TECHNICAL FIELD
[0001] The present invention relates to an austenitic heat resistant alloy, and more specifically
a NiCrFe alloy.
BACKGROUND ART
[0002] Conventionally, facilities such as thermal power generation boilers and chemical
plants are operated in high temperature environments (such as 400 to 800°C) and, in
addition, they are brought into contact with process fluids including sulfides and/or
chlorides. Therefore, materials to be used in such facilities require their creep
strength and corrosion resistance at high temperatures.
[0003] Examples of the material for use in such facilities include 18-8 stainless steel
such as SUS304H, SUS316H, SUS321H, and SUS347H, and NiCrFe alloys represented by Alloy
800H, which is specified as NCF800H by the JIS standard.
[0004] A NiCrFe alloy excels in corrosion resistance and high temperature strength compared
to an 18-8 stainless steel. Further, a NiCrFe alloy excels in economic efficiency
compared to a Ni-base alloy represented by Alloy617. Therefore, NiCrFe alloys are
widely used in regions of severe use environments.
[0005] NiCrFe alloys used in such severe use environments are proposed in Japanese Patent
Application Publication No.
2013-227644 (Patent Literature 1), Japanese Patent Application Publication No.
06-264169 (Patent Literature 2), Japanese Patent Application Publication No.
2002-256398 (Patent Literature 3), and Japanese Patent Application Publication No.
08-13104 (Patent Literature 4).
[0006] An austenitic heat resistant alloy disclosed in Patent Literature 1 consists of,
in mass%, C: less than 0.02%, Si: not more than 2%, Mn: not more than 2%, Cr: not
less than 20% and less than 28%, Ni: more than 35% and not more than 50%, W: 2.0 to
7.0%, Mo: less than 2.5% (including 0%), Nb: less than 2.5% (including 0%), Ti: less
than 3.0% (including 0%), Al: not more than 0.3%, P: not more than 0.04%, S: not more
than 0.01%, and N: not more than 0.05%, with the balance being Fe and impurities,
wherein f1 = (1/2)W + Mo is 1.0 to 5.0, f2 = (1/2)W + Mo + Nb + 2Ti is 2.0 to 8.0,
and f3 = Nb + 2Ti is 0.5 to 5.0.
[0007] A heat resistant and corrosion resistant alloy disclosed in Patent Literature 2 consists
of, in weight%, 55 to 65% of Nickel, 19 to 25% of Chromium, 1 to 4.5% of Aluminum,
0.045 to 0.3% of Yttrium, 0.15 to 1% of Titan, 0.005 to 0.5% of Carbon, 0.1 to 1.5%
of Silicon, not more than 1% of Manganese, a total of 0.005% of at least one element
selected from the group consisting of Magnesium, Calcium, and Cerium, a total of less
than 0.5% of Magnesium and Calcium, less than 1% of Cerium, 0.0001 to 0.1% of Boron,
not more than 0.5% of Zirconium, 0.0001 to 0.2% of Nitrogen, and not more than 10%
of Cobalt, with the balance being Fe and accompanying impurities.
[0008] An austenitic alloy disclosed in Patent Literature 3 contains, in mass%, C: 0.01
to 0.1%, Mn: 0.05 to 2%, Cr: 19 to 26%, and Ni: 10 to 35%, with the Si content satisfying
a formula of 0.01 < Si < (Cr + 0.15 × Ni - 18)/10.
[0009] A heat resistant alloy disclosed in Patent Literature 4 consists of, in weight%,
C: 0.02 to 0.15%, Si: 0.70 to 3.00%, Mn: not more than 0.50%, Ni: 30.0 to 40.0%, Cr:
18.0 to 25.0%, Al: 0.50 to 2.00%, and Ti: 0.10 to 1.00%, with the balance being Fe
and inevitable impurities.
CITATION LIST
PATENT LITERATURE
[0010]
Patent Literature 1: Japanese Patent Application Publication No. 2013-227644
Patent Literature 2: Japanese Patent Application Publication No. 06-264169
Patent Literature 3: Japanese Patent Application Publication No. 2002-256398
Patent Literature 4: Japanese Patent Application Publication No. 08-13104
NON PATENT LITERATURE
SUMMARY OF INVENTION
TECHNICAL PROBLEM
[0012] The austenitic heat resistant alloy disclosed in Patent Literature 1 controls the
formation of Laves phase by specifying the contents of W, Mo, Nb, and Ti, thereby
improving creep strength and toughness. The heat resistant and corrosion resistant
alloy disclosed in Patent Literature 2 improves high-temperature oxidation resistance
by causing γ' to be precipitated during creep. The austenitic alloy disclosed in Patent
Literature 3 improves carburizing properties by suppressing exfoliation of the oxide
film dominantly composed of Cr
2O
3 and formed on the material surface. The heat resistant alloy disclosed in Patent
Literature 4 contains a specific amount of Cr, a reduced amount of Mn, and a fixed
amount of Si, thereby making it possible to obtain excellent oxidation resistance
even in a case in which the Ni content is reduced.
[0013] On the other hand, Non Patent Literature 1 discloses that a NiCrFe alloy has a high
susceptibility to stress relaxation cracking. This means that a NiCrFe alloy requires
stress relief heat treatment, after working, for a bent part and welded part, in which
residual stress is present. Therefore, a NiCrFe alloy requires not only excellent
creep strength but also excellent stress relaxation cracking resistance.
[0014] An objective of the present invention is to provide a NiCrFe alloy which excels
in creep strength and stress relaxation cracking resistance.
SOLUTION TO PROBLEM
[0015] A NiCrFe alloy according to the present invention has a chemical composition consisting
of, in mass%, C: 0.03 to 0.15%, Si: not more than 1.00%, Mn: not more than 2.00%,
P: not more than 0.040%, S: not more than 0.0050%, Cr: 18.0 to 25.0%, Ni: 25.0 to
40.0%, Ti: 0.10 to 1.60%, Al: 0.05 to 1.00%, N: not more than 0.020%, O: not more
than 0.008%, rare earth metal (REM): 0.001 to 0.100%, B: 0 to 0.010%, Ca: 0 to 0.010%,
Mg: 0 to 0.010%, V: 0 to 0.5%, Nb: 0 to 1.0%, Ta: 0 to 1.0%, Hf: 0 to 1.0%, Mo: 0
to 1.0%, W: 0 to 2.0%, Co: 0 to 3.0%, and Cu: 0 to 3.0%, with the balance being Fe
and impurities, the chemical composition satisfying Formulae (1) to (3):
where, each symbol of element in the above described formulae is substituted by the
content (mass%) of the corresponding element. A(REM) in Formula (3) is substituted
by the atomic weight of each rare earth metal.
ADVANTAGEOUS EFFECT OF INVENTION
[0016] A NiCrFe alloy according to the present invention excels in creep strength and stress
relaxation cracking resistance.
BRIEF DESCRIPTION OF DRAWINGS
[0017] FIG. 1 is a diagram to show a relation between fn2 of each Reference mark of Examples,
and a sum (mass%) of γ' and η phase after aging treatment.
DESCRIPTION OF EMBODIMENTS
[0018] The present inventors have conducted detailed study on the creep strength and the
stress relaxation cracking resistance of NiCrFe alloys. As a result, the present inventors
have obtained the following findings.
- (A) To obtain excellent creep strength, the precipitation amount of γ' (intermetallic
compound: Ni3(Ti, Al)), which precipitates during creep under a high-temperature environment, may
be increased. If γ' precipitates sufficiently during creep under a high-temperature
environment, the creep strength of the alloy is increased by precipitation hardening.
However, if γ' precipitates excessively, deformability within an austenite grain deteriorates,
thus causing stress concentration in grain boundary surfaces. As a result, the stress
relaxation cracking resistance of the alloy deteriorates. Therefore, to achieve excellent
creep strength and excellent stress relaxation cracking resistance at the same time,
it is necessary to adjust the amount of γ' which precipitates during creep under a
high temperature environment. To obtain an appropriate amount of γ' precipitation,
the contents of Ti and Al, which constitute γ', may be adjusted.
Specifically, the chemical composition of NiCrFe alloy satisfies Formula (1) to maintain
stress relaxation cracking resistance while ensuring creep strength:
where, each symbol of element in Formula (1) is substituted by the content (mass%)
of the corresponding element.
Now define fn1 as fn1 = Ti + 48Al/27. fn1 is an index to indicate the amount of γ'
that precipitates during creep. fn1 is a total content of Al and Ti, the content of
Al being converted into the amount of Ti. When fn1 is less than 0.50, a sufficient
precipitation amount of γ' will not be obtained. For that reason, the NiCrFe alloy
cannot obtain excellent creep strength. On the other hand, when fn1 is more than 2.20,
the stress relaxation cracking resistance of the NiCrFe alloy will deteriorate due
to a large amount of precipitation of γ'.
- (B) The γ' that has precipitated during creep under a hot-temperature environment
may change in its form over time. Specifically, while fine γ' precipitates in an early
stage of creep, the γ' may change to a coarse and acicular η phase (Ni3Ti) during creep under a high-temperature environment over time. Formation of η phase
will decrease the creep strength of the NiCrFe alloy.
Then, the present inventors have investigated in detail on a case in which γ' phase
changes to η phase under a high-temperature environment. As a result, they supposed
that the Ti content with respect to the total content of Al and Ti, the content of
Al being converted into the amount of Ti is related to the change from γ' phase to
η phase. Accordingly, the present inventors have investigated in detail on the Ti
content with respect to the total content of Al and Ti, the content of Al being converted
into the amount of Ti, and the microstructure during creep.
Now define fn2 as fn2 = Ti/(Ti + 48Al/27). fn2 is a ratio of the Ti content with respect
to the total content of Al and Ti, the content of Al being converted into the amount
of Ti. FIG. 1 shows a relation between fn2 and a sum of γ' and η phase after aging
treatment. FIG. 1 is obtained by the following method. It is created by using fn2,
and Ti, Al, and Ni contents in the γ' and η phase after aging treatment, which are
obtained by the below described method, for NiCrFe alloys whose chemical compositions
are within the range of the present invention, and in which the above described Formula
(1) and the below described Formula (3) are within the range of the present invention.
Further, γ' and η phase are discriminated by using a method to be described below.
The symbol "○" in FIG. 1 indicates an Example in which the number density of η phase
after aging treatment is less than 5/100 µm2. On the other hand, the symbol "●" in FIG. 1 indicates an Example in which the number
density of η phase after aging treatment is not less than 5/100 µm2.
Referring to FIG. 1, when fn2 is less than 0.40, a sufficient precipitation amount
of γ' will not be obtained. In this case, the NiCrFe alloy cannot obtain excellent
creep strength. On the other hand, when fn2 is more than 0.80, γ' changes to η phase.
As a result, the NiCrFe alloy cannot obtain excellent creep strength. Therefore, when
fn2 is 0.40 to 0.80, it is possible to increase the creep strength of the NiCrFe alloy.
As described so far, if the chemical composition of the NiCrFe alloy of the present
invention satisfies Formula (2), γ' precipitates in an appropriate amount, and the
precipitation of η phase will be suppressed even after time has passed so that excellent
creep strength will be obtained:
where, each symbol of element in Formula (2) is substituted by the content (mass%)
of the corresponding element.
- (C) One cause of stress relaxation cracking is segregation of S in grain boundaries.
Therefore, it is possible to increase the stress relaxation cracking resistance of
the NiCrFe alloy by decreasing an impurity S, which segregates in grain boundaries,
thereby causing grain boundary embrittlement. On the other hand, rare earth metals
(REM) combine with a minute amount of S, which cannot be removed by refining, in the
alloy thereby forming inclusions. In other words, a REM can immobilize S as inclusions.
[0019] Therefore, adjusting the content of REM to be an appropriate amount will allow improving
the stress relaxation cracking resistance of the NiCrFe alloy. REM combines with S
and is also likely to combine with O easily. Therefore, to immobilize S by REM, the
REM content should be adjusted while the amount of REM that combines with O is taken
into consideration.
[0020] If the chemical composition of the NiCrFe alloy of the present invention satisfies
Formula (3), S will be sufficiently immobilized by REM, and excellent stress relaxation
cracking resistance will be obtained:
where, each symbol of element in Formula (3) is substituted by the content (mass%)
of the corresponding element, and A(REM) is substituted by the atomic weight of each
rare earth metal.
[0021] ∑[REM/(A(REM))] is substituted by an addition sum of values which are obtained by
dividing each REM content (mass%) contained in the NiCrFe alloy by the atomic weight
of the REM.
[0022] Now define fn3 as fn3 = ∑[REM/(A(REM))] - S/32 - 2/3•O/16. REM is a generic name
of a total of 17 elements of Sc, Y, and lanthanoids. When fn3 is not less than 0,
REM can sufficiently immobilize S as inclusions, thereby improving the stress relaxation
cracking resistance.
[0023] The NiCrFe alloy according to the present invention, which has been completed based
on the above described findings, has a chemical composition consisting of, in mass%,
C: 0.03 to 0.15%, Si: not more than 1.00%, Mn: not more than 2.00%, P: not more than
0.040%, S: not more than 0.0050%, Cr: 18.0 to 25.0%, Ni: 25.0 to 40.0%, Ti: 0.10 to
1.60%, Al: 0.05 to 1.00%, N: not more than 0.020%, O: not more than 0.008%, rare earth
metal (REM): 0.001 to 0.100%, B: 0 to 0.010%, Ca: 0 to 0.010%, Mg: 0 to 0.010%, V:
0 to 0.5%, Nb: 0 to 1.0%, Ta: 0 to 1.0%, Hf: 0 to 1.0%, Mo: 0 to 1.0%, W: 0 to 2.0%,
Co: 0 to 3.0%, and Cu: 0 to 3.0%, with the balance being Fe and impurities, the chemical
composition satisfying Formulae (1) to (3):
where, each symbol of element in Formulae (1) to (3) is substituted by the content
(mass%) of the corresponding element. A(REM) in Formula (3) is substituted by an atomic
weight of each rare earth metal.
[0024] The above described chemical composition may contain B: 0.0001 to 0.010%.
[0025] The above described chemical composition may contain one or two types selected from
the group consisting of Ca: 0.0001 to 0.010%, and Mg: 0.0001 to 0.010%.
[0026] The above described chemical composition may contain one or more types selected from
the group consisting of V: 0.01 to 0.5%, Nb: 0.01 to 1.0%, Ta: 0.01 to 1.0%, and Hf:
0.01 to 1.0%.
[0027] The above described chemical composition may contain one or more types selected from
the group consisting of Mo: 0.01 to 1.0%, W: 0.01 to 2.0%, Co: 0.01 to 3.0%, and Cu:
0.01 to 3.0%.
[0028] The NiCrFe alloy according to the present invention has excellent creep strength
and excellent stress relaxation cracking resistance. To be more specific, the NiCrFe
alloy will not rupture for 300 or more hours even if it is subjected to tensile strain
of 10% at a strain rate of 0.05 min
-1 and is kept as is under air atmosphere of 650°C after being subjected to cold rolling
at a reduction of area of 20%.
[0029] Hereinafter, the NiCrFe alloy according to the present invention will be described
in detail. The symbol"%" regarding elements means, unless otherwise stated, mass%.
[Chemical composition]
[0030] The chemical composition of NiCrFe alloy of the present invention contains the following
elements.
C: 0.03 to 0.15%
[0031] Carbon (C) stabilizes austenite, and increases high temperature creep strength of
the alloy. When the C content is too low, these effects cannot be obtained. On the
other hand, when the C content is too high, coarse carbide will precipitate in large
amount, thus deteriorating the ductility of grain boundaries. Further, the toughness
and creep strength of the alloy decrease. Therefore, the C content is 0.03 to 0.15%.
The lower limit of the C content is preferably 0.04%, more preferably more than 0.04%,
and further preferably 0.05%, and further preferably 0.06%. The upper limit of the
C content is preferably 0.12%, and more preferably 0.10%.
Si: not more than 1.00%
[0032] Silicon (Si) is inevitably contained. Si deoxidizes the alloy, and improves the corrosion
resistance and oxidation resistance at high temperatures of the alloy. However, when
the Si content is too high, the stability of austenite deteriorates, and toughness
and creep strength of the alloy decrease. Therefore, the Si content is not more than
1.00%. The upper limit of the Si content is preferably 0.80%, more preferably 0.60%,
and further preferably less than 0.60%. Excessive reduction of the Si content deteriorates
deoxidization effect, thus deteriorating the corrosion resistance and oxidization
resistance at high temperatures of the alloy. And further, the production cost is
significantly increased. Therefore, the lower limit of the Si content is preferably
0.02%, and more preferably 0.05%.
Mn: not more than 2.00%
[0033] Manganese (Mn) is inevitably contained. Mn deoxidizes the alloy, and stabilizes austenite.
However, when the Mn content is too high, embrittlement is caused and the toughness
and creep ductility of the alloy deteriorate. Therefore, the Mn content is not more
than 2.00%. The upper limit of the Mn content is preferably 1.80%, and more preferably
1.50%. Excessive reduction of the Mn content deteriorates the deoxidization effect
and stabilization of austenite, and further causes significant increase of the production
cost. Therefore, the lower limit of the Mn content is preferably 0.10%, more preferably
0.30%, and further preferably more than 0.50%.
P: not more than 0.040%
[0034] Phosphorous (P) is an impurity. P deteriorates hot workability and weldability of
the alloy, and also deteriorates creep ductility of the alloy after long hours of
usage. Therefore, the P content is not more than 0.040%. The upper limit of the P
content is preferably 0.035%, and more preferably 0.030%. The P content is preferably
as low as possible. However, excessive reduction of the P content will increase the
production cost. Therefore, the lower limit of the P content is preferably 0.0005%,
and more preferably 0.0008%.
S: not more than 0.0050%
[0035] Sulfur (S) is an impurity. S deteriorates the stress relaxation cracking resistance
of the alloy, and also deteriorates the hot workability, weldability, and creep ductility
of the alloy. Therefore, the S content is not more than 0.0050%. The upper limit of
the S content is preferably 0.0030%. The S content is preferably as low as possible.
However, excessive reduction of the S content will increase the production cost. Therefore,
the lower limit of the S content is preferably 0.0002%, and more preferably 0.0003%.
Cr: 18.0 to 25.0%
[0036] Chromium (Cr) improves the oxidation resistance and corrosion resistance at high
temperatures of the alloy. When the Cr content is too low, these effects cannot be
obtained. On the other hand, when the Cr content is too high, the stability of austenite
at high temperatures deteriorates and the creep strength of the alloy decreases. Therefore,
the Cr content is 18.0 to 25.0%. The lower limit of the Cr content is preferably 18.5%,
and more preferably 19.0%. The upper limit of the Cr content is preferably 24.5%,
and more preferably 24.0%.
Ni: 25.0 to 40.0%
[0037] Nickel (Ni) stabilizes austenite structure. Further, Ni forms γ', thereby increasing
the creep strength of the alloy. When the Ni content is too low, γ' is not likely
to be formed, and the aforementioned effect cannot be obtained. On the other hand,
when the Ni content is too high, the production cost increases. Therefore, the Ni
content is 25.0 to 40.0%. The lower limit of the Ni content is preferably 26.0%, and
more preferably 27.0%. The upper limit of the Ni content is preferably 37.0%, and
more preferably 35.0%.
Ti: 0.10 to 1.60%
[0038] Titanium (Ti) combines with Ni to form γ'. Further, Ti combines with C to form TiC,
thereby increasing the creep strength and tensile strength of the alloy at high temperatures.
When the Ti content is too low, such effects cannot be obtained. On the other hand,
when the Ti content is too high, γ' precipitates excessively, thereby deteriorating
the stress relaxation cracking resistance of the alloy. Therefore, the Ti content
is 0.10 to 1.60%. The lower limit of the Ti content is preferably 0.20%, more preferably
0.30%, and further preferably more than 0.60%. The upper limit of the Ti content is
preferably 1.50%, more preferably less than 1.50%, and further preferably 1.40%.
Al: 0.05 to 1.00%
[0039] Aluminum (Ai) deoxidizes the alloy. Further, Al combines with Ni to form γ' and increases
the creep strength and tensile strength of the alloy at high temperatures. When the
Al content is too low, such effects cannot be obtained. On the other hand, when the
Al content is too high, γ' precipitates in a large amount, thereby deteriorating the
stress relaxation cracking resistance, creep ductility, and toughness of the alloy.
Therefore, the Al content is 0.05 to 1.00%. The lower limit of the Al content is preferably
0.08%, and more preferably 0.10%. The upper limit of the Al content is preferably
0.90%, and more preferably 0.80%.
N: not more than 0.020%
[0040] Nitrogen (N) is an impurity. N precipitates as coarse TiN, and decreases the amount
of dissolved Ti, thereby decreasing the creep strength of the alloy. Further, N deteriorates
toughness and hot workability of the alloy. Therefore, the N content is not more than
0.020%. The upper limit of the N content is preferably 0.017%, and more preferably
0.015%. The N content is preferably as low as possible. However, excessive reduction
thereof will increase production cost. Therefore, the lower limit of the N content
is preferably 0.002%, and more preferably 0.004%.
O: not more than 0.008%
[0041] Oxygen (O) is an impurity. Oxygen deteriorates the hot workability of the alloy,
and also deteriorates the toughness and ductility of the alloy. Therefore, the O content
is not more than 0.008%. The upper limit of the O content is preferably 0.006%, and
more preferably 0.005%. The O content is preferably as low as possible. However, excessive
reduction thereof will increase the production cost. Therefore, the lower limit of
the O content is preferably 0.0005%, and more preferably 0.0008%.
REM: 0.001 to 0.100%
[0042] Rare earth metal (REM) forms a compound with S, thereby decreasing the content of
S which has dissolved into the matrix, and improving the stress relaxation cracking
resistance of the alloy. Further, REM improves the hot workability and oxidization
resistance of the alloy. When the REM content is too low, these effects cannot be
obtained. On the other hand, when the REM content is too high, the hot workability
and weldability of the alloy will deteriorate. Therefore, the REM content is 0.001
to 0.100%. The lower limit of the REM content is preferably 0.003%, and more preferably
0.005%. The upper limit of the REM content is preferably 0.090%, and more preferably
0.080%.
[0043] REM is a generic name of a total of 17 elements of Sc, Y, and lanthanoids, and the
REM content refers to a total content of one or more elements of REM. Moreover, REM
is generally contained in misch metal. For that reason, REM may be added to molten
metal as misch metal, and may be adjusted such that the REM content is within the
above described range.
[0044] The balance of the chemical composition of the NiCrFe alloy according to the present
invention consists of Fe and impurities. Here, the term impurity means an element
which is introduced from ores and scraps as the raw material, or from a production
environment, etc., when the NiCrFe alloy is industrially produced, and which is permitted
within a range not adversely affecting the NiCrFe alloy of the present embodiment.
[Optional elements]
[0045] The NiCrFe alloy according to the present invention may contain B in place of part
of Fe.
B: 0 to 0.010%
[0046] Boron (B) is an optional element and may not be contained. When contained, B increases
the creep strength of the alloy by causing grain boundary carbides to be finely dispersed.
Further, B segregates in grain boundaries to assist the effects of REM. When B is
contained in any small amount, the aforementioned effects will be obtained to some
degree. However, when the B content is too high, the weldability and hot workability
of the alloy will deteriorate. Therefore, the B content is 0 to 0.010%. The upper
limit of the B content is preferably 0.008%. The lower limit of the B content to effectively
obtain the aforementioned effects is preferably 0.0001%, and more preferably 0.0005%.
[0047] The NiCrFe alloy according to the present invention may contain one or two types
selected from the group consisting of Ca and Mg in place of part of Fe. Each of these
elements forms a compound with S, thereby assisting the effects of REM.
Ca: 0 to 0.010%
[0048] Calcium (Ca) is an optional element and may not be contained. When contained, Ca
forms a compound with S, thereby assisting the S immobilizing effect of REM. If Ca
is contained in any small amount, the aforementioned effect will be obtained to some
degree. However, when the Ca content is too high, Ca forms oxide, and deteriorates
the hot workability of the alloy. Therefore, the Ca content is 0 to 0.010%. The upper
limit of the Ca content is preferably 0.008%. The lower limit of the Ca content to
effectively obtain the aforementioned effect is preferably 0.0001%, more preferably
0.0002% and further preferably 0.0003%.
Mg: 0 to 0.010%
[0049] Magnesium (Mg) is an optional element and may not be contained. When contained, Mg
forms a compound with S, thereby assisting the S immobilizing effect of REM. When
Mg is contained in any small amount, the aforementioned effect will be obtained to
some degree. However, when the Mg content is too high, Mg forms oxide, thereby deteriorating
the hot workability of the alloy. Therefore, the Mg content is 0 to 0.010%. The upper
limit of the Mg content is preferably 0.008%. The lower limit of the Mg content to
effectively obtain the aforementioned effect is preferably 0.0001%, more preferably
0.0002% and further preferably 0.0003%.
[0050] The NiCrFe alloy according to the present invention may contain one or more types
selected from the group consisting of V, Nb, Ta, and Hf in place of part of Fe. Each
of these elements forms carbide and carbonitride, thereby increasing the creep strength
of the alloy.
V: 0 to 0.5%
[0051] Vanadium (V) is an optional element and may not be contained. When contained, V forms
fine carbide and carbonitride with C and N, thereby increasing the creep strength
of the alloy. When V is contained in any small amount, the aforementioned effect will
be obtained to some degree. However, when the V content is too high, a large amount
of carbide and carbonitride will precipitate, thereby deteriorating the creep ductility
of the alloy. Therefore, the V content is 0 to 0.5%. The upper limit of the V content
is preferably 0.4%. The lower limit of the V content to effectively obtain the aforementioned
effect is 0.01%.
Nb: 0 to 1.0%
[0052] Niobium (Nb) is an optional element and may not be contained. When contained, Nb
forms fine carbide and carbonitride with C and N, thereby increasing the creep strength
of the alloy. When Nb is contained in any small amount, the aforementioned effect
will be obtained to some degree. However, when the Nb content is too high, a large
amount of carbide and carbonitride will precipitate, thereby deteriorating the creep
ductility and toughness of the alloy. Therefore, the Nb content is 0 to 1.0%. The
upper limit of the Nb content is preferably 0.4%. The lower limit of the Nb content
to effectively obtain the aforementioned effect is 0.01%.
Ta: 0 to 1.0%
[0053] Tantalum (Ta) is an optional element and may not be contained. When contained, Ta
forms fine carbide and carbonitride with C and N, thereby increasing the creep strength
of the alloy. When Ta is contained in any small amount, the aforementioned effect
will be obtained to some degree. However, when the Ta content is too high, a large
amount of carbide and carbonitride will precipitate, thereby deteriorating the creep
ductility and toughness of the alloy. Therefore, the Ta content is 0 to 1.0%. The
upper limit of the Ta content is preferably 0.4%. The lower limit of the Ta content
to effectively obtain the aforementioned effect is 0.01%.
Hf: 0 to 1.0%
[0054] Hafnium (Hf) is an optional element and may not be contained. When contained, Hf
forms fine carbide and carbonitride with C and N, thereby increasing the creep strength
of the alloy. When Hf is contained in any small amount, the aforementioned effect
will be obtained to some degree. However, when the Hf content is too high, a large
amount of carbide and carbonitride will precipitate, thereby deteriorating the creep
ductility and toughness of the alloy. Therefore, the Hf content is 0 to 1.0%. The
upper limit of the Hf content is preferably 0.4%. The lower limit of the Hf content
to effectively obtain the aforementioned effect is 0.01%.
[0055] The NiCrFe alloy according to the present invention may contain one or more types
selected from the group consisting of Mo, W, Co, and Cu in place of part of Fe.
Mo: 0 to 1.0%
[0056] Molybdenum (Mo) is an optional element and may not be contained. When contained,
Mo dissolves into the alloy, thereby increasing the creep strength of the alloy at
high temperatures. When Mo is contained in any small amount, such effect will be obtained
to some degree. However, when the Mo content is too high, the stability of austenite
will be lost, thereby deteriorating the toughness of the alloy. Therefore, the Mo
content is 0 to 1.0%. The upper limit of the Mo content is preferably 0.9%. The lower
limit to effectively obtain the aforementioned effect is preferably 0.01%.
W: 0 to 2.0%
[0057] Tungsten (W) is an optional element and may not be contained. When contained, W dissolves
into the alloy, thereby increasing the creep strength of the alloy at high temperatures.
When W is contained in any small amount, such effect will be obtained to some degree.
However, when the W content is too high, the stability of austenite will be lost,
thereby deteriorating the toughness of the alloy. Therefore, the W content is 0 to
2.0%. The upper limit of the W content is preferably 1.8%. The lower limit of the
W content to effectively obtain the aforementioned effect is preferably 0.01%.
Co: 0 to 3.0%
[0058] Cobalt (Co) is an optional element and may not be contained. When contained, Co stabilizes
austenite and dissolves into the alloy, thereby increasing the creep strength of the
alloy at high temperatures. When Co is contained in any small amount, such effects
will be obtained to some degree. However, when the Co content is too high, the production
cost increases. Therefore, the Co content is 0 to 3.0%. The upper limit of the Co
content is preferably 2.8%. The lower limit of the Co content to effectively obtain
the aforementioned effects is preferably 0.01%.
Cu: 0 to 3.0%
[0059] Cupper (Cu) is an optional element and may not be contained. When contained, Cu stabilizes
austenite and suppresses precipitation of brittle phase such as σ phase during use
at high temperatures. When Cu is contained in any small amount, such effects will
be obtained to some degree. However, when the Cu content is too high, the hot workability
of the alloy deteriorates. Therefore, the Cu content is 0 to 3.0%. The upper limit
of the Cu content is preferably 2.5%, and more preferably less than 2.0%. The lower
limit of the Cu content to effectively obtain the aforementioned effects is preferably
0.01%.
[Formula (1)]
[0060] The NiCrFe alloy according to the present invention further satisfies Formula (1):
where, each symbol of element in Formula (1) is substituted by the content (mass%)
of the corresponding element.
[0061] fn1 = Ti + 48Al/27 is an index to indicate the precipitation amount of γ'. fn1 indicates
a total amount of Ti when the amount of Al is converted into the amount of Ti. When
fn1 is less than 0.50, a sufficient precipitation amount of γ' will not be obtained,
so that the NiCrFe alloy cannot obtain excellent creep resistance. On the other hand,
when fn1 is more than 2.20, the stress relaxation cracking resistance, creep ductility,
and toughness of the alloy will deteriorate due to an excessive precipitation amount
of γ'. Therefore, fn1 is 0.50 to 2.20. In this range, an appropriate amount of γ'
is precipitated, and excellent creep resistance is obtained. The upper limit of the
fn1 is preferably 2.00. The lower limit of fn1 is preferably 0.65.
[Formula (2)]
[0062] The above described chemical composition further satisfies Formula (2):
where, each symbol of element in Formula (2) is substituted by the content (mass%)
of the corresponding element.
[0063] fn2 = Ti/(Ti + 48Al/27) is a ratio of the Ti content with respect to the total content
of Al and Ti, the content of Al being converted into the amount of Ti. When fn2 is
less than 0.40, the Ti content is too low with respect to the Al content, and the
precipitation amount of γ' decreases. As a result, the NiCrFe alloy cannot obtain
excellent creep strength. On the other hand, when fn2 is more than 0.80, the Ti content
becomes excessive with respect to the Al content so that although fine γ' precipitates
in an early stage of creep, the γ' changes to coarse and acicular η phase over time.
As a result, the creep strength and toughness of the alloy deteriorate. Therefore,
fn2 is 0.40 to 0.80. In this range, γ' precipitates in an appropriate amount, and
will not change to η phase even when further time passes so that excellent creep strength
is obtained. The upper limit of fn2 is preferably 0.75.
[Formula (3)]
[0064] The above described chemical composition further satisfies Formula (3):
where, each symbol of element in Formula (3) is substituted by the content (mass%)
of the corresponding element, and A(REM) is substituted by the atomic weight of each
REM.
[0065] fn3 = ∑[REM/(A(REM))] - S/32 - 2/3•O/16 is an index to indicate the amount of S that
segregates in grain boundaries. When fn3 is a negative value, S segregates in grain
boundaries, thereby resulting in grain boundary embrittlement so that the stress relaxation
cracking resistance of the alloy deteriorates. On the other hand, when fn3 is not
less than 0, REM immobilizes S as inclusions, thereby decreasing the S content in
the matrix. As a result, it is possible to improve the stress relaxation cracking
resistance of the alloy. Therefore, fn3 is not less than 0.
[Production method]
[0066] One example of production method of the NiCrFe alloy of the present embodiment will
be described. The production method of the present embodiment comprises a process
of producing an ingot (steelmaking process), and a process of producing a hot-rolled
plate (hot working process). Hereinafter, each process will be described in detail.
[Steelmaking process]
[0067] First, alloys having the above described chemical compositions are melted. The melting
is performed by using, for example, the high-frequency induction vacuum melting. Next,
an ingot is produced by an ingot-making method.
[Hot working process]
[0068] In the hot working process, normally, hot working is performed once or multiple times.
First, the ingot is heated, and thereafter hot working is performed. The hot working
refers to, for example, hot forging and hot rolling. The hot working may be performed
by a well-known method.
[0069] Further, the hot-worked NiCrFe alloy may be subjected to cold working. The cold working
is, for example, cold rolling.
[0070] Further, the NiCrFe alloy, which has been subjected to the above described working,
may be subjected to heat treatment. The heat treatment temperature is preferably 1050
to 1200°C. Further, after being heated and held, the NiCrFe alloy is preferably water
cooled.
[0071] In the above described exemplary production method, a production method of a NiCrFe
alloy plate has been described. However, the NiCrFe alloy may be a bar or an alloy
pipe. In other words, the shape of the product will not be limited. Moreover, in the
case of the alloy pipe, it is preferable that hot working by hot extrusion is performed.
[0072] The NiCrFe alloy produced by the processes described so far has excellent creep strength
and excellent stress relaxation cracking resistance.
[Micro structure]
[0073] In the NiCrFe alloy according to the present invention, γ' and η phase precipitate
in a use environment at high temperatures. In other words, the microstructure of the
NiCrFe alloy according to the present invention after being kept at 650°C for 3000
hours contains a total of 2 to 6 mass% of γ' and η phase, wherein the number density
of η phase is less than 5/100 µm
2. Note that the γ' and η phase are herein also collectively referred to as "aging
precipitates".
[0074] In a case where the NiCrFe alloy according to the present invention is subjected
to aging treatment for keeping the alloy at 650°C for 3000 hours and then a total
of γ' and η phase is less than 2 mass%, the precipitation amount of γ' in the alloy
will be decreased. As a result, the NiCrFe alloy cannot obtain excellent creep strength.
On the other hand, in a case where the same aging treatment is performed and then
a total of γ' and η phase is more than 6 mass%, the precipitation amount of γ' may
excessively increase. In that case, the alloy cannot obtain excellent stress relaxation
cracking resistance. Therefore, the total of γ' and η phase after aging treatment
is 2 to 6 mass%.
[0075] Specifically, the total of γ' and η phase can be measured by the following method.
The NiCrFe alloy according to the present invention is subjected to aging treatment
for keeping the alloy at 650°C for 3000 hours. A test specimen of 10 mm × 5 mm × 50
mm is sampled from the NiCrFe alloy after the aging treatment. When the alloy is an
alloy plate, the test specimen is sampled from a middle part of plate thickness of
the alloy pipe. On the other hand, when the alloy is an alloy pipe, the test specimen
is sampled from a middle part of wall thickness. Note that the weight of the test
specimen is measured in advance.
[0076] The sampled test specimen is electrolyzed in a 1% tartaric acid-1% (NH
4)
2SO
4-water solution to sample the residue from the electrolyte. The sampled residue is
melted by HCl (1+4)-20% tartaric acid solution of 60°C and the solution is filtered.
The filtrate is measured by ICP emission spectrometry to determine Ti, Al, and Ni
concentrations in the residue. From the determined Ti, Al, and Ni concentrations in
the residue, and the weight of the test specimen, Ti, Al, and Ni contents in the γ'
and η phase of the test specimen are determined. The sum of Ti, Al, and Ni contents,
which have been determined by the method described so far, is defined as a sum of
γ' and η phase (mass%).
[0077] In a case where the NiCrFe alloy according to the present invention is subjected
to aging treatment for keeping the alloy at 650°C for 3000 hours and then the number
density of η phase is not less than 5/100 µm
2, part of γ' has changed to η phase. For that reason, the NiCrFe alloy cannot obtain
excellent creep strength. Therefore, the number density of η phase after aging treatment
is less than 5/100 µm
2.
[0078] Specifically, the number density of η phase can be measured by the following method.
The NiCrFe alloy according to the present invention is subjected to aging treatment
for keeping the alloy at 650°C for 3000 hours. Microscopic observation is performed
on the NiCrFe alloy after aging treatment. Specifically, a microscopic test specimen
is sampled from the NiCrFe alloy after aging treatment. When the alloy is an alloy
plate, the test specimen is sampled from a middle part of the plate thickness. On
the other hand, when the alloy is an alloy pipe, the microscopic test specimen is
sampled from a middle part of wall thickness of the alloy pipe. The sampled microscopic
test specimen is subjected to mechanical polishing. The surface of the microscopic
test specimen after mechanical polishing is electrolytically corroded by 10% oxalic
acid. The microscopic test specimen after electrolytic corrosion is observed by a
scanning electron microscope (SEM) in 5 visual fields, and an SEM image is created
for each visual field. The observation magnification is 10000 times, and observation
field is, for example, 12 µm × 9 µm.
[0079] The γ' and η phase differ in their shapes. Specifically, γ' is observed to be spherical
and η phase be acicular. More specifically, an aspect ratio of γ' is less than 3,
and an aspect ratio of η phase is not less than 3. Here, the term "aspect ratio" means
a value obtained by dividing the major axis length by the minor axis length for each
aging precipitate.
[0080] In the above described SEM image of each visual field, aging precipitates (γ' and
η phase) are identified from contrast. Further, by image processing, aspect ratios
are calculated for the identified aging precipitates. To calculate an aspect ratios,
general purpose application software may be used. When a calculated aspect ratio is
not less than 3, the aging precipitate is identified to be η phase.
[0081] For an SEM image of each visual field, the number of identified η phase is counted
to determine a sum of the numbers in all visual fields. By using the number of η phase
in all visual fields and the area of the all visual fields, the number density of
η phase in an observation field of 100 µm
2 (number/100 µm
2) is determined.
EXAMPLES
[0082] Alloys having chemical compositions indicated by Reference marks 1 to 15 shown in
Table 1 were melted by the high-frequency induction vacuum melting method.
[Table 1]
[0083]
TABLE 1
Reference mark |
Chemical composition (in mass%, with the balance being Fe and impurities) |
fn1 |
fn2 |
fn3 |
REM element |
C |
Si |
Mn |
P |
S |
Cr |
Ni |
Ti |
Al |
N |
O |
REM |
Others |
1 |
0.06 |
0.41 |
0.85 |
0.013 |
0.0020 |
20.4 |
34.8 |
0.85 |
0.37 |
0.008 |
0.003 |
0.028 |
- |
1.51 |
0.56 |
0.0000069 |
Nd |
2 |
0.07 |
0.38 |
1.04 |
0.011 |
0.0006 |
23.5 |
35.0 |
1.35 |
0.36 |
0.012 |
0.002 |
0.015 |
B:0.002,Mo:0.5 |
1.99 |
0.68 |
0.0000021 |
Nd |
3 |
0.07 |
0.25 |
0.98 |
0.014 |
0.0010 |
19.8 |
31.4 |
0.95 |
0.43 |
0.007 |
0.005 |
0.035 |
Ca:0.003 |
1.71 |
0.55 |
0.0000035 |
Nd |
4 |
0.08 |
0.29 |
1.48 |
0.008 |
0.0003 |
19.3 |
29.5 |
0.84 |
0.55 |
0.006 |
0.002 |
0.019 |
W:1.2 |
1.82 |
0.46 |
0.0000392 |
Nd |
5 |
0.07 |
0.25 |
0.78 |
0.011 |
0.0007 |
21.7 |
30.7 |
1.19 |
0.43 |
0.009 |
0.005 |
0.044 |
Mg:0.004, V:0.3 |
1.95 |
0.61 |
0.0000753 |
Nd |
6 |
0.05 |
0.36 |
0.89 |
0.010 |
0.0006 |
23.6 |
33.0 |
0.64 |
0.35 |
0.009 |
0.004 |
0.031 |
Nb:0.3,Co:2.3 |
1.26 |
0.51 |
0.0000376 |
La |
7 |
0.05 |
0.38 |
0.95 |
0.012 |
0.0010 |
20.5 |
27.3 |
0.72 |
0.45 |
0.003 |
0.003 |
0.030 |
Ta:0.1,Hf:0.4 |
1.52 |
0.47 |
0.0000580 |
Ce |
8 |
0.08 |
0.22 |
0.99 |
0.015 |
0.0010 |
22.1 |
32.9 |
0.77 |
0.65 |
0.004 |
0.003 |
0.018 |
Cu:1.3,Co:2.1 |
1.93 |
0.40 |
0.0000460 |
Y |
9 |
0.08 |
0.55 |
1.45 |
0.013 |
0.0010 |
21.8 |
31.6 |
0.28 |
0.11 |
0.013 |
0.001 |
0.028 |
Co:2.5 |
0.48 |
0.59 |
0.0001215 |
Nd |
10 |
0.09 |
0.21 |
0.64 |
0.016 |
0.0010 |
21.8 |
29.8 |
1.33 |
0.70 |
0.008 |
0.002 |
0.032 |
Nb:0.3 |
2.57 |
0.52 |
0.0001076 |
Nd |
11 |
0.10 |
0.25 |
1.12 |
0.008 |
0.0008 |
21.3 |
34.5 |
0.45 |
0.84 |
0.008 |
0.005 |
0.042 |
- |
1.94 |
0.23 |
0.0000583 |
Nd |
12 |
0.08 |
0.18 |
1.51 |
0.009 |
0.0004 |
19.5 |
29.9 |
0.65 |
0.57 |
0.008 |
0.002 |
0.024 |
W:1.5 |
1.66 |
0.39 |
0.0000708 |
Nd |
13 |
0.10 |
0.33 |
0.88 |
0.011 |
0.0005 |
20.9 |
38.2 |
1.67 |
0.11 |
0.011 |
0.002 |
0.065 |
B:0.002,Cu:1.3 |
1.87 |
0.90 |
0.0003524 |
Nd |
14 |
0.10 |
0.25 |
1.07 |
0.009 |
0.0010 |
19.9 |
32.8 |
0.88 |
0.41 |
0.009 |
0.005 |
0.026 |
B:0.003 |
1.61 |
0.55 |
-0.0000590 |
Nd |
15 |
0.07 |
0.42 |
0.88 |
0.013 |
0.0010 |
20.8 |
34.5 |
0.88 |
0.36 |
0.007 |
0.003 |
- |
- |
1.52 |
0.58 |
-0.0001563 |
- |
[0084] An ingot of 50 kg was produced by using an alloy of each Reference mark. The ingot
was subjected to hot forging and hot rolling to obtain a plate material having a thickness
of 15 mm. Each plate material was kept at 1150°C for 30 minutes, and thereafter the
plate material was rapidly cooled (water cooling) and subjected to solution treatment.
By the production processes described so far, NiCrFe alloy plate materials were produced.
Using thus produced NiCrFe alloy plate materials, the following tests were conducted.
[Creep rupture test]
[0085] A test specimen was fabricated from the produced alloy plate material. The test specimen
was sampled from a central part of thickness of the alloy plate material in parallel
with the longitudinal direction (rolling direction). The specimen was a round bar
test specimen, whose parallel part had a diameter of 6 mm, and which had a gauge length
of 30 mm. By using the test specimen, a creep rupture test was conducted. The creep
rupture test was performed under a tensile load of 70 MPa in the air atmosphere of
750°C. A test specimen whose rupture time was not less than 3000 hours was evaluated
as "E" (Excellent), and those whose rupture time was less than 3000 hours as "NA"
(Not Acceptable).
[Table 2]
[0086]
TABLE 2
Reference mark |
Microstructure evaluation |
Sum of γ' and η phase (mass%) |
Creep rupture test |
Stress relaxation cracking test |
1 |
E |
3.0 |
E |
E |
2 |
E |
5.2 |
E |
E |
3 |
E |
3.7 |
E |
E |
4 |
E |
3.3 |
E |
E |
5 |
E |
5.0 |
E |
E |
6 |
E |
3.4 |
E |
E |
7 |
E |
2.9 |
E |
E |
8 |
E |
2.8 |
E |
E |
9 |
L |
0.0 |
NA |
E |
10 |
TM |
7.6 |
E |
NA |
11 |
L |
0.8 |
NA |
E |
12 |
L |
1.9 |
NA |
E |
13 |
TM, η |
7.0 |
NA |
E |
14 |
E |
3.6 |
E |
NA |
15 |
E |
3.8 |
E |
NA |
[Microstructure observation]
[0087] From thus produced alloy plate materials, test specimens were fabricated by the above
described method. The fabricated test specimens were subjected to aging treatment
to keep them at 650°C for 3000 hours, and the sum (mass%) of γ' and η phase of each
test specimen was determined by the above described method. Further, the number density
of η phase (number/100 µm
2) was determined by the above described method. A sum of γ' and η phase of less than
2 mass% was evaluated as "L" (Less), that of 2 to 6 mass% as "E" (Excellent), and
that of more than 6 mass% as "TM" (Too Much). Further, those showed a number density
of η phase of not less than 5/100 µm
2 were evaluated as "η".
[Stress relaxation cracking test]
[0088] The produced alloy plate material was further subjected to cold working. Specifically,
cold rolling was performed on the alloy plate material until its thickness became
12 mm. The reduction of area of this cold rolling was 20%. A test specimen was fabricated
from this alloy plate material. The test specimen was sampled from a central part
of thickness of the alloy plate material in parallel with the longitudinal direction
(rolling direction). The specimen was a round bar test specimen, whose parallel part
had a diameter of 6 mm, and which had a gauge length of 30 mm. By using the specimen,
a stress relaxation cracking test was conducted. The stress relaxation cracking test
was conducted such that the test specimen is subjected to tensile strain 10% at a
strain rate of 0.05 min
-1 and is kept as is for 300 hours in air atmosphere of 650°C. A specimen which did
not rupture after being kept for 300 hours was evaluated as "E" (Excellent), and one
which ruptured as "NA" (Not Acceptable).
[Test results]
[0089] Test results are shown in Table 2.
[0090] Referring to Table 2, the chemical compositions of Reference marks 1 to 8 were appropriate,
so that fn1 was 0.50 to 2.20, fn2 was 0.40 to 0.80, and fn3 was not less than 0. For
that reason, in the microstructure, the sum of γ' and η phase was 2 to 6 mass%. Further,
the number density of η phase was less than 5/100 µm
2. As a result, the creep rupture time was not less than 3000 hours, thus exhibiting
excellent creep strength. Further, none of the specimens ruptured in the stress relaxation
cracking test, exhibiting excellent stress relaxation cracking resistance.
[0091] On the other hand, in Reference mark 9, the value of fn1 was too low. For that reason,
in the microstructure, the sum of γ' and η phase was less than 2 mass%, which was
too low. As a result, the creep rupture time was less than 3000 hours, not exhibiting
excellent creep strength.
[0092] In Reference mark 10, the value of fn1 was too high. For that reason, in the microstructure,
the sum of γ' and η phase was more than 6 mass%. Further, the number density of η
phase was less than 5/100 µm
2. In other words, in the microstructure, γ' was more than 6 mass%, which was too high.
As a result, the test specimen ruptured in the stress relaxation cracking test, thus
not exhibiting excellent stress relaxation cracking resistance.
[0093] In Reference marks 11 and 12, the value of fn2 was too low. For that reason, in the
microstructure, the sum of γ' and η phase was less than 2 mass%, which was too low.
As a result, the creep rupture time was less than 3000 hours, not exhibiting excellent
creep strength.
[0094] In Reference mark 13, the value of fn2 was too high. For that reason, in the microstructure,
the number density of η phase was not less than 5/100 µm
2. As a result, the creep rupture time was less than 3000 hours, not exhibiting excellent
creep strength.
[0095] In Reference mark 14, the value of fn3 was too low. As a result, the specimen ruptured
in the stress relaxation cracking test, thus not exhibiting excellent stress relaxation
cracking resistance. This is considered because S in the matrix could not be immobilized.
[0096] In Reference mark 15, the REM content was too low. Further the value of fn3 was too
low. As a result, the test specimen ruptured in the stress relaxation cracking test,
not exhibiting excellent stress relaxation cracking resistance. This is considered
because S in the matrix could not be immobilized.
[0097] So far, the embodiments of the present invention have been described. However, the
above described embodiments are merely illustration for practicing the present invention.
Therefore, the present invention will not be limited to the above described embodiments,
and can be practiced by appropriately modifying the above described embodiments within
a range not departing from the spirit of the invention.
INDUSTRIAL APPLICABILITY
[0098] The present invention can be widely applied to uses for which high creep strength
and stress relaxation cracking resistance are demanded. Particularly, the present
invention can be suitably used for high temperature members of thermal power generation
boilers, petroleum refining and chemical industry plants, or the like.