite stability
(see Sample Nos. 29 and 30 in Table 2B of after-mentioned Experimental Example 1).
The L treatment temperature is preferably from about 620°C to about 650°C.
[0088] The A
c1 and A
c3 temperatures herein are calculated based on formulae below (see "Koza Gendai no Kinzoku-gaku,
Zairyo-hen 4, Tekko Zairyo" (in Japanese), The Japan Institute of Metals and Materials).
A
c3 temperature = 910-203×[C]
1/2-15.2×[Ni]+44.7×[Si]+104x[V]+31.5×[Mo]-30×[Mn]+11×[Cr]+20×[Cu] where [Mn], [Ni], [Si],
[Cr], [As], [W], [C], [V], [Mo], and [Cu] are contents (in mass percent) respectively
of alloy elements Mn, Ni, Si, Cr, As, W, C, V, Mo, and Cu in the steel. The steel
plate according to the present invention does not contain As and W as steel chemical
compositions, and the calculation according to the formula is performed while defining
[As] and [W] each as 0%.
[0089] The heating at a temperature in the two-phase region is preferably performed for
a time (holding time) of about 10 to about 50 minutes. The heating, if performed for
a time shorter than 10 minutes, may fail to allow alloy elements to sufficiently concentrate
in the austenite phase. In contrast, the heating, if performed for a time longer than
50 minutes, may cause the α phase to be annealed and cause the steel plate to have
lower strengths. The heating time is preferably 30 minutes in its upper limit.
[0090] In addition, the L parameter specified by Formula (5) is herein controlled to be
0.25 to 0.45, where Formula (5) is defined by individual chemical compositions. The
L parameter is defined to efficiently use the alloy element concentration during the
L treatment so as to allow the steel plate to finally have a retained austenite volume
fraction and retained austenite stability both at certain levels. A steel plate having
an L parameter out of the range fails to have a desired retained austenite volume
fraction and/or sufficient retained austenite stability, as demonstrated in after-mentioned
Experimental Examples. The L parameter is preferably 0.28 to 0.42, and more preferably
0.30 to 0.40.
[0091] The λ
L parameter is herein controlled to be 7 or less. The λ
L parameter is defined by the contents of Mn, Cr, and Mo, and the L parameter, as specified
by Formula (6). The λ
L parameter is defined so as to restrain adverse effects of temper embrittlement occurring
in a portion where Mn and Cr are excessively concentrated, where the Mn and Cr concentration
is caused typically by phosphorus segregation at prior austenite grain boundaries
during the L treatment. It is difficult to directly measure the content of phosphorus
segregated at prior austenite grain boundaries. Accordingly, the λ
L parameter can act as, so to speak, an alternate parameter for the content of phosphorus
segregated at prior austenite grain boundaries. A steel including a smaller amount
of phosphorus segregated at prior austenite grain boundaries has a lower λ
L parameter. The λ
L parameter is preferably 0.0 or less, and more preferably -10.0 or less. The λ
L parameter is not critical in its lower limit, but is preferably about -30 or more.
This is preferred in synthetic consideration that the amount of Mo to be added is
preferably minimized from the viewpoint of cost; and that the contents and the L parameter
preferably fall within the specific ranges.
[0092] Specifically, at a temperature of -196°C in the cryogenic temperature range, adverse
effects of trace impurities such as phosphorus readily become obvious, and, when phosphorus
is significantly segregated at prior austenite grain boundaries (i.e., when the λ
L parameter is high), the temper embrittlement probably adversely affects the cryogenic
toughness. Typically, Sample Nos. 1,2, and 25 (all according to the present invention)
in Table 1A in after-mentioned Experimental Example 1 are compared These samples have
retained austenite volume fractions and retained austenite stabilization parameters
at the same levels. Specifically, Sample No. 1 has a retained austenite volume fraction
of 8.0% and a retained austenite stabilization parameter of 3.7. Sample No. 2 has
a retained austenite volume fraction of 9.4% and a retained austenite stabilization
parameter of 3.8. Sample No. 25 has a retained austenite volume fraction of 7.9% and
a retained austenite stabilization parameter of 3.7. However, these samples have significantly
different λ
L parameters of -6.8 (Sample No. 1), -10.9 (Sample No. 2), and 5.2 (Sample No. 25).
Accordingly, among the three samples, Sample No. 2 having the lowest λ
L parameter is most excellent in cryogenic toughness.
[0093] Next, the workpiece is water-cooled down to room temperature and subjected to a tempering
treatment (T treatment).
[0094] The tempering treatment is performed at a temperature of equal to or lower than the
A
c1 temperature for 10 to 60 minutes. Such low-temperature tempering allows carbon to
concentrate in the metastable retained austenite and thereby further stabilizes the
metastable retained austenite phase to give a retained austenite phase that stably
exists at -196°C. In addition, the low-temperature tempering helps the steel plate
to have a low Ms temperature.
[0095] The tempering, if performed at a temperature higher than the A
c1 temperature, may cause the metastable retained austenite phase to decompose into
a ferrite (α) phase and a cementite phase and may cause the steel plate to fail to
include the retained austenite phase existing at -196°C in a sufficient volume fraction,
where the metastable retained austenite phase is formed during holding in the two-phase
region. In contrast, the tempering, if performed at a temperature lower than 540°C
or for a time shorter than 10 minutes, may cause carbon to fail to sufficiently concentrate
into the metastable retained austenite phase and may cause the steel plate to fail
to have a desired volume fraction of retained austenite existing at -196°C. The tempering,
if performed for a time longer than 60 minutes, may cause excessive reduction of α
phase dislocation density and may thereby cause the steel plate to fail to surely
have predetermined strengths (TS and YS) (see Sample No. 33 of Table 2B in Experimental
Example 1).
[0096] The tempering treatment is preferably performed at a temperature of 540°C to 560°C
for a time of 15 minutes to 45 minutes (more preferably 35 minutes or shorter, and
furthermore preferably 25 minutes or shorter).
[0097] After undergoing the tempering treatment as above, the workpiece is cooled down to
room temperature. The cooling after the tempering is performed not by water cooling,
but by air cooling. This is because carbon is concentrated into the retained austenite
during air cooling, and the steel plate cooled by air cooling has a higher retained
austenite stabilization parameter as compared with a steel plate cooled by water cooling.
[0098] Next, a method for producing the steel plate according to the second embodiment of
the present invention will be illustrated.
[0099] This production method according to the fourth embodiment of the present invention
includes the steps a') and b). In the step a), the L treatment temperature and the
steel chemical compositions are adjusted so that an L parameter as specified by Formula
(5) is 0.6 to1.1 and a λ
L parameter as specified by Formula (6) is 0 or less. Formula (6) is defined by the
L parameter and the steel chemical compositions. The step b) is performed after the
L treatment, in which the steel plate is water-cooled down to room temperature and
subjecting to a tempering treatment (T treatment) at a temperature equal to or lower
than the A
c1 temperature for 10 to 60 minutes.
[0100] The individual steps or processes will be described in detail below. However, description
will be omitted for conditions the same as in the method for producing the steel plate
according to the first embodiment of the present invention. The conditions are exemplified
by the conditions in the rolling step, and conditions of the temperature and holding
time of the L treatment.
[0101] The method for producing the steel plate according to the second embodiment of the
present invention controls the L parameter specified by Formula (5) to be 0.6 to 1.1.
The L parameter is defined so as to allow the final steel plate to have a sufficient
retained austenite volume fraction and satisfactory retained austenite stability (in
particular one determined by the Di value and the Mn content in the retained austenite).
In particular, the L parameter is specified in its upper limit to be 1.1 or less from
the viewpoints of the chemical compositions and desired microstructure conditions
of the steel plate according to the embodiment of the present invention. The L treatment
increases the retained austenite stability (namely, allows Mn to be concentrated into
the retained austenite). Conversely, the L treatment causes the Mn content in the
matrix (in the steel) to be reduced. Such reduced Mn content in the matrix steel may
adversely affect the steel plate and may cause the steel plate to have insufficient
strengths or to have a retained austenite volume fraction and retained austenite stability
both at unsatisfactory levels. To prevent this, the L parameter is controlled in its
lower limit (0.6 or more) in the embodiment of the present invention. The L parameter
is preferably 0.7 to 1.0.
[0102] In addition, the λ
L parameter is herein controlled to be 0 or less. The λ
L parameter is defined by the L parameter and the contents of Mn, Cr, and Mo in the
steel, as specified by Formula (6). The λ
L parameter is defined so as to restrain adverse effects of temper embrittlement occurring
in a portion where Mn and Cr are excessively concentrated, where the Mn and Cr concentration
is caused typically by phosphorus segregation at prior austenite grain boundaries
during the L treatment, as is described above. It is difficult to directly measure
the content of phosphorus segregated at prior austenite grain boundaries. Accordingly,
the λ
L parameter can act as, so to speak, an alternate parameter for the content of phosphorus
segregated at prior austenite grain boundaries. A steel including a smaller amount
of phosphorus segregated at prior austenite grain boundaries has a lower λ
L parameter. The λ
L parameter is preferably -10.0 or less. The λ
L parameter is not critical in its lower limit, but is preferably about -30 or more.
This is preferred in synthetic consideration that the amount of Mo to be added is
preferably minimized from the viewpoint of cost; and that the contents and the L parameter
preferably fall within the specific ranges.
[0103] Next, the workpiece is water-cooled down to room temperature and subjected to a tempering
treatment (T treatment).
[0104] The tempering treatment is performed at a temperature of equal to or lower than the
A
c1 temperature for 10 to 60 minutes. Such low-temperature tempering allows carbon to
be concentrated in the metastable retained austenite and thereby further stabilizes
the metastable retained austenite phase to give a retained austenite phase that stably
exists at -196°C, as is described above. In addition, the low-temperature tempering
helps the steel plate to have a low Ms temperature.
[0105] The tempering, if performed at a temperature higher than the A
c1 temperature, may cause the metastable retained austenite phase to decompose into
a ferrite (α) phase and a cementite phase and may cause the steel plate to fail to
include the retained austenite phase at -196°C in a sufficient volume fraction, where
the metastable retained austenite phase is formed during holding in the two-phase
region. In contrast, the tempering, if performed for a time shorter than 10 minutes,
may cause carbon to fail to be sufficiently concentrated into the metastable retained
austenite phase and may cause the steel plate to fail to have a desired volume fraction
of retained austenite existing at -196°C. The tempering, if performed for a time longer
than 60 minutes, may cause excessive reduction of α phase dislocation density and
may thereby cause the steel plate to fail to surely have predetermined strength (TS)
(see Sample No. 7 of Table 2B in after-mentioned Experimental Example 3). The tempering
is preferably performed for a time of 15 minutes to 45 minutes, and more preferably
20 minutes to 35 minutes.
[0106] The tempering is performed at a temperature equal to or lower than the A
c1 temperature, and preferably at a temperature of 510°C to 520°C.
[0107] After undergoing the tempering treatment as above, the workpiece is cooled down to
room temperature. The cooling after the tempering is performed not by water cooling,
but by air cooling. This is because carbon is concentrated into the retained austenite
during air cooling, and the steel plate cooled by air cooling has higher stability
of retained austenite as compared with a steel plate cooled by water cooling.
Experimental Examples
[0108] The present invention will be illustrated in further detail with reference to several
examples (experimental examples) below. It should be noted, however, that the examples
are by no means intended to limit the scope of the invention; that various changes
and modifications can naturally be made therein without deviating from the spirit
and scope of the invention as described herein; and all such changes and modifications
should be considered to be within the scope of the invention.
Experimental Example 1
[0109] This experimental example relates to steel plates having a percent brittle fracture
at - 196°C of equal to or less than 10%, a tensile strength TS of greater than 741
MPa, and a yield strength YS of greater than 590 MPa.
[0110] Molten steels as test samples having chemical compositions given in Table 1 (with
the remainder consisting of iron and inevitable impurities, in mass percent) were
made using a vacuum induction furnace (150-kg VIF). The molten steels were cast, subjected
to hot forging, and yielded ingots of a size of 150 mm by 150 mm by 600 mm. REM used
in this experimental example was a misch metal containing about 50% of Ce and about
25% of La.
[0111] Next, the ingots were heated to 1100°C and rolled at a temperature of 830°C or higher
to a thickness of 75 mm. The workpieces were rolled at a finish rolling temperature
(FRT) of 700°C and water-cooled from a start cooling temperature (SCT) of 650°C within
60 seconds after the finish rolling. Thus, the workpieces were rolled down to a thickness
of 25 mm with a rolling reduction of 83%. The cooling in the range from 800°C down
to 500°C was performed at an average cooling rate of 19°C/s, and the cold rolling
was performed to a stop temperature of 200°C or lower to give steel plates.
[0112] The above-prepared steel plates were each subjected to an L treatment by heating
to and holding at an L treatment temperature given in Table 2 for 30 minutes, followed
by water cooling. The steel plates were further subjected to a T treatment (tempering)
at a temperature (T treatment temperature) for a time (T time) given in Table 2, and
air-cooled down to room temperature.
[0113] The prepared steel plates were evaluated on the amount (volume fraction) of retained
austenite phase existing at -196°C, retained austenite stabilization parameter, tensile
properties (tensile strength TS and yield strength YS), and cryogenic toughness (percent
brittle fracture in the C-direction at -196°C or -233°C) in the following manner.
(1) Measurement of Amount (Volume Fraction) of Retained Austenite Phase Existing at
-196°C
[0114] A test specimen of a size of 10 mm by 10 mm by 55 mm was sampled from each steel
plate at a position one-fourth the thickness, held at the liquid nitrogen temperature
(-196°C) for 5 minutes, and subjected to X-ray diffractometry using a two-dimensional
micro-diffraction X-ray diffractometer (RINT-RAPID II) supplied by Rigaku Corporation.
Next, integrated intensity ratios of peaks of (111), (200), (220), and (311) planes
of the retained austenite phase to peaks of (110), (200), (211), and (220) planes
of the ferrite phase were respectively determined, based on which the (111), (200),
(220), and (311) volume fractions of the retained austenite phase were calculated
and averaged, and the average was defined as a "retained austenite volume fraction".
(2) Retained Austenite Stabilization Parameter Measurement
[0115] To determine the retained austenite stabilization parameter as specified by Formula
(2), the contents of chemical compositions in the retained austenite existing at -196°C
to define Formula (2) were measured Specifically, <C>, <Mn>, <Al>, <Cu>, <Ni>, <Cr>,
<Mo>, and <V> as contents (in mass percent) respectively of C, Mn, Al, Cu, Ni, Cr,
Mo, and V were measured in manners as follows.
(2-1) Measurement of C content <C> in retained austenite existing at -196°C
[0116] Simultaneously with the measurement (1), calibrator silicon (Si) was applied onto
each test sample steel, and a precise γ-Fe lattice constant [ao (in angstrom)] was
determined with angle correction with the Si peak. The carbon (C) content in the retained
austenite was determined by inverse calculation from the precisely determined γ-Fe
lattice constant and the contents of following chemical compositions excluding carbon.
(2-2) Measurement of Ni content <Ni> in retained austenite existing at -196°C
[0117] Test specimens of a size of 10 mm by 10 mm by 55 mm were sampled from each steel
plate at a position one-fourth the thickness, held at the liquid nitrogen temperature
(-196°C) for 5 minutes, and subjected to Ni content measurement using an electron
probe microanalyzer (EPMA) JXA-8500F supplied by JEOL Ltd. at an acceleration voltage
of 15 kV and an applied current of 50 nA with a minimum beam diameter. The measurement
was performed three times per each sample (steel plate), and the maximum among the
measured values was defined as the Ni content in the retained austenite.
(2-3) Measurement of Al content <Al> in retained austenite existing at -196°C
[0118] The Al content in the retained austenite was determined as zero (0) on the assumption
that all Al formed oxides and/or nitrides and existed therein.
(2-4) Measurement of contents <Mn>, <Cu>, <Cr>, <Mo>, and <V> of Mn, Cu, Cr, Mo, and
V in retained austenite existing at -196°C
[0119] In this experimental example, the alloy element contents <Mn>, <Cu>, <Cr>, <Mo>,
and <V> after the L treatment and the subsequent T treatment were considered to be
proportional to the measured Ni content <Ni> as measured by the measurement (2-2)
and were calculated in a manner as follows.
[0120] The Ni content in the heat treatments, i.e., the L treatment and the T treatment
may behave (vary) as specified by the formula:
[0121] The term "driving force of austenite reverse transformation" in the formula was calculated
based on the temperature in the heat treatment using a commercially available computational
software (Thermo-Calc). The term "diffusion constant of each alloy element" in the
formula was calculated based on the temperature and holding time in each heat treatment
using a value found in "Diffusion in Solid Metals and Alloys", H. Mehrer, 1990.
[0122] The term "constant in each heat treatment" was experimentally determined in the following
manner. The measured Ni content after the L treatment and the subsequent T treatment
is specified by the formula as the product of [(Constant in L treatment)×(Driving
force of austenite reverse transformation)×(Ni diffusion constant in L treatment)]
and [(Constant in T treatment)×(Driving force of austenite reverse transformation)×(Ni
diffusion constant in L treatment)]. Specifically, the measured Ni content after the
L treatment and the subsequent T treatment includes both the terms "constant in L
treatment" and "constant in T treatment", and the "constant in T treatment" varies
with the "constant in L treatment". Based on these, the constants in the individual
heat treatments ["constant in L treatment" and "constant in T treatment"] were recursively
determined so that the product be most approximal to the measured Ni content after
the L treatment and the subsequent T treatment. Using the thus-determined constants,
the alloy element contents <Mn>, <Cu>, <Cr>, <Mo>, and <V> were calculated.
(3) Measurement of tensile properties (tensile strength TS and yield strength YS)
[0123] A JIS Z2241 No. 4 test specimen was sampled from each steel plate at a position one-fourth
the thickness in a direction parallel to the C-direction and subjected to a tensile
test by the method prescribed in JIS Z 2241 to measure the tensile strength TS and
yield strength YS. In this experimental example, a sample having a tensile strength
TS of greater than 740 MPa and a yield strength YS of greater than 590 MPa was evaluated
as having satisfactory base metal strengths.
(4) Measurement of cryogenic toughness (percent brittle fracture in C-direction)
[0124] Three Charpy impact test specimens (V-notched test specimens according to JIS Z 2242)
were sampled from each steel plate in a direction parallel to the C-direction each
at a position one-fourth the plate thickness and one-fourth the plate width and at
a position one-fourth the plate thickness and half the plate width. The percents brittle
fracture at -196°C (%) of the test specimens at the two positions were measured by
the method prescribed in JIS Z2242 and were averaged independently. Of the two averages
thus calculated, one indicating inferior properties (namely, one with a higher percent
brittle fracture) was employed. A sample having an employed average of 10% or less
was evaluated as having excellent cryogenic toughness in this experimental example.
[0125] Results of these measurements and evaluations are together indicated in Tables 2A
and 2B.
[Table 1A]
No. |
C |
Si |
Mn |
P |
S |
Al |
Ni |
N |
Cu |
Cr |
Mo |
Ti |
Nb |
V |
B |
Ca |
REM |
Zr |
1 |
0.05 |
0.06 |
0.90 |
<0.005 |
<0.0005 |
0.032 |
5.64 |
0.0033 |
|
0.40 |
0.30 |
|
|
|
|
|
|
|
2 |
0.05 |
0.06 |
1.05 |
<0.004 |
0.001 |
0.030 |
5.66 |
0.0032 |
|
0.42 |
0.43 |
|
|
|
|
|
|
|
3 |
0.05 |
0.06 |
0.55 |
<0.004 |
0.001 |
0.030 |
5.62 |
0.0034 |
|
0.82 |
0.43 |
|
|
|
|
|
|
|
4 |
0.06 |
0.06 |
1.01 |
<0.004 |
0.001 |
0.032 |
5.61 |
0.0033 |
|
0.75 |
|
|
|
|
|
|
|
|
5 |
0.05 |
0.06 |
0.79 |
0.005 |
0.001 |
0.033 |
5.88 |
0.0032 |
|
0.98 |
0.05 |
|
|
|
|
|
|
|
6 |
0.04 |
0.09 |
1.28 |
<0.004 |
0.004 |
0.035 |
6.13 |
0.0039 |
0.72 |
|
0.35 |
|
|
|
|
|
|
|
7 |
0.04 |
0.15 |
0.89 |
<0.004 |
0.001 |
0.032 |
5.72 |
0.0033 |
|
0.52 |
0.33 |
|
|
|
|
|
|
|
8 |
0.09 |
0.05 |
0.87 |
<0.004 |
0.002 |
0.020 |
6.15 |
0.0031 |
0.21 |
0.36 |
0.05 |
|
|
|
|
|
|
|
9 |
0.02 |
0.20 |
1.48 |
<0.004 |
0.001 |
0.030 |
5.33 |
0.0032 |
|
|
0.82 |
|
|
|
|
|
0.0042 |
|
10 |
0.03 |
0.38 |
0.80 |
<0.004 |
0.004 |
0.038 |
7.25 |
0.0031 |
|
|
0.75 |
|
0.027 |
|
|
|
0.0023 |
|
11 |
0.05 |
0.29 |
0.54 |
<0.004 |
0.001 |
0.033 |
6.24 |
0.0038 |
0.15 |
0.13 |
0.67 |
|
|
|
|
|
|
|
12 |
0.04 |
0.12 |
1.58 |
<0.004 |
0.001 |
0.034 |
5.92 |
0.0029 |
|
0.41 |
0.06 |
|
|
|
|
0.0019 |
|
|
13 |
0.05 |
0.06 |
0.79 |
0.007 |
0.002 |
0.038 |
7.08 |
0.0029 |
|
|
0.83 |
|
0.012 |
|
|
|
|
|
14 |
0.04 |
0.19 |
1.19 |
<0.004 |
0.007 |
0.011 |
5.65 |
0.0045 |
|
|
0.52 |
0.012 |
|
|
|
|
|
0.0022 |
15 |
0.07 |
0.10 |
0.66 |
<0.004 |
0.003 |
0.047 |
5.88 |
0.0034 |
|
1.05 |
|
|
|
0.25 |
|
|
|
|
16 |
0.04 |
0.02 |
0.66 |
0.005 |
0.002 |
0.041 |
5.09 |
0.0033 |
0.51 |
0.50 |
0.53 |
|
|
|
0.0019 |
|
|
|
17 |
0.03 |
0.22 |
1.72 |
<0.004 |
0.001 |
0.019 |
5.70 |
0.0031 |
|
|
0.61 |
|
|
0.02 |
|
|
|
0.0046 |
18 |
0.04 |
0.05 |
0.70 |
<0.004 |
0.001 |
0.031 |
5.41 |
0.0075 |
|
0.71 |
0.22 |
0.016 |
|
|
0.0007 |
0.0011 |
|
|
19 |
0.05 |
0.11 |
0.74 |
<0.004 |
0.003 |
0.035 |
5.86 |
0.0058 |
0.10 |
0.26 |
0.38 |
|
|
|
|
|
|
|
20 |
0.04 |
0.19 |
0.82 |
<0.004 |
0.001 |
0.035 |
5.62 |
0.0033 |
|
1.15 |
|
|
|
0.05 |
|
0.0022 |
|
|
21 |
0.04 |
0.15 |
0.54 |
0.01 |
0.005 |
0.029 |
5.90 |
0.0031 |
|
0.83 |
0.52 |
0.021 |
|
|
|
|
0.0030 |
|
22 |
0.04 |
0.22 |
1.03 |
<0.004 |
0.002 |
0.034 |
6.89 |
0.0035 |
0.42 |
|
0.43 |
|
0.055 |
|
|
|
|
|
23 |
0.05 |
0.08 |
0.65 |
<0.004 |
0.001 |
0.031 |
5.68 |
0.0032 |
|
0.96 |
0.05 |
|
|
0.38 |
|
|
|
|
24 |
0.06 |
0.08 |
0.89 |
0.005 |
0.001 |
0.030 |
6.37 |
0.0032 |
0.15 |
|
0.50 |
|
|
|
0.0033 |
|
|
|
25 |
0.04 |
0.13 |
1.13 |
<0.004 |
0.001 |
0.031 |
6.30 |
0.0032 |
|
0.70 |
|
|
|
0.11 |
|
0.0027 |
|
|
[Table 1B]
No. |
C |
Si |
Mn |
P |
S |
Al |
Ni |
N |
Cu |
Cr |
Mo |
Ti |
Nb |
V |
B |
Ca |
REM |
Zr |
26 |
0.05 |
0.09 |
0.88 |
<0.004 |
0.001 |
0.032 |
5.45 |
0.0035 |
|
|
|
|
|
|
|
|
|
|
27 |
0.11 |
0.10 |
0.75 |
<0.004 |
0.004 |
0.027 |
6.21 |
0.0038 |
|
0.65 |
|
|
|
|
|
|
|
|
28 |
0.05 |
0.08 |
0.89 |
0.008 |
0.001 |
0.033 |
5.75 |
0.0032 |
0.21 |
0.33 |
0.49 |
|
|
|
|
|
|
|
29 |
0.07 |
0.15 |
0.80 |
<0.004 |
0.001 |
0.031 |
6.32 |
0.0033 |
|
0.36 |
0.15 |
|
|
|
|
|
|
|
30 |
0.10 |
0.42 |
1.12 |
<0.004 |
0.001 |
0.031 |
6.25 |
0.0032 |
|
0.20 |
0.12 |
|
|
|
0.0021 |
|
|
|
31 |
0.04 |
0.06 |
1.33 |
0.005 |
0.001 |
0.030 |
6.24 |
0.0032 |
|
|
0.45 |
|
|
|
|
|
|
|
32 |
0.03 |
0.04 |
2.15 |
<0.004 |
0.002 |
0.029 |
6.33 |
0.0032 |
|
0.63 |
|
|
|
|
|
|
|
|
33 |
0.05 |
0.06 |
0.96 |
<0.004 |
0.002 |
0.032 |
5.45 |
0.0040 |
0.50 |
0.29 |
0.18 |
|
|
|
|
|
|
|
34 |
0.05 |
0.08 |
0.48 |
0.005 |
0.001 |
0.032 |
5.79 |
0.0038 |
|
0.62 |
|
0.015 |
|
|
|
|
|
|
35 |
0.04 |
0.10 |
1.42 |
<0.004 |
0.008 |
0.038 |
6.50 |
0.0051 |
|
0.25 |
0.17 |
|
|
|
|
|
|
|
36 |
0.05 |
0.06 |
0.70 |
<0.004 |
0.001 |
0.030 |
5.03 |
0.0030 |
|
0.50 |
0.30 |
|
|
|
|
|
|
|
37 |
0.01 |
0.08 |
1.05 |
<0.004 |
0.001 |
0.052 |
4.80 |
0.0045 |
|
0.83 |
0.46 |
|
|
0.22 |
|
|
|
|
38 |
0.08 |
0.12 |
0.88 |
<0.004 |
0.001 |
0.004 |
6.40 |
0.0107 |
|
0.25 |
0.15 |
|
0.016 |
|
|
|
|
|
39 |
0.05 |
0.08 |
0.79 |
<0.004 |
0.004 |
0.037 |
6.49 |
0.0032 |
1.08 |
0.23 |
0.16 |
|
|
|
|
0.0032 |
|
|
40 |
0.06 |
0.11 |
0.93 |
0.006 |
0.001 |
0.029 |
5.55 |
0.0029 |
|
1.25 |
|
|
|
|
|
|
|
0.0057 |
41 |
0.04 |
0.10 |
1.34 |
<0.004 |
0.001 |
0.045 |
6.25 |
0.0030 |
|
0.59 |
|
|
0.104 |
|
|
|
0.0051 |
|
42 |
0.04 |
0.22 |
1.21 |
<0.004 |
0.003 |
0.037 |
5.90 |
0.0034 |
|
0.05 |
1.07 |
|
|
|
|
|
|
|
43 |
0.05 |
0.12 |
1.05 |
<0.004 |
0.003 |
0.028 |
6.47 |
0.0036 |
|
0.66 |
|
0.027 |
|
|
|
|
0.0021 |
|
44 |
0.05 |
0.32 |
0.82 |
<0.004 |
0.001 |
0.015 |
5.83 |
0.0029 |
|
|
0.15 |
|
|
0.52 |
|
|
|
|
45 |
0.05 |
0.20 |
0.90 |
<0.004 |
0.001 |
0.035 |
6.04 |
0.0033 |
|
0.31 |
0.22 |
|
|
|
0.0052 |
|
|
|
[Table 2A]
No. |
Ac1 |
Ac3 |
(Ac1+Ac3)/2 |
L treatment temperature (°C) |
T treatment temperature (°C) |
T time (min) |
Cooling after tempering |
L parameter |
λL parameter |
Di value |
Retained γ (%) |
Retained γ stabilization parameter |
Retained γ volume fraction-retained γ stabilization parameter |
Cryogenic toughness |
Tensile properties |
Percent brittle fracture at -196 °C (%) |
YS |
TS |
1 |
627 |
769 |
698 |
640 |
550 |
30 |
Air cooling |
0.34 |
-6.8 |
3.5 |
8.0% |
3.7 |
18 |
3 |
720 |
775 |
2 |
625 |
769 |
697 |
640 |
550 |
30 |
Air cooling |
0.35 |
-10.9 |
4.8 |
9.4% |
3.8 |
17 |
0 |
753 |
811 |
3 |
638 |
789 |
714 |
640 |
560 |
30 |
Air cooling |
0.26 |
-12.5 |
4.3 |
5.9% |
3.9 |
21 |
0 |
743 |
799 |
4 |
632 |
756 |
694 |
640 |
550 |
35 |
Air cooling |
0.32 |
4.7 |
3.1 |
8.8% |
4.2 |
16 |
5 |
711 |
765 |
5 |
633 |
767 |
700 |
640 |
550 |
30 |
Air cooling |
0.30 |
2.2 |
3.3 |
7.4% |
4.1 |
18 |
5 |
715 |
769 |
6 |
608 |
767 |
688 |
620 |
550 |
20 |
Air cooling |
0.32 |
-7.7 |
3.2 |
6.8% |
3.6 |
20 |
3 |
717 |
771 |
7 |
630 |
779 |
704 |
640 |
550 |
25 |
Air cooling |
0.32 |
-7.9 |
4.0 |
7.0% |
3.7 |
20 |
3 |
734 |
790 |
8 |
617 |
741 |
679 |
620 |
560 |
20 |
Air cooling |
0.27 |
1.6 |
3.0 |
8.5% |
4.6 |
16 |
0 |
705 |
792 |
9 |
623 |
791 |
707 |
650 |
550 |
40 |
Air cooling |
0.41 |
-23.7 |
3.4 |
12.0% |
4.0 |
15 |
0 |
726 |
781 |
10 |
603 |
781 |
692 |
620 |
550 |
30 |
Air cooling |
0.35 |
-24.1 |
3.3 |
2.3% |
3.3 |
36 |
5 |
717 |
771 |
11 |
622 |
792 |
707 |
630 |
550 |
30 |
Air cooling |
0.29 |
-22.0 |
3.5 |
3.7% |
3.6 |
27 |
0 |
721 |
776 |
12 |
616 |
744 |
680 |
630 |
550 |
15 |
Air cooling |
0.36 |
4.9 |
3.3 |
10.4% |
3.8 |
16 |
5 |
726 |
781 |
13 |
597 |
763 |
680 |
610 |
550 |
15 |
Air cooling |
0.33 |
-27.0 |
3.6 |
3.0% |
3.4 |
32 |
5 |
725 |
779 |
14 |
620 |
773 |
697 |
640 |
550 |
30 |
Air cooling |
0.38 |
-14.0 |
3.1 |
10.1% |
3.6 |
16 |
0 |
710 |
764 |
15 |
637 |
789 |
713 |
640 |
560 |
30 |
Air cooling |
0.27 |
3.5 |
4.7 |
8.3% |
4.5 |
16 |
0 |
752 |
809 |
16 |
639 |
806 |
722 |
640 |
560 |
20 |
Air cooling |
0.26 |
-16.1 |
4.1 |
2.1% |
3.8 |
35 |
0 |
761 |
819 |
17 |
615 |
768 |
691 |
630 |
550 |
20 |
Air cooling |
0.35 |
-15.2 |
4.2 |
9.0% |
3.7 |
17 |
0 |
755 |
813 |
18 |
638 |
783 |
710 |
650 |
550 |
40 |
Air cooling |
0.34 |
-4.4 |
3.0 |
7.5% |
3.7 |
19 |
3 |
709 |
763 |
19 |
624 |
775 |
699 |
640 |
550 |
55 |
Air cooling |
0.36 |
-10.5 |
3.2 |
9.1% |
3.7 |
17 |
0 |
617 |
747 |
20 |
644 |
786 |
715 |
650 |
550 |
20 |
Air cooling |
0.29 |
4.3 |
3.4 |
7.3% |
4.0 |
19 |
5 |
719 |
774 |
21 |
636 |
796 |
716 |
640 |
550 |
20 |
Air cooling |
0.28 |
-15.8 |
4.9 |
5.1% |
3.7 |
23 |
0 |
775 |
834 |
22 |
602 |
766 |
684 |
610 |
550 |
30 |
Air cooling |
0.30 |
-11.7 |
3.3 |
5.8% |
3.6 |
22 |
0 |
716 |
770 |
23 |
639 |
814 |
726 |
640 |
560 |
20 |
Air cooling |
0.26 |
1.5 |
4.7 |
6.2% |
4.1 |
20 |
5 |
753 |
810 |
24 |
608 |
759 |
684 |
620 |
550 |
25 |
Air cooling |
0.33 |
-14.7 |
3.1 |
7.7% |
3.7 |
19 |
0 |
753 |
810 |
25 |
620 |
765 |
692 |
630 |
550 |
20 |
Air cooling |
0.32 |
5.2 |
3.6 |
7.9% |
3.7 |
18 |
10 |
723 |
778 |
[Table 2B]
No. |
Ac1 |
Ac3 |
(Ac1+Ac3)/2 |
L treatment temperature (°C) |
T treatment temperature (°C) |
T time (min) |
Cooling after tempering |
L parameter |
λL parameter |
Di value |
Retained γ (%) |
Retained γ stabilization parameter |
Retained γ volume fraction-retained γ stabilization parameter |
Cryogenic toughness |
Tensile properties |
Percent brittle fracture at-196°C (%) |
YS |
TS |
26 |
624 |
759 |
692 |
650 |
550 |
20 |
Air cooling |
0.44 |
4.3 |
1.0 |
1.8% |
2.5 |
47 |
23 |
656 |
705 |
27 |
624 |
737 |
681 |
630 |
550 |
15 |
Air cooling |
0.30 |
3.5 |
3.4 |
10.0% |
8.0 |
11 |
17 |
850 |
900 |
28 |
624 |
777 |
701 |
640 |
550 |
20 |
Air cooling |
0.35 |
-13.9 |
4.6 |
0.7% |
3.0 |
69 |
21 |
750 |
807 |
29 |
618 |
752 |
685 |
590 |
550 |
35 |
Air cooling |
0.04 |
-3.3 |
3.2 |
1.5% |
1.0 |
82 |
42 |
713 |
767 |
30 |
621 |
742 |
681 |
770 |
550 |
30 |
Air cooling |
1.48 |
9.4 |
4.3 |
0.8% |
0.5 |
158 |
30 |
773 |
831 |
31 |
605 |
752 |
678 |
630 |
500 |
25 |
Air cooling |
0.42 |
-10.6 |
3.0 |
1.5% |
2.5 |
52 |
29 |
713 |
767 |
32 |
605 |
723 |
664 |
610 |
580 |
20 |
Air cooling |
0.29 |
8.6 |
4.0 |
11.8% |
4.0 |
15 |
29 |
709 |
862 |
33 |
627 |
775 |
701 |
630 |
550 |
65 |
Air cooling |
0.27 |
-2.8 |
3.0 |
5.9% |
4.4 |
20 |
10 |
588 |
712 |
34 |
633 |
773 |
703 |
650 |
550 |
20 |
Air cooling |
0.37 |
2.8 |
1.6 |
1.2% |
2.0 |
65 |
34 |
672 |
723 |
35 |
605 |
741 |
673 |
620 |
550 |
20 |
Air cooling |
0.36 |
0.0 |
3.4 |
9.7% |
3.6 |
17 |
16 |
723 |
778 |
36 |
641 |
785 |
713 |
680 |
550 |
20 |
Air cooling |
0.52 |
-7.0 |
3.0 |
0.4% |
3.3 |
86 |
17 |
784 |
825 |
37 |
647 |
835 |
741 |
690 |
550 |
25 |
Air cooling |
0.48 |
-11.0 |
4.2 |
0.3% |
0.5 |
258 |
50 |
623 |
687 |
38 |
613 |
742 |
677 |
620 |
550 |
40 |
Air cooling |
0.30 |
-1.8 |
3.1 |
12.0% |
3.1 |
16 |
35 |
711 |
766 |
39 |
611 |
775 |
693 |
620 |
550 |
20 |
Air cooling |
0.30 |
-2.5 |
3.1 |
6.8% |
3.8 |
20 |
47 |
803 |
900 |
40 |
644 |
767 |
705 |
650 |
550 |
15 |
Air cooling |
0.30 |
4.9 |
4.2 |
9.0% |
4.5 |
16 |
58 |
780 |
906 |
41 |
616 |
745 |
681 |
630 |
550 |
40 |
Air cooling |
0.36 |
6.3 |
3.0 |
12.0% |
3.9 |
15 |
55 |
775 |
907 |
42 |
618 |
788 |
703 |
630 |
550 |
20 |
Air cooling |
0.32 |
-34.0 |
6.0 |
6.9% |
3.6 |
20 |
12 |
816 |
906 |
43 |
617 |
747 |
682 |
620 |
550 |
15 |
Air cooling |
0.27 |
4.4 |
3.1 |
6.5% |
4.1 |
19 |
24 |
817 |
902 |
44 |
625 |
825 |
725 |
640 |
550 |
35 |
Air cooling |
0.33 |
-2.2 |
3.1 |
7.7% |
3.6 |
19 |
21 |
788 |
894 |
45 |
622 |
765 |
694 |
630 |
550 |
20 |
Air cooling |
0.30 |
-4.2 |
3.2 |
6.5% |
3.7 |
20 |
25 |
809 |
895 |
Considerations can be made from Tables 2A and 2B as follows.
[0126] Sample Nos. 1 to 25 in Table 2A are samples meeting all the conditions specified
in the present invention. These samples could provide steel plates having excellent
cryogenic toughness at -196°C even though having high base metal strengths. Specifically,
the samples each had an average of percent brittle fracture in the C-direction of
equal to or less than 10%.
[0127] In contrast, Sample Nos. 26 to 45 in Table 2B are comparative examples not meeting
one or more of the conditions specified in the present invention, because the samples
did not meet either one of the steel chemical compositions and the preferred production
conditions specified in the present invention. The samples failed to have desired
property or properties.
[0128] Sample No. 26 had a Di value not meeting the condition specified in the present invention.
The sample failed to have a desired retained austenite volume fraction and had a low
retained austenite stabilization parameter. In addition, the sample had a retained
austenite volume fraction-retained austenite stabilization parameter of greater than
the predetermined range. As a result, the sample had a high percent brittle fracture
and failed to achieve desired cryogenic toughness at -196°C. In addition, the sample
had a low Di value and therefore had a low tensile strength TS.
[0129] Sample No. 27 employed Steel No. 27 in Table 1B having an excessively high carbon
content and had inferior cryogenic toughness.
[0130] Sample No. 28 employed Steel No. 28 in Table 1B having an excessively high phosphorus
content. The sample failed to have a desired retained austenite volume fraction and
had a low retained austenite stabilization parameter. In addition, the sample had
a retained austenite volume fraction-retained austenite stabilization parameter of
greater than the predetermined range and, as a result, had inferior cryogenic toughness.
[0131] Sample No. 29 employed Steel No. 29 in Table 1B having chemical compositions meeting
the conditions specified in the present invention, but underwent heating at a temperature
lower than the two-phase region temperature (L treatment temperature), and had a low
L parameter. The sample therefore contained retained austenite in an insufficient
volume fraction and had a low retained austenite stabilization parameter. In addition,
the sample had a retained austenite volume fraction-retained austenite stabilization
parameter of greater than the predetermined range. As a result, the sample had inferior
cryogenic toughness.
[0132] Sample No. 30 employed Steel No. 30 in Table 1B having an excessively high Si content,
underwent heating at a temperature higher than the two-phase region temperature (L
treatment temperature), and had excessively high L parameter and λ
L parameter. The sample therefore contained retained austenite in an insufficient volume
fraction, a low retained austenite stabilization parameter, and a retained austenite
volume fraction-retained austenite stabilization parameter of greater than the predetermined
range. As a result, the sample had inferior cryogenic toughness.
[0133] Sample No. 31 employed Steel No. 31 in Table 1B having chemical compositions meeting
the conditions specified in the present invention, but underwent tempering (T treatment)
at an excessively low temperature. The sample therefore contained retained austenite
in an insufficient volume fraction and had a low retained austenite stabilization
parameter. In addition, the sample had a retained austenite volume fraction-retained
austenite stabilization parameter of greater than the predetermined range. As a result,
the sample had inferior cryogenic toughness.
[0134] Sample No. 32 employed Steel No. 32 in Table 1B having an excessively high Mn content
and had an excessively high λ
L parameter. The sample therefore had inferior cryogenic toughness.
[0135] Sample No. 33 employed Steel No. 33 in Table 1B having chemical compositions meeting
the conditions specified in the present invention, but underwent tempering for an
excessively long time (T time). As a result, the sample had low strengths (TS and
YS).
[0136] Sample No. 34 employed Steel No. 34 in Table 1B having an excessively low Mn content
and had an excessively low Di value. The sample failed to have a desired retained
austenite volume fraction and had a low retained austenite stabilization parameter.
In addition, the sample had a retained austenite volume fraction-retained austenite
stabilization parameter of greater than the predetermined range. As a result, the
sample had a high percent brittle fracture and failed to achieve desired cryogenic
toughness at -196°C. The sample also had a low tensile strength TS due to the low
Di value.
[0137] Sample No. 35 employed Steel No. 35 in Table 1B having an excessively high sulfur
content. As a result, the sample had a high percent brittle fracture and failed to
achieve desired cryogenic toughness.
[0138] Sample No. 36 employed Steel No. 36 in Table 1B having chemical compositions meeting
the conditions specified in the present invention, but had an excessively high L parameter.
The sample therefore contained retained austenite in an insufficient volume fraction
and had a retained austenite volume fraction-retained austenite stabilization parameter
of greater than the predetermined range. As a result, the sample had inferior cryogenic
toughness.
[0139] Sample No. 37 employed Steel No. 37 in Table 1B having a low carbon content, a high
Al content, and a low Ni content. The sample therefore contained retained austenite
in an insufficient volume fraction and had a low retained austenite stabilization
parameter. In addition, the sample had a retained austenite volume fraction-retained
austenite stabilization parameter of greater than the predetermined range. As a result,
the sample had inferior cryogenic toughness and also had a low tensile strength TS.
[0140] Sample No. 38 employed Steel No. 38 in Table 1B having a low Al content and a high
nitrogen content and therefore had inferior cryogenic toughness.
[0141] Sample No. 39 employed Steel No. 39 in Table 1B having excessively high contents
of selective compositions Cu and Ca and therefore had inferior cryogenic toughness.
[0142] Sample No. 40 employed Steel No. 40 in Table 1B having excessively high contents
of selective compositions Cr and Zr and therefore had inferior cryogenic toughness.
[0143] Sample No. 41 employed Steel No. 41 in Table 1B having excessively high contents
of selective compositions Nb and REM and therefore had inferior cryogenic toughness.
[0144] Sample No. 42 employed Steel No. 42 in Table 1B having an excessively high content
of selective composition Mo, had a high Di value, and therefore had inferior cryogenic
toughness.
[0145] Sample No. 43 employed Steel No. 43 in Table 1B having an excessively high content
of selective composition Ti and therefore had inferior cryogenic toughness.
[0146] Sample No. 44 employed Steel No. 44 in Table 1B having an excessively high content
of selective composition V and therefore had inferior cryogenic toughness.
[0147] Sample No. 45 employed Steel No. 45 in Table 1B having an excessively high content
of selective composition boron (B) and therefore had inferior cryogenic toughness.
Experimental Example 2
[0148] In this experimental example, part of the samples according to the present invention
used in Experimental Example 1 were examined and evaluated on percent brittle fracture
at -233°C.
[0149] Specifically, each three test specimens were sampled from each of samples given in
Table 3 at a position one-fourth the plate thickness and one-fourth the plate width,
subjected to a Charpy impact test at -233°C by a method mentioned below, an average
of measured percent brittle fracture values was calculated and evaluated. The sample
numbers in Table 3 correspond to the sample numbers (steel numbers) in Tables 1A and
2A.
[0150] In this experimental example, a sample having a percent brittle fracture equal to
or less than 50% was evaluated as being excellent in percent brittle fracture at -233°C.
[0152] Results of these determinations and evaluations are indicated in Table 3.
[Table 3]
No. |
Retained γ volume fraction-retained γ stabilization parameter |
Cryogenic toughness |
Percent brittle fracture at -196°C (%) |
Percent brittle fracture at -233°C (%) |
1 |
18 |
3 |
49 |
2 |
17 |
0 |
32 |
3 |
21 |
0 |
24 |
6 |
20 |
3 |
48 |
8 |
16 |
0 |
40 |
9 |
15 |
0 |
15 |
10 |
36 |
5 |
56 |
14 |
16 |
0 |
21 |
16 |
35 |
0 |
58 |
18 |
19 |
3 |
47 |
20 |
19 |
5 |
43 |
[0153] Sample Nos. 1 to 3, 6, 8, 9, 14, 18, and 20 in Table 3 have satisfactorily low percent
brittle fracture not only at -196°C, but also at a lower temperature of -233°C and
could achieved extremely excellent cryogenic toughness. This is probably because each
of these samples had a low retained austenite volume fraction-retained austenite stabilization
parameter of 21 or less.
[0154] In contrast, Sample Nos. 10 and 16 had higher retained austenite volume fraction-retained
austenite stabilization parameter of about 35 and therefore had higher percent brittle
fracture at -233°C as compared with the above-mentioned samples.
[0155] The experimental results demonstrate that it is effective to minimize the retained
austenite volume fraction-retained austenite stabilization parameter, particularly
among the conditions specified in the present invention, so as to give steel plates
that have good (low) percent brittle fracture not only at -196°C, but also at a lower
temperature of -233°C.
Experimental Example 3
[0156] This experimental example relates to steel plates having a percent brittle fracture
at -196°C of equal to or less than 10%, a tensile strength TS of greater than 830
MPa, and a yield strength YS of greater than 690 MPa.
[0157] Molten steels as test samples having chemical compositions given in Table 4 (with
the remainder consisting of iron and inevitable impurities, in mass percent) were
made using a vacuum induction furnace (150-kg VIF). The molten steels were cast, subjected
to hot forging, and yielded ingots of a size of 150 mm by 150 mm by 600 mm. REM used
in this experimental example was a misch metal containing about 50% of Ce and about
25% of La.
[0158] Next, the ingots were heated to 1100°C and rolled at a temperature of 830°C or higher
to a thickness of 75 mm. The workpieces were rolled at a finish rolling temperature
(FRT) of 700°C and water-cooled from a start cooling temperature (SCT) of 650°C within
60 seconds after the finish rolling. Thus, the workpieces were rolled to a thickness
of 25 mm with a rolling reduction of 85%. The cooling in the range from 800°C down
to 500°C was performed at an average cooling rate of 19°C/s, and the cold rolling
was performed to a stop temperature of 200°C or lower to give steel plates.
[0159] The above-prepared steel plates were each subjected to an L treatment by heating
to and holding at an L treatment temperature given in Table 5 for 30 minutes, followed
by water cooling. The steel plates were further subjected to a T treatment (tempering)
at a temperature (T treatment temperature) for a time (T time) given in Table 2, and
air-cooled down to room temperature.
[0160] The above-prepared steel plates were examined and evaluated on the amount (volume
fraction) of retained austenite phase existing at -196°C, Mn content in the retained
austenite phase, tensile properties (tensile strength TS and yield strength YS), and
cryogenic toughness (percent brittle fracture in the C-direction at -196°C or -233°C).
[0161] The amount (volume fraction) of retained austenite phase existing at 196°C, tensile
properties (tensile strength TS and yield strength YS), and cryogenic toughness (percent
brittle fracture in the C-direction) were measured by the procedure of Experimental
Example 1. How the Mn content in the retained austenite phase existing at -196°C was
measured will be described below.
[0162] An average Mn content in the retained austenite phase was measured by transmission
electron microscopy-energy dispersive X-ray spectroscopy (TEM-EDX) and calculated
by a procedure as follows. The calculation was performed assuming that the retained
austenite phase includes Fe, Mn, and Ni as chemical compositions. An actual retained
austenite phase may include other elements such as C and Si, in addition to Fe, Mn,
and Ni. However, these elements are present in small amounts and are appiroximately
trivial in this experimental example.
[0163] A test specimen of a size of 10 mm by 10 mm by 55 mm was sampled from each steel
plate at a position one-fourth the thickness, held at the liquid nitrogen temperature
(-196°C) for 5 minutes, cut to a size of 10 mm by 10 mm by 2 mm, mechanically polished
to reduce the thickness "t" from 2 mm to 0.1 mm, blanked into a disc having a size
of 3 mm in diameter, electrically polished, and yielded a thin-film specimen. The
above-prepared thin-film specimen was analyzed using a transmission electron microscope
H-800 supplied by Hitachi, Ltd, based on which an austenite phase was identified using
a transition image and a reciprocal lattice, and the Mn content in the austenite phase
was measured using an EDX analyzer F.MAX7000 supplied by HORIBA, Ltd The measurement
using the EDX was performed at an acceleration voltage of 200 kV and a 75000-fold
observation magnification on five points per sample. The measured values at the five
points were averaged, and the average was defined as the Mn content in the retained
austenite.
[0164] In Experimental Example 3, a sample having a tensile strength TS of greater than
830 MPa and a yield strength YS of greater than 690 MPa was evaluated as having excellent
base metal strengths, differing from Experimental Example 1.
[0165] Results of these measurements and evaluations are together indicated in Tables 5A
and 5B.
[Table 4A]
No. |
C |
Si |
Mn |
P |
S |
Al |
Ni |
Mo |
Cr |
N |
Cu |
Ti |
Nb |
V |
B |
Ca |
REM |
Zr |
1 |
0.05 |
0.06 |
1.27 |
<0.004 |
0.001 |
0.032 |
5.66 |
0.43 |
0.42 |
0.0032 |
|
|
|
|
|
|
|
|
2 |
0.05 |
0.06 |
1.27 |
<0.004 |
0.001 |
0.030 |
5.62 |
0.43 |
0.82 |
0.0034 |
|
|
|
|
|
|
|
|
3 |
0.05 |
0.06 |
1.10 |
<0.004 |
0.001 |
0.030 |
5.84 |
0.30 |
0.62 |
0.0033 |
|
|
|
|
|
|
|
|
4 |
0.09 |
0.05 |
0.87 |
<0.004 |
0.002 |
0.020 |
6.15 |
0.38 |
0.36 |
0.0031 |
0.21 |
|
|
|
|
|
|
|
5 |
0.02 |
0.20 |
1.48 |
<0.004 |
0.001 |
0.030 |
5.33 |
0.82 |
0.55 |
0.0032 |
|
|
|
|
|
|
0.0042 |
|
6 |
0.03 |
0.38 |
0.90 |
<0.004 |
0.004 |
0.038 |
7.25 |
0.75 |
0.33 |
0.0031 |
|
|
0.027 |
|
|
|
0.0023 |
|
7 |
0.05 |
0.29 |
0.82 |
<0.004 |
0.001 |
0.033 |
6.24 |
0.67 |
0.55 |
0.0038 |
0.15 |
|
|
|
|
|
|
|
8 |
0.04 |
0.12 |
1.58 |
<0.004 |
0.001 |
0.034 |
5.92 |
0.33 |
0.35 |
0.0029 |
|
|
|
|
|
0.0019 |
|
|
9 |
0.05 |
0.06 |
1.06 |
0.007 |
0.002 |
0.038 |
7.08 |
0.83 |
0.38 |
0.0029 |
|
|
0.012 |
|
|
|
|
|
10 |
0.04 |
0.19 |
1.19 |
<0.004 |
0.007 |
0.011 |
5.65 |
0.52 |
0.33 |
0.0045 |
|
0.012 |
|
|
|
|
|
0.0022 |
11 |
0.07 |
0.10 |
1.10 |
<0.004 |
0.003 |
0.047 |
5.88 |
0.30 |
0.30 |
0.0034 |
|
|
|
0.25 |
|
|
|
|
12 |
0.04 |
0.02 |
1.15 |
0.005 |
0.002 |
0.041 |
5.09 |
0.53 |
0.69 |
0.0033 |
0.51 |
|
|
|
0.0019 |
|
|
|
13 |
0.03 |
0.22 |
1.72 |
<0.004 |
0.001 |
0.019 |
5.70 |
0.61 |
0.33 |
0.0031 |
|
|
|
0.02 |
|
|
|
0.0046 |
14 |
0.04 |
0.05 |
1.09 |
<0.004 |
0.001 |
0.031 |
5.41 |
0.45 |
0.58 |
0.0075 |
|
0.016 |
|
|
0.0007 |
0.0011 |
|
|
15 |
0.05 |
0.11 |
1.06 |
<0.004 |
0.003 |
0.035 |
5.86 |
0.38 |
0.45 |
0.0058 |
0.10 |
|
|
|
|
|
|
|
16 |
0.04 |
0.19 |
1.06 |
<0.004 |
0.001 |
0.035 |
5.62 |
0.30 |
1.15 |
0.0033 |
|
|
|
0.05 |
|
0.0022 |
|
|
17 |
0.04 |
0.15 |
1.07 |
0.006 |
0.005 |
0.029 |
5.90 |
0.52 |
0.83 |
0.0031 |
|
0.021 |
|
|
|
|
0.0030 |
|
18 |
0.04 |
0.22 |
1.07 |
<0.004 |
0.002 |
0.034 |
6.89 |
0.43 |
0.32 |
0.0035 |
0.42 |
|
0.055 |
|
|
|
|
|
19 |
0.05 |
0.08 |
1.10 |
<0.004 |
0.001 |
0.031 |
5.68 |
0.30 |
0.64 |
0.0032 |
|
|
|
0.38 |
|
|
|
|
20 |
0.06 |
0.08 |
1.07 |
0.005 |
0.001 |
0.030 |
6.37 |
0.35 |
0.52 |
0.0032 |
0.15 |
|
|
|
0.0033 |
|
|
|
21 |
0.04 |
0.13 |
1.02 |
<0.004 |
0.001 |
0.031 |
6.30 |
0.38 |
0.70 |
0.0032 |
|
|
|
0.11 |
|
0.0027 |
|
|
[Table 4B]
No. |
C |
Si |
Mn |
P |
S |
Al |
Ni |
Mo |
Cr |
N |
Cu |
Ti |
Nb |
V |
B |
Ca |
REM |
Zr |
1 |
0.05 |
0.09 |
1.00 |
<0.004 |
0.001 |
0.032 |
5.45 |
0.53 |
0.30 |
0.0035 |
|
|
|
|
|
|
|
|
2 |
0.11 |
0.10 |
1.01 |
<0.004 |
0.004 |
0.027 |
6.21 |
0.24 |
0.65 |
0.0038 |
|
|
|
|
|
|
|
|
3 |
0.05 |
0.08 |
1.12 |
0.008 |
0.001 |
0.033 |
5.75 |
0.49 |
0.33 |
0.0032 |
0.21 |
|
|
|
|
|
|
|
4 |
0.07 |
0.15 |
1.08 |
<0.004 |
0.001 |
0.031 |
6.32 |
0.34 |
0.36 |
0.0033 |
|
|
|
|
|
|
|
|
5 |
0.10 |
0.42 |
1.03 |
<0.004 |
0.001 |
0.031 |
6.25 |
0.30 |
0.20 |
0.0032 |
|
|
|
|
0.0021 |
|
|
|
6 |
0.03 |
0.04 |
2.15 |
<0.004 |
0.002 |
0.029 |
8.33 |
0.22 |
0.49 |
0.0032 |
|
|
|
|
|
|
|
|
7 |
0.05 |
0.06 |
0.96 |
<0.004 |
0.002 |
0.032 |
5.45 |
0.36 |
0.55 |
0.0040 |
0.50 |
|
|
|
|
|
|
|
8 |
0.05 |
0.08 |
0.51 |
0.005 |
0.001 |
0.032 |
5.20 |
0.58 |
0.86 |
0.0038 |
|
0.015 |
|
|
|
|
|
|
9 |
0.04 |
0.10 |
1.42 |
<0.004 |
0.008 |
0.038 |
6.50 |
0.35 |
0.34 |
0.0051 |
|
|
|
|
|
|
|
|
10 |
0.01 |
0.08 |
1.18 |
<0.004 |
0.001 |
0.052 |
4.80 |
0.82 |
0.55 |
0.0045 |
|
|
|
0.22 |
|
|
|
|
11 |
0.08 |
0.12 |
1.01 |
<0.004 |
0.001 |
0.004 |
6.40 |
0.20 |
0.57 |
0.0107 |
|
|
0.016 |
|
|
|
|
|
12 |
0.05 |
0.08 |
1.01 |
<0.004 |
0.004 |
0.037 |
6.49 |
0.32 |
0.33 |
0.0032 |
1.08 |
|
|
|
|
0.0032 |
|
|
13 |
0.06 |
0.11 |
1.05 |
0.006 |
0.001 |
0.029 |
5.55 |
0.21 |
1.25 |
0.0029 |
|
|
|
|
|
|
|
0.0057 |
14 |
0.04 |
0.10 |
1.34 |
<0.004 |
0.001 |
0.045 |
6.25 |
0.36 |
0.47 |
0.0030 |
|
|
0.104 |
|
|
|
0.0051 |
|
15 |
0.04 |
0.22 |
1.21 |
<0.004 |
0.003 |
0.037 |
5.90 |
1.07 |
0.10 |
0.0034 |
|
|
|
|
|
|
|
|
16 |
0.05 |
0.12 |
1.02 |
<0.004 |
0.003 |
0.028 |
6.47 |
0.38 |
0.47 |
0.0036 |
|
0.027 |
|
|
|
|
0.0021 |
|
17 |
0.05 |
0.32 |
1.02 |
<0.004 |
0.001 |
0.015 |
5.83 |
0.31 |
0.36 |
0.0029 |
|
|
|
0.52 |
|
|
|
|
18 |
0.05 |
0.20 |
1.07 |
<0.004 |
0.001 |
0.035 |
6.04 |
0.37 |
0.52 |
0.0033 |
|
|
|
|
0.0052 |
|
|
|
19 |
0.05 |
0.17 |
0.95 |
<0.004 |
0.001 |
0.032 |
5.70 |
0.42 |
0.50 |
0.0027 |
|
|
|
|
|
|
|
|
20 |
0.05 |
0.24 |
0.75 |
<0.004 |
0.001 |
0.029 |
5.65 |
0.39 |
0.67 |
0.0030 |
|
|
|
|
|
|
|
|
21 |
0.05 |
0.05 |
1.98 |
<0.004 |
0.001 |
0.035 |
5.10 |
0.25 |
0.30 |
0.0034 |
|
|
|
|
|
|
|
|
[Table 5A]
No. |
Di value |
Mn content in retained γ (%) |
[Mn]-[0.31*(7.20-[Ni])+ 0.50] |
Ac1 (°C) |
Ac3 (°C) |
(Ac1+Ac3) /2 (°C) |
L treatment temperature (°C) |
T treatment temperature (°C) |
T time (min) |
Cooling after tempering |
L parameter |
λL parameter |
Retained γ (%) |
Cryogenic toughness |
Tensile properties |
Percent brittle fracture at -196°C (%) |
YS (MPa) |
TS (MPa) |
1 |
5.5 |
1.66 |
0.29 |
623 |
763 |
693 |
700 |
520 |
30 |
Air cooling |
0.80 |
-7.3 |
2.1% |
3% |
825 |
860 |
2 |
8.0 |
1.66 |
0.27 |
630 |
768 |
699 |
690 |
510 |
30 |
Air cooling |
0.69 |
-7.5 |
3.0% |
3% |
850 |
882 |
3 |
5.1 |
1.44 |
0.11 |
628 |
766 |
697 |
700 |
520 |
30 |
Air cooling |
0.77 |
-3.2 |
2.5% |
3% |
814 |
852 |
4 |
5.6 |
1.14 |
0.04 |
617 |
752 |
885 |
700 |
510 |
20 |
Air cooling |
0.86 |
-8.3 |
2.6% |
7% |
827 |
862 |
5 |
7.4 |
1.93 |
0.39 |
632 |
797 |
714 |
730 |
520 |
40 |
Air cooling |
0.84 |
-21.5 |
4.5% |
0% |
845 |
878 |
6 |
6.2 |
1.17 |
0.42 |
607 |
782 |
695 |
710 |
520 |
30 |
Air cooling |
0.84 |
23.6 |
4.6% |
1% |
836 |
869 |
7 |
7.9 |
1.07 |
0.02 |
626 |
788 |
707 |
720 |
520 |
30 |
Air cooling |
0.83 |
-20.6 |
2.7% |
5% |
849 |
881 |
8 |
5.2 |
2.06 |
0.68 |
615 |
751 |
683 |
700 |
510 |
15 |
Air cooling |
0.87 |
0.0 |
4.3% |
0% |
818 |
854 |
9 |
8.2 |
1.39 |
0.53 |
600 |
759 |
680 |
700 |
520 |
15 |
Air cooling |
0.88 |
-25.6 |
4.7% |
1% |
853 |
884 |
10 |
5.3 |
1.55 |
0.20 |
626 |
776 |
701 |
720 |
520 |
30 |
Air cooling |
0.88 |
-11.5 |
2.7% |
1% |
820 |
857 |
11 |
6.6 |
1.44 |
0.19 |
620 |
777 |
699 |
710 |
510 |
30 |
Air cooling |
0.82 |
-3.3 |
2.2% |
3% |
839 |
873 |
12 |
7.4 |
1.50 |
0.00 |
637 |
793 |
715 |
730 |
520 |
20 |
Air cooling |
0.85 |
-11.8 |
2.2% |
1% |
845 |
878 |
13 |
7.3 |
2.24 |
0.75 |
620 |
771 |
696 |
710 |
510 |
20 |
Air cooling |
0.84 |
-11.1 |
4.6% |
0% |
844 |
877 |
14 |
5.3 |
1.42 |
0.03 |
631 |
777 |
704 |
720 |
510 |
40 |
Air cooling |
0.86 |
-9.1 |
2.0% |
3% |
819 |
856 |
15 |
5.2 |
1.38 |
0.14 |
624 |
768 |
696 |
710 |
510 |
55 |
Air cooling |
0.85 |
-6.6 |
2.9% |
7% |
789 |
835 |
16 |
7.9 |
1.39 |
0.07 |
642 |
788 |
715 |
730 |
510 |
20 |
Air cooling |
0.85 |
-2.2 |
2.7% |
7% |
849 |
881 |
17 |
7.9 |
1.39 |
0.16 |
630 |
780 |
705 |
720 |
510 |
20 |
Air cooling |
0.85 |
-11.9 |
2.7% |
5% |
849 |
881 |
18 |
5.7 |
1.40 |
0.48 |
607 |
768 |
687 |
700 |
510 |
30 |
Air cooling |
0.83 |
-8.9 |
4.7% |
3% |
829 |
864 |
19 |
8.9 |
1.44 |
0.12 |
628 |
805 |
717 |
730 |
510 |
20 |
Air cooling |
0.83 |
-2.8 |
2.0% |
3% |
865 |
892 |
20 |
6.3 |
1.40 |
0.31 |
615 |
755 |
685 |
700 |
520 |
25 |
Air cooling |
0.86 |
-4.9 |
4.1% |
3% |
836 |
870 |
21 |
7.1 |
1.33 |
0.23 |
621 |
780 |
701 |
720 |
520 |
20 |
Air cooling |
0.87 |
-6.5 |
3.3% |
7% |
843 |
876 |
[Table 5B]
No. |
Di value |
Mn content in retained γ (%) |
[Mn]-[0.31*(7.20-[Ni])+ 0.50] |
Ac1 (°C) |
Ac3 (°C) |
(Ac1+Ac3) /2 (°C) |
L treatment temperature (°C) |
T treatment temperature (°C) |
T time (min) |
Cooling after tempering |
L parameter |
λL parameter |
Retained γ (%) |
Cryogenic toughness |
Tensile properties |
Percent brittle fracture at -196°C (%) |
YS (MPa) |
TS (MPa) |
1 |
4.6 |
1.31 |
-0.05 |
628 |
776 |
702 |
720 |
520 |
20 |
Air cooling |
0.87 |
-13.5 |
1.7% |
15% |
798 |
840 |
2 |
7.4 |
1.32 |
0.20 |
621 |
737 |
679 |
690 |
510 |
15 |
Air cooling |
0.84 |
-1.1 |
2.6% |
25% |
895 |
928 |
3 |
5.5 |
1.46 |
0.16 |
622 |
770 |
696 |
710 |
510 |
20 |
Air cooling |
0.84 |
-10.9 |
2.0% |
63% |
825 |
860 |
4 |
5.6 |
1.41 |
0.30 |
615 |
749 |
682 |
590 |
520 |
35 |
Air cooling |
0.06 |
-9.1 |
0.5% |
55% |
826 |
861 |
5 |
5.6 |
1.35 |
0.24 |
622 |
750 |
686 |
790 |
510 |
30 |
Air cooling |
1.56 |
0.2 |
1.8% |
20% |
828 |
862 |
6 |
5.7 |
2.81 |
1.38 |
602 |
728 |
665 |
680 |
580 |
20 |
Air cooling |
0.87 |
9.6 |
13.2% |
70% |
829 |
863 |
7 |
5.4 |
1.25 |
-0.09 |
632 |
783 |
707 |
720 |
510 |
65 |
Air cooling |
0.83 |
-6.5 |
4.5% |
13% |
765 |
823 |
8 |
5.1 |
0.67 |
-0.62 |
647 |
802 |
724 |
740 |
510 |
20 |
Air cooling |
0.85 |
-18.9 |
1.5% |
27% |
814 |
851 |
9 |
5.1 |
1.85 |
0.70 |
607 |
747 |
677 |
690 |
510 |
20 |
Air cooling |
0.84 |
-2.6 |
4.7% |
55% |
816 |
853 |
10 |
5.3 |
1.54 |
-0.07 |
641 |
840 |
740 |
760 |
510 |
25 |
Air cooling |
0.85 |
-23.9 |
1.1% |
25% |
770 |
806 |
11 |
5.5 |
1.31 |
0.25 |
617 |
743 |
680 |
690 |
520 |
40 |
Air cooling |
0.83 |
0.4 |
4.0% |
23% |
826 |
861 |
12 |
5.6 |
1.32 |
0.29 |
610 |
774 |
692 |
710 |
510 |
20 |
Air cooling |
0.86 |
-4.7 |
4.0% |
19% |
877 |
912 |
13 |
7.6 |
1.38 |
0.04 |
642 |
770 |
706 |
720 |
520 |
15 |
Air cooling |
0.86 |
1.7 |
2.9% |
17% |
896 |
929 |
14 |
5.6 |
1.75 |
0.54 |
614 |
755 |
684 |
700 |
510 |
40 |
Air cooling |
0.86 |
-3.2 |
4.9% |
18% |
877 |
912 |
15 |
6.5 |
1.58 |
0.30 |
618 |
788 |
703 |
720 |
520 |
20 |
Air cooling |
0.85 |
-34.8 |
2.5% |
12% |
889 |
923 |
16 |
5.4 |
1.34 |
0.29 |
614 |
758 |
686 |
700 |
510 |
15 |
Air cooling |
0.85 |
-6.9 |
2.9% |
20% |
873 |
908 |
17 |
8.6 |
1.33 |
0.09 |
629 |
827 |
728 |
750 |
520 |
35 |
Air cooling |
0.86 |
-4.0 |
2.1% |
19% |
908 |
937 |
18 |
5.8 |
1.39 |
0.20 |
624 |
767 |
696 |
710 |
520 |
20 |
Air cooling |
0.85 |
-6.1 |
2.6% |
17% |
880 |
915 |
19 |
5.3 |
1.01 |
-0.02 |
630 |
776 |
703 |
790 |
520 |
20 |
Air cooling |
1.35 |
-7.4 |
1.3% |
14% |
870 |
906 |
20 |
5.2 |
0.98 |
-0.24 |
638 |
787 |
712 |
730 |
640 |
20 |
Air cooling |
0.87 |
-9.2 |
12.3% |
12% |
688 |
774 |
21 |
5.1 |
2.58 |
0.82 |
622 |
741 |
682 |
730 |
520 |
20 |
Air cooling |
1.16 |
10.3 |
4.8% |
64% |
864 |
902 |
Considerations can be made from Tables 5A and 5B as follows.
[0166] Sample Nos. 1 to 21 in Table 5A respectively employed Steel Nos. 1 to 21 in Table
4A having chemical compositions meeting the conditions specified in the present invention
and were prepared under the production conditions specified in the present invention.
These samples could provide steel plates having excellent cryogenic toughness at -196°C
even though having high base metal strengths. Specifically, the samples each had an
average percent brittle fracture in the C-direction of equal to or less than 10%.
[0167] In contrast, Sample Nos. 1 to 21 in Table 5B are comparative examples not meeting
one or more of the conditions specified in the present invention, including the steel
chemical compositions and production conditions, and failed to have desired properties.
[0168] Sample No. 1 in Table 5B employed Steel No. 1 in Table 4B having chemical compositions
meeting the conditions specified in the present invention, but had a Di value not
meeting the condition specified in the present invention. The sample failed to have
a desired retained austenite volume fraction. As a result, the sample had a high percent
brittle fracture and failed to achieve desired cryogenic toughness at -196°C.
[0169] Sample No. 2 in Table 5B employed Steel No. 2 in Table 4B having a high carbon content
and a low Mo content and had inferior cryogenic toughness.
[0170] Sample No. 3 in Table 5B employed Steel No. 3 in Table 4B having an excessively high
phosphorus content and had inferior cryogenic toughness.
[0171] Sample No. 4 in Table 5B employed Steel No. 4 in Table 4B having chemical compositions
meeting the conditions specified in the present invention, but underwent heating at
a temperature lower than the two-phase region temperature (L treatment temperature),
and had a low L parameter. The sample therefore contained retained austenite in an
insufficient volume fraction and had inferior cryogenic toughness.
[0172] Sample No. 5 in Table 5B employed Steel No. 5 in Table 4B having excessively high
Si and Mo contents, underwent heating at a temperature higher than the two-phase region
temperature (L treatment temperature), and had excessively high L parameter and λ
L parameter. The sample therefore contained retained austenite in an insufficient volume
fraction and had inferior cryogenic toughness.
[0173] Sample No. 6 in Table 5B employed Steel No. 6 in Table 4B having a high Mn content
and a low Mo content, underwent tempering at an excessively high temperature (T treatment
temperature), had an excessively high λ
L parameter, and failed to have a desired retained austenite volume fraction. As a
result, the sample had inferior cryogenic toughness.
[0174] Sample No. 7 in Table 5B employed Steel No. 7 in Table 4B having chemical compositions
meeting the conditions specified in the present invention, but underwent tempering
for an excessively long time (T time). The sample had a Ni-Mn balance as specified
by Formula (2) lower than the preferred range. As a result, the sample had inferior
low-temperature toughness (cryogenic toughness) and also had a low strength (TS).
[0175] Sample No. 8 in Table 5B employed Steel No. 8 in Table 4B having an excessively low
Mn content. The sample had a Ni-Mn balance as specified by Formula (2) lower than
the preferred range, had a low Mn content in the retained austenite, and contained
retained austenite in an insufficient volume fraction. As a result, the sample had
inferior cryogenic toughness.
[0176] Sample No. 9 in Table 5B employed Steel No. 9 in Table 4B having an excessively high
sulfur content. As a result, the sample had a high percent brittle fracture and failed
to achieve desired cryogenic toughness.
[0177] Sample No. 10 in Table 5B employed Steel No. 10 in Table 4B having an excessively
low carbon content, an excessively high Al content, and an excessively low Ni content
and having a Ni-Mn balance specified by Formula (2) lower than the preferred range.
The sample contained retained austenite in a low volume fraction because of excessively
low contents of C and Ni that are useful for ensuring retained austenite in a sufficient
volume fraction. As a result, the sample had inferior cryogenic toughness, although
having a yield strength YS at good level. However, the steel had a low tensile strength
TS because of excessively low contents of C and Ni that are effective for higher strength.
[0178] Sample No. 11 in Table 5B employed Steel No. 11 in Table 4B having excessively low
Al and Mo contents and an excessively high nitrogen content and had an excessively
high λ
L parameter. The sample therefore had inferior cryogenic toughness.
[0179] Sample No. 12 in Table 5B employed Steel No. 12 in Table 4B having excessively high
contents of selective compositions Cu and Ca. The sample therefore had inferior cryogenic
toughness.
[0180] Sample No. 13 in Table 5B employed Steel No. 13 in Table 4B having an excessively
low content of Mo and excessively high contents of Cr and Zr each added as selective
compositions and had an excessively high λ
L parameter. The sample therefore had inferior cryogenic toughness.
[0181] Sample No. 14 in Table 5B employed Steel No. 14 in Table 4B having excessively high
contents of selective compositions Nb and REM. The sample therefore had inferior cryogenic
toughness.
[0182] Sample No. 15 in Table 5B employed Steel No. 15 in Table 4B having an excessively
high content of selective composition Mo. The sample therefore had inferior cryogenic
toughness.
[0183] Sample No. 16 in Table 5B employed Steel No. 16 in Table 4B having an excessively
high content of elective composition Ti. The sample therefore had inferior cryogenic
toughness.
[0184] Sample No. 17 in Table 5B employed Steel No. 17 in Table 4B having an excessively
high content of selective composition vanadium (V). The sample therefore had inferior
cryogenic toughness.
[0185] Sample No. 18 in Table 5B employed Steel No. 18 in Table 4B having an excessively
high content of selective composition boron (B). The sample therefore had inferior
cryogenic toughness.
[0186] Sample No. 19 in Table 5B employed Steel No. 19 in Table 4B having chemical compositions
meeting the conditions specified in the present invention, but had an excessively
high L parameter, and underwent an L treatment at an excessively high temperature.
The sample therefore had an excessively low Mn content in the retained austenite,
contained the retained austenite in an insufficient volume fraction, and had inferior
cryogenic toughness.
[0187] Sample No. 20 in Table 5B employed Sample No. 20 in Table 5B having chemical compositions
meeting the conditions specified in the present invention, but underwent tempering
at an excessively high temperature (T treatment temperature), and had a Ni-Mn balance
specified by Formula (2) lower than the preferred range. The sample failed to have
a desired retained austenite volume fraction and had a low Mn content in the retained
austenite. As a result, the sample had a high percent brittle fracture, failed to
achieve desired cryogenic toughness at -196°C, and was inferior in yield strength
YS and tensile strength TS.
[0188] Sample No. 21 in Table 5B employed Steel No. 21 in Table 4B having an excessively
low Mo content and had an excessively high L parameter and an excessively high λ
L parameter. As a result, the sample had a high percent brittle fracture and failed
to achieve desired cryogenic toughness at -196°C.
Experimental Example 4
[0189] In this experimental example, the samples according to the present invention in Table
5A used in Experimental Example 3 were further examined and evaluated on percent brittle
fracture at -233°C.
[0190] Specifically, each three test specimens were sampled from each of samples given in
Table 6 at a position one-fourth the plate thickness and one-fourth the plate width,
subjected to a Charpy impact test at -233°C by a method described below, and an average
of measured percent brittle fracture values was calculated and evaluated. The sample
numbers in Table 6 correspond to the sample numbers (steel numbers) in Tables 4A and
5A In this experimental example, a sample having a percent brittle fracture equal
to or less than 50% was evaluated as being excellent in percent brittle fracture at
-233°C. "
Cryogenic-Temperature Impact Test of Austenitic Stainless Cast Steel", Journal of
the High Pressure Gas Safety Institute of Japan, vol. 24, p. 181.
[0191] Results of the measurements and evaluations are indicated in Table 6. For reference,
Table 6 also indicates data of the (i) retained austenite volume fraction, (ii) Mn
content in retained austenite, and (iii) λ
L parameter as abstracted from Table 5A Details of them are as follows.
[Table 6]
No. |
Retained γ |
Mn content in retained γ |
λL parameter |
Cryogenic toughness |
Percent brittle fracture at -196°C (%) |
Percent brittle fracture at -233°C (%) |
(%) |
(%) |
1 |
2.1% |
1.66 |
-7.3 |
3% |
50% |
2 |
3.0% |
1.66 |
-7.5 |
3% |
50% |
3 |
2.5% |
1.44 |
-3.2 |
3% |
50% |
4 |
2.6% |
1.14 |
-8.3 |
7% |
60% |
5 |
4.5% |
1.93 |
-21.5 |
0% |
15% |
6 |
4.6% |
1.17 |
-23.6 |
1% |
40% |
7 |
2.7% |
1.07 |
-20.6 |
5% |
50% |
8 |
4.3% |
2.06 |
0.0 |
0% |
25% |
9 |
4.7% |
1.39 |
-25.6 |
1% |
40% |
10 |
2.7% |
1.55 |
-11.5 |
1% |
40% |
11 |
2.2% |
1.44 |
-3.3 |
3% |
50% |
12 |
2.2% |
1.50 |
-11.8 |
1% |
40% |
13 |
4.6% |
2.24 |
-11.1 |
0% |
15% |
14 |
2.0% |
1.42 |
-9.1 |
3% |
50% |
15 |
2.9% |
1.38 |
-6.6 |
7% |
60% |
16 |
2.7% |
1.39 |
-2.2 |
7% |
60% |
17 |
2.7% |
1.39 |
-11.9 |
5% |
50% |
18 |
4.7% |
1.40 |
-8.9 |
3% |
50% |
19 |
2.0% |
1.44 |
-2.8 |
3% |
50% |
20 |
4.1% |
1.40 |
-4.9 |
3% |
50% |
21 |
3.3% |
1.33 |
-6.5 |
7% |
60% |
[0192] Sample No. 1 to 3, 5 to 14, and 17 to 20 in Table 6 respectively employed Steel Nos.
1 to 3, 5 to 14, and 17 to 20 in Table 5A meeting at least one of the preferred conditions
(i) to (iii). The samples each had a good percent brittle fracture at -233°C of 50%
or less. In contrast, Sample Nos. 4,15,16, and 21 in Table 6 respectively employed
Steel Nos. 4,15,16, and 21 in Table 5A meeting none of the preferred conditions (i)
to (iii). The samples failed to have desired toughness at -233°C.
[0193] Specifically, Sample Nos. 1 to 3 in Table 6 respectively employed Steel Nos. 1 to
3 in Table 5A and had a good percent brittle fracture at -233°C of 50%.
[0194] In contrast, Sample No. 4 in Table 6 employed Steel No. 4 in Table 5A meeting none
of the preferred conditions (i) to (iii) and failed to have desired toughness at -233°C.
[0195] Sample No. 5 in Table 6 employed Steel No. 5 in Table 5A meeting all the preferred
conditions (i) to (iii) and having a Mn content in the retained austenite controlled
within the more preferred range of 1.75% to 2.50% in the condition (ii). The sample
could have still better toughness (lower percent of brittle fracture) at -233°C of
15%.
[0196] Sample No. 6 in Table 6 employed Steel No. 6 in Table 5A meeting the preferred conditions
(i) and (iii). The sample could have still better toughness (lower percent of brittle
fracture) at -233°C of 40%.
[0197] Sample No. 7 in Table 6 employed Steel No. 7 in Table 5A meeting the preferred condition
(iii) and had a good percent brittle fracture at -233°C of 50%.
[0198] Sample No. 8 in Table 6 employed Steel No. 8 in Table 5A meeting the preferred conditions
(i) and (ii) and having a Mn content in the retained austenite controlled within the
more preferred range of 1.75% to 2.50% in the condition (ii). The sample could have
still better toughness (lower percent of brittle fracture) at -233°C of 25%.
[0199] Sample No. 9 in Table 6 employed Steel No. 9 in Table 5A meeting the preferred conditions
(i) and (iii) and could have still better toughness (lower percent of brittle fracture)
at -233°C of 40%.
[0200] Sample No. 10 in Table 6 employed Steel No. 10 in Table 5A meeting the preferred
conditions (ii) and (iii) and could have still better toughness (lower percent of
brittle fracture) at -233°C of 40%.
[0201] Sample No.11 in Table 6 employed Steel No. 11 in Table 5A meeting the preferred condition
(ii) and had a good percent brittle fracture at -233°C of 50%.
[0202] Sample No. 12 in Table 6 employed Steel No. 12 in Table 5A meeting the preferred
conditions (ii) and (iii) and could have still better toughness (lower percent of
brittle fracture) at -233°C of 40%.
[0203] Sample No. 13 in Table 6 employed Steel No. 13 in Table 5A meeting all the preferred
conditions (i) to (iii) and having a Mn content in the retained austenite controlled
within the more preferred range of 1.75% to 2.50% in the condition (ii). The sample
could have still better toughness (lower percent of brittle fracture) at -233°C of
15%.
[0204] Sample No. 14 in Table 6 employed Steel No. 14 in Table 5A meeting the preferred
condition (ii) and had a good percent brittle fracture at -233°C of 50%.
[0205] In contrast, Sample Nos. 15 and 16 in Table 6 respectively employed Steel Nos. 15
and 16 in Table 5A meeting none of the preferred conditions (i) to (iii) and failed
to have desired toughness at -233°C.
[0206] In contrast, Sample No. 17 in Table 6 employed Steel No. 17 in Table 5A meeting the
preferred condition (iii) and had a good percent brittle fracture at -233°C of 50%.
[0207] Sample No. 18 in Table 6 employed Steel No. 18 in Table 5A meeting the preferred
condition (i) and had a good percent brittle fracture at -233°C of 50%.
[0208] Sample No. 19 in Table 6 employed Steel No. 19 in Table 5A meeting the preferred
condition (ii) and had a good percent brittle fracture at -233°C of 50%.
[0209] Sample No. 20 in Table 6 employed Steel No. 20 in Table 5A meeting the preferred
condition (i) and had a good percent brittle fracture at -233°C of 50%.
[0210] In contrast, Sample No. 21 in Table 6 employed Steel No. 21 in Table 5A meeting none
of the preferred conditions (i) to (iii) and failed to have desired toughness at -233°C.
[0211] While the present invention has been particularly described with reference to specific
embodiments thereof, it is obvious to those skilled in the art that various changes
and modifications may be made without departing from the spirit and scope of the present
invention.
Industrial Applicability
[0213] The steel plates according to the present invention are useful as steel plates that
are in contact with substances at cryogenic temperatures, such as in liquefied natural
gas storage tanks.