FIELD
[0001] Embodiments described herein relate generally to an Ni-based alloy for forging, a
method for manufacturing the same, and a turbine component.
BACKGROUND
[0002] A thermal power generation plant has been made highly efficient in recent years in
view of reduction of carbon dioxide emissions into the atmosphere. Thus, it is required
to make a steam turbine and a gas turbine which are installed in the thermal power
generation plant highly efficient. Further, it is also required to make a CO
2 turbine which is installable in the thermal power generation plant highly efficient.
Here, the CO
2 turbine is one which drives a turbine, using CO
2 generated by combustion of fuel such as natural gas and oxygen as a working fluid.
In the CO
2 turbine, great part of generated CO
2 can be easily separated and recovered. Thus, the CO
2 turbine attracts attention in view of global environmental protection.
[0003] In order to raise an efficiency in each turbine described above, it is effective
to raise an inlet temperature of a working fluid introduced into the turbine. For
example, with regard to the steam turbine, an operation at a temperature of 700°C
or higher of steam being the working fluid is expected in the future. With regard
to the gas turbine or the CO
2 turbine, the inlet temperatures of the introduced working fluids tend to be higher.
[0004] Thus, a component constructing a hot section of each turbine is desirable to be constituted
with an Ni-based alloy, which is used for a component of a power generation gas turbine
or an air plane engine, and is time-proven in usage in a high-temperature place.
[0005] As a representative example of the Ni-based alloy, there can be cited Inconel 718
and Inconel 617 (manufactured by Special Metals Corporation). A strengthening mechanism
of the Ni-based alloy can be broadly classified into a precipitation strengthening
type and a solid solution strengthening type.
[0006] In the precipitation strengthening type Ni-based alloy, as a result that a precipitates
called as a γ' (gamma prime: Ni
3(Al, Ti)) phase or a γ" (gamma double prime: Ni
3Nb) phase is precipitated by adding Al, Ti, Ta, and Nb to Ni, a mechanical strength
under a high temperature is improved. As the representative precipitation strengthening
type Ni-based alloy, the above-described Inconel 718 can be cited.
[0007] On the other hand, in the solid solution strengthening type Ni-based alloy, a matrix
itself is strengthened by adding Co, Mo, and the like to Ni. As the representative
solid solution strengthening type Ni-based alloy, the above-described Inconel 617
can be cited.
[0008] As described above, as a material for a component of a turbine used under a high-temperature
environment, application of an Ni-based alloy is discussed. As for the Ni-based alloy,
a sufficient mechanical strength under the high-temperature environment is required,
and further, productivity in manufacturing of a large-sized forged component and the
like is required.
[0009] EP 2 233 594 A1 discloses a nickel (Ni)-base alloy for a turbine rotor of a steam turbine containing,
in mass%, carbon (C): 0.01% to 0.15%, chromium (Cr): 18% to 28%, cobalt (Co): 10%
to 15%, molybdenum (Mo): 8% to 12%, aluminum (Al): 0.5% to less than 1.5%, titanium
(Ti): 0.7% to 3.0%, and boron (B): 0.001% to 0.006%, the balance being nickel (Ni)
and unavoidable impurities.
BRIEF DESCRIPTION OF THE DRAWINGS
[0010]
Fig. 1 is a diagram schematically showing a microstructure of an Ni-based alloy in
an embodiment;
Fig. 2 is a view showing an electron micrograph of an Ni-based alloy, in order to
explain a precipitation form of a carbide precipitated into a grain boundary by a
condition of an aging treatment; and
Fig. 3 is a view showing an electron micrograph of an Ni-based alloy, in order to
explain a precipitation form of a carbide precipitated into a grain boundary by a
condition of the aging treatment.
DETAILED DESCRIPTION
[0011] In one embodiment, a nickel (Ni)-based alloy for forging contains, in mass%, 0.01
to 0.07% of carbon (C), 14 to 26% of chromium (Cr), 10 to 15% of cobalt (Co), 5 to
12% of molybdenum (Mo), 0.8 to 3% of aluminum (Al), 0.8 to 3% of titanium (Ti), and
0.001 to 0.006% of boron (B), optionally 0.05 to 0.7% of tantalum (Ta), optionally
0.1 to 0.7% of niobium (Nb), and the balance being nickel (Ni) and unavoidable impurities,
the relationship 11 mass% ≤ Mo + 0.176Cr + 0.037Co≤13.5 mass% being satisfied, wherein
an average thickness of a carbide precipitated along a grain boundary is 250 nm or
less.
[0012] Hereinafter, an embodiment according to the present invention will be described.
[0013] In the Ni-based alloy, a material strength at a room temperature and a high temperature
is improved by solid solution strengthening by a solid solution strengthening element
such as Mo and W and by precipitation strengthening by fine precipitation of a γ'
(gamma prime: Ni
3(Al, Ti)) phase obtained by addition of Al, Ti and the like. On the other hand, excessive
strengthening deteriorates hot workability of a material and reduces productivity.
[0014] For example, Inconel 617 in which a precipitation strengthening amount by a γ' phase
is minor has better forgeability compared with Udimet 520 (manufactured by Special
Metals Corporation) or the like in which a precipitation strengthening amount by a
γ' phase is large. On the other hand, Inconel 738LC (manufactured by Special Metals
Corporation) in which a precipitation amount of a γ' phase is large is not able to
be formed by forging and is generally formed by casting.
[0015] As described above, a method for manufacturing an Ni-based alloy is determined mainly
by a precipitation amount of a γ' phase. For example, in a case of an Ni-based alloy
for forging, an alloy composition is set in a manner not to generate excessive precipitation
of a γ' phase during forging process.
[0016] A large-sized forging such as a turbine rotor of a steam turbine or a CO
2 turbine installed in a thermal power generation plant is larger compared with a forging
such as a gas turbine or a jet engine for which an Ni-based alloy is conventionally
used. Thus, in order to manufacture such a large-sized forging, a forging of over
10 tons is necessary, for example.
[0017] In forging such a large-sized forging, there is a case where a good forged product
is not obtained even with Inconel 617 or the like, which has been conventionally considered
forgeable, due to a cause such as capacity lack of a forging press. As described above,
with regard to an Ni-based alloy used for a large-sized forging, it is necessary to
consider not only a precipitation amount of a γ' phase but also a solid solution strengthening
amount which influences a deformation resistance at a high temperature.
[0019] Further, as a factor to influence a characteristic of a metal material substantially,
a microstructure of a material can be cited. In an Ni-based alloy, a characteristic
of a material is dependent on structures of not only within a grain but also in a
grain boundary. In particular, it is known that an M
23C
6 type carbide precipitated on a grain boundary in a film shape reduces a toughness
of a material. Therefore, in order to secure reliability of a material, it is necessary
to control a microstructure properly by optimizing a heat treatment condition.
[0020] Under the circumstances, the present inventors find a parameter indicating a solid
solution strengthening amount by quantitatively evaluating a solid solution amount
of each additive element and a misfit strain which influence an Ni-based alloy. Further,
the present inventors carry out various material testings about materials whose chemical
compositions are changed, and find a chemical composition which has excellent forgeability
while maintaining a sufficient material strength.
[0021] Further, as a result of investigation of grain boundary structures of Ni-based alloys
to which various heat treatments are applied, the present inventors find "an average
thickness of a carbide on a grain boundary" as a factor dominating a toughness of
an Ni-based alloy. Besides, the present inventors clarify a range of the carbide thickness
on the grain boundary by which the toughness can be secured.
[0022] Next, an Ni-based alloy for forging of an embodiment will be described concretely.
[0023] Fig. 1 is a diagram schematically showing a microstructure of an Ni-based alloy in
the embodiment. Note that "%" indicating a composition component in the following
explanation means "mass%" if not specifically mentioned.
[0024] The Ni-based alloy of the embodiment contains 0.01 to 0.07% of carbon (C), 14 to
26% of chromium (Cr), 10 to 15% of cobalt (Co), 5 to 12% of molybdenum (Mo), 0.8 to
3% of aluminum (Al), 0.8 to 3% of titanium (Ti), and 0.001 to 0.006% of boron (B),
optionally 0.05 to 0.7% of tantalum (Ta), optionally 0.1 to 0.7% of niobium (Nb),
and the balance being nickel (Ni) and unavoidable impurities, the relationship 11
mass% ≤ Mo + 0.176Cr + 0.037Co ≤13.5 mass% being satisfied.
[0025] Further, in the Ni-based alloy of the embodiment, as shown in Fig. 1, a carbide 11
is precipitated along a grain boundary 10. An average thickness t of the carbide 11
is preferable to be 250 nm or less. The carbide 11 is continuously precipitated along
the grain boundary 10. Further, in a grain 12, a precipitate 13 is precipitated in
a grain shape.
[0026] The carbide 11 is a carbide whose major constituents are Cr and Mo, and concretely,
is an M
23C
6 type carbide. The reason why the average thickness t of the carbide 11 is preferable
to be 250 nm or less is that, for example, such a thickness not reducing a toughness,
a toughness for manufacturing a turbine component properly can be secured.
[0027] The precipitate 13 is constituted with a γ' (gamma prime: Ni
3(Al, Ti)) phase. A diameter of the γ' phase is preferable to be small, in view of
precipitation strengthening. An average diameter of the γ' phase is preferable to
be 150 nm or less, for example.
[0028] Here, the Ni-based alloy in the embodiment can contain 0.05 to 0.7% of Ta, in addition
to the aforementioned chemical composition. The Ni-based alloy in the embodiment can
contain 0.1 to 0.7% of Nb. The Ni-based alloy in the embodiment can contain 0.05 to
0.7% of Ta and 0.1 to 0.7% of Nb.
[0029] Note that as the unavoidable impurity there can be cited Si, Mn, N, Cu, Fe, S, and
the like, for example. A remaining content ratio of the unavoidable impurity as above
is preferable to be approximated to 0% to the extent possible.
[0030] The above-described Ni-based alloy of the embodiment is suitable as a material constituting
a turbine component constructed by forging, such as a power generation turbine, for
example, which is used under a temperature of 650°C or higher, for example. As the
turbine component, there can be cited a turbine rotor, a rotor blade, a stationary
blade, a screwing member, a pipe, and the like, for example. Those forged components
are each disposed in a high temperature and high pressure environment.
[0031] Here, as the screwing member, there can be exemplified a bolt and a nut fixing a
turbine casing or various component parts inside the turbine, for example. Further,
as the pipe, there can be exemplified a pipe which is disposed in a power generation
turbine plant or the like and through which a high temperature and high pressure working
fluid passes, for example.
[0032] Note that it is possible to construct all portions of the turbine components of the
power generation turbine described above with the above-described Ni-based alloy.
Further, it is also possible to construct a limited portion of the turbine component
to be subjected to a high temperature in particular with the above-described Ni-based
alloy.
[0033] The Ni-based alloy for forging of the embodiment described above is superior to a
conventional Ni-based alloy for forging in a strength characteristic and is superior
in forgeability. Thus, the turbine components fabricated by using the Ni-based alloy
for forging of the embodiment, such as a turbine rotor, a rotor blade, a stationary
blade, a screwing member, and a pipe, have high reliability even under a high temperature
environment.
[0034] Next, there will be explained a reason for limiting a range of each composition component
in the Ni-based alloy for forging of the embodiment described above.
(1) C (carbon)
[0035] C is effective as a constituent element of a carbide being a strengthening phase.
Further, C has a function to suppress coarsening of a grain under a high temperature,
by a pinning effect of the carbide to prevent movement of a grain boundary. When a
content ratio of C is less than 0.01%, strengthening by the carbide is not sufficient.
Further, when the content ratio of C is less than 0.01%, there is a possibility that
failure in securing of a sufficient precipitation amount of the carbide causes coarsening
of the grain. On the other hand, when the content ratio of C is over 0.07%, forgeability
is reduced. Thus, the content ratio of C is set to be 0.01 to 0.07%. Further, the
more preferable content ratio of C is 0.03 to 0.07%.
(2) Cr (chromium)
[0036] Cr is an element indispensable for heightening an oxidation resistance, a corrosive
resistance, and a high temperature strength characteristic of an Ni-based alloy. When
a content ratio of Cr is less than 14%, the oxidation resistance and the corrosive
resistance are reduced. On the other hand, when the content ration of Cr is over 26%,
precipitation of a σ phase which causes reduction of a creep strength becomes prominent
and forgeability is deteriorated. Thus, the content ratio of Cr is set to be 14 to
26%. Further, the more preferable content ratio of Cr is 16 to 20%.
(3) Co (cobalt)
[0037] Co solid-dissolves in a matrix in an Ni-based alloy and improves a creep strength
and a tensile strength. When a content ratio of Co is less than 10%, a sufficient
mechanical strength cannot be obtained. On the other hand, when the content ratio
of Co is over 15%, forgeability is reduced. Thus, the content ratio of Co is set to
be 10 to 15%. Further, the more preferable content ratio of Co is 11 to 14%.
(4) Mo (molybdenum)
[0038] Mo solid-dissolves in an Ni matrix and improves a creep strength and a tensile strength.
Further, by part of Mo substituting in an M
23C
6 type carbide, stability of the carbide is heightened. When a content ratio of Mo
is over 12%, hot workability is reduced. On the other hand, when the content ratio
of Mo is less than 5%, improvement of a mechanical strength cannot be obtained. Thus,
the content ratio of Mo is set to be 5 to 12%. Further, the more preferable content
ratio of Mo is 7 to 10%.
(5) A1 (aluminum)
[0039] A1 generates a γ' (Ni
3Al) phase with Ni and improves a mechanical strength of an Ni-based alloy by precipitation.
When a content ratio of A1 is less than 0.8%, an effect by precipitation of the γ'
phase is not exhibited. On the other hand, when the content ratio of A1 is over 3%,
precipitation of a σ phase is promoted, and the mechanical strength is reduced. Further,
when the content ratio of A1 is over 3%, hot workability is substantially reduced.
Thus, the content ratio of A1 is set to be 0.8% to 3%. Further, the more preferable
content ratio of A1 is 1 to 2%.
(6) Ti (titanium)
[0040] Ti, similarly to A1, generates a γ' (Ni
3(Al, Ti)) phase with Ni, and improves a mechanical strength of an Ni-based alloy.
When a content ratio of Ti is less than 0.8%., an effect by precipitation of the γ'
phase is not exhibited. On the other hand, when the content ratio of Ti is over 3%,
precipitation of a σ phase or a η phase is promoted, and the mechanical strength is
reduced and hot workability is reduced. Thus, the content ratio of Ti is set to be
0.8 to 3%. Further, the more preferable content ratio of Ti is 1 to 2%.
(7) B (boron)
[0041] B segregates in a grain boundary and improves a high temperature strength characteristic.
When a content ratio of B is less than 0.001%, such an effect to improve the high
temperature strength characteristic is not exhibited. On the other hand, when the
content ratio of B is over 0.006%, intergranular embrittlement occurs. Thus, the content
ratio of B is set to be 0.001 to 0.006%. Further, the more preferable content ratio
of B is 0.002 to 0.004%.
(8) Ta (tantalum)
[0042] Ta solid-dissolves in a γ' (Ni
3(Al,Ti)) phase and stabilizes the γ' phase. When a content ratio of Ta is less than
0.05%, the above-described effect is not exhibited. On the other hand, when the content
ratio of Ta is over 0.7%, forgeability is reduced. Thus, the content ratio of Ta is
set to be 0.05% to 0.7%. Further, the more preferable content ratio of Ta is 0.08
to 0.12%.
(9) Nb (niobium)
[0043] Nb, similarly to Ta, solid-dissolves in a γ' (Ni
3(Al,Ti)) phase and stabilizes the γ' phase. When a content ratio of Nb is less than
0.1%, the above-described effect is not exhibited. On the other hand, when the content
ratio of Nb is over 0.7%, segregation occurs at a time of dissolving or forging, and
forgeability is reduced. Thus, the content ratio of Nb is set to be 0.1 to 0.7%. Further,
the more preferable content ratio ofNb is 0.2 to 0.5%.
(10) Mo + 0.176Cr + 0.037Co
[0044] As stated above, it is considered that a solid solution strengthening amount in a
highly concentrated solid solution is proportional to the two-thirds power of a solute
atom concentration and is proportional to the four-thirds power of a misfit strain
due to an atom size difference. Thus, in the present embodiment, for Mo, Cr, and Co,
which are considered to contribute to solid solution strengthening, a parameter representing
solid solution strengthening is defined from the number of atoms per one mass% and
each atomic radius. Note that since the content ratio of C (carbon) is small in the
present embodiment, C is excluded from the parameter.
[0045] Atomic weights of Mo, Cr, and Co are 95.9, 52.0, and 58.9, respectively. A ratio
of the number of atoms in a case where the same amounts of the respective elements
are added is, in a case of Mo being 1, Cr and Co being 1.84 and 1.62, respectively.
Values of the two-thirds power of this ratio are 1, 1.50, and 1.38, respectively.
[0046] Further, a misfit strain occurring at a time that each element is added is determined
by an atomic size difference from an Ni atom. Atomic radius differences between the
Ni atom and Mo, Cr, Co atoms is 0.15 A (angstrom), 0.03 A, and 0.01 A, respectively.
Thus, a ratio of the misfit strain amounts in a case where the respective elements
are added is, in a case of Mo being 1, Cr and Co being 0.200 and 0.067, respectively.
Values of the four-thirds power of this ratio are 1, 0.117, and 0.027, respectively.
[0047] Therefore, a ratio of solid solution strengthening amount per one mass% of the respective
elements is, in a case of Mo being 1, Cr being 0.176 (1.50 × 0.117 = 0.176), and Co
being 0.037 (1.38 × 0.027 = 0.037). From those results, as a parameter representing
a solid solution strengthening amount, "Mo + 0.0176Cr + 0.037Co" is set.
[0048] When a value (content ratio) of this parameter is over 13.5%, the solid solution
strengthening amount becomes excessive, and deteriorates ductility during forging
process. On the other hand, when the value of the parameter is less than 11%, the
solid solution strengthening amount becomes substantially low, and a sufficient strength
cannot be obtained. Thus, the value of the above-described parameter is set to be
11 to 13.5%.
[0049] Note that a misfit strain by addition of an element is considered, in a strict sense,
to be influenced not only by an atomic size but also by interaction or the like with
Ni and other atoms. However, here, for the sake of simplicity, a misfit strain value
is determined from a difference between each solute atom and an Ni atom. Further,
though it is known that Mo and Cr in combination with C form carbides, consumption
of Mo and Cr by the carbide is ignored since the content ratio of C is low.
(11) Si (silicon), Mn (manganese), N (nitrogen), Cu (copper), Fe (iron), and S (sulfur)
[0050] Si, Mn, N, Cu, Fe, and S are classified into unavoidable impurities in the Ni-based
alloy for forging of the embodiment. Remaining content ratios of those unavoidable
impurities are desirable to be approximated to 0% to the extent possible. Further,
it is preferable that among those unavoidable impurities, at least Si and Mn are restricted
to be 0.1% or less and that N is restricted to be 0.01% or less.
[0051] Si is added in order to supplement a corrosive resistance in a case of ordinary steel.
However, an Ni-based alloy has a large Cr content and a sufficient corrosive resistance
can be secured. Thus, it is desirable that a remaining content ratio of Si is set
to be 0.1% or less and that the remaining content ratio thereof is approximated to
0% to the extent possible.
[0052] Mn makes S (sulfur) causing brittleness into MnS, to prevent brittleness, in a case
of ordinary steel. However, a content of S in an Ni-based alloy is quite low and it
is not necessary to add Mn. Thus, it is desirable that a remaining content ratio of
Mn is set to be 0.1% or less and that the remaining content ratio thereof is approximated
to 0% to the extent possible.
[0053] N forms TiN by reacting with Ti in a material and decreases Ti which contributes
to generation of a γ' phase. Consequently, a mechanical strength is reduced. Thus,
it is desirable that a remaining content ratio of N is set to be 0.01% or less and
that the remaining content ratio thereof is approximated to 0% to the extent possible.
[0054] Here, there will be described a method for manufacturing an Ni-based alloy for forging
of the embodiment and a turbine component manufactured by using the Ni-based alloy
for forging.
[0055] The above-described Ni-based alloy for forging of the embodiment is manufactured
as follows, for example.
[0056] First, composition components to constitute the Ni-based alloy are vacuum induction
melted (VIM) and molten metal thereof is poured into a predetermined mold form, to
form an ingot. Then, the ingot is soaking treated and hot forged, and subjected to
a solution treatment, an ageing treatment, and the like, so that the Ni-based alloy
for forging is fabricated.
[0057] A turbine rotor being a turbine component is fabricated as below, for example.
[0058] For example, as one method (double melt), composition components to constitute the
Ni-based alloy for forging of the embodiment are vacuum induction melted (VIM), electroslag
remelted (ESR), and poured into a predetermined mold form. Subsequently, a soaking
treatment, a forging treatment, a solution treatment, an aging treatment, and the
like are carried out, so that the turbine rotor is fabricated.
[0059] As another method (double melt), composition components to constitute the Ni-based
alloy for forging of the embodiment are vacuum induction melted (VIM), vacuum arc
remelted (VAR), and poured into a predetermined mold form. Subsequently, a soaking
treatment, a forging treatment, a solution treatment, an aging treatment, and the
like are carried out, so that the turbine rotor is fabricated.
[0060] Further, as still another method (triple melt), composition components to constituted
the Ni-based alloy for forging of the embodiment are vacuum induction melted (VIM),
electroslag remelted (ESR), vacuum arc remelted (VAR), and poured into a predetermined
mold form. Subsequently, a soaking treatment, a forging treatment, a solution treatment,
an aging treatment, and the like are carried out, so that the turbine rotor is fabricated.
[0061] At least a predetermined portion of the turbine rotor is manufactured by the above-described
method for manufacturing the turbine rotor. As the predetermined portion, there can
be cited a portion exposed to a high temperature of 700°C or higher, for example,
of the turbine rotor. In this case, a portion exposed to a temperature of about 600°C,
for example, of the turbine rotor is manufactured with a conventional heat resistant
alloy. Then, a component made of the Ni-based alloy for forging of the embodiment
manufactured by the above-described manufacturing method and a component made of the
conventional heat resistant alloy are joined by welding, for example, to construct
a turbine rotor. Note that a method for joining the component made of the Ni-based
alloy for forging of the embodiment and the component made of the conventional heat
resistant alloy is not limited to welding, but the components can be fastened by a
bolt and a nut, for example.
[0062] As a result that components to construct the turbine rotor are fabricated separately
as above, it is possible to manufacture a turbine rotor usable under a high temperature
environment of 700°C or higher with a small steel ingot of an Ni-based alloy. Note
that depending on a temperature condition to be used, it is possible to manufacture
the whole turbine rotor by the above-described method for manufacturing the turbine
rotor.
[0063] A rotor blade, a stationary blade, and a screwing member being turbine components
are fabricated as below, for example.
[0064] First, composition components to constitute the Ni-based alloy for forging of the
embodiment are vacuum induction melted (VIM) and electroslag remelted (ESR). Subsequently,
a molten alloy is poured into a predetermined mold form under a reduced pressure atmosphere
to fabricate an ingot, and a soaking treatment is performed. Then, the ingot is disposed
in a mold form corresponding to a shape of the above-described turbine component and
a forging treatment, a solution treatment, an aging treatment, and the like are carried
out, so that the rotor blade, the stationary blade, and the screwing member are fabricated.
In other words, the rotor blade, the stationary blade, and the screwing member are
fabricated by die forging.
[0065] Further, as another method (double melt), composition components to constitute the
Ni-based alloy for forging of the embodiment are vacuum induction melted (VIM) and
vacuum arc remelted (VAR), for example. Subsequently, a molten alloy is poured into
a predetermined mold form under a reduced pressure atmosphere to fabricate an ingot.
Then, a soaking treatment is applied to the ingot, and similarly to the above, a forging
treatment, a solution treatment, an aging treatment, and the like are carried out,
so that the rotor blade, the stationary blade, and the screwing member can be fabricated.
[0066] Further, as another method (triple melt), composition components to constitute the
Ni-based alloy for forging of the embodiment are vacuum induction melted (VIM), electroslag
remelted (ESR), and vacuum arc remelted (VAR), for example. Subsequently, a molten
alloy is poured into a predetermined mold form under a reduced pressure atmosphere
to fabricate an ingot. Then, a soaking treatment is applied to the ingot, and similarly
to the above, a forging treatment, a solution treatment, an aging treatment, and the
like are carried out, so that the rotor blade, the stationary blade, and the screwing
member can be fabricated.
[0067] A pipe being a forged component of the embodiment is fabricated as below, for example.
[0068] First, composition components to constitute the Ni-based alloy for forging of the
embodiment are electric furnace melted (EF) and argon-oxide decarburization (AOD)
is performed, so that an ingot is fabricated. Subsequently, a soaking treatment is
applied to the ingot. The ingot is drilled by a vertical press, so that an element
pipe of a cup shape is fabricated. Then, treatment and reheating by a mandrel and
a die are repeated by a horizontal press, so that the element pipe is formed in a
shape of a pipe. This treatment method is an Ehrhardt push bench pipe manufacturing
method. Then, a solution treatment, an aging treatment, and the like are carried out,
so that the pipe is fabricated.
[0069] Note that the methods for fabricating the turbine rotor, the rotor blade, the stationary
blade, the screwing member, and the pipe are not limited to the methods described
above. Further, the above-described forged components such as a turbine rotor, a rotor
blade, a stationary blade, a screwing member, and a pipe can be applied to power generation
turbines, such as a steam turbine, a gas turbine, and a CO
2 turbine, for example.
[0070] Here, the above-described each heat treatment performed in manufacturing the Ni-based
alloy for forging and the turbine component will be explained. Note that a temperature
in each heat treatment is set in each range described below, in correspondence with
the Ni-based alloy for forging, the turbine component, and the like to be treated.
Further, a time for each treatment is set properly in correspondence with the Ni-based
alloy for forging, the turbine component, and the like to be treated.
[0071] In the soaking treatment, it is necessary to heat the alloy at a high temperature
for a sufficient time, in order to decrease segregation of a chemical component by
thermal diffusion. Thus, the soaking treatment is preferable to be performed in a
temperature range of 1000 to 1200°C.
[0072] Forging is required to be performed in a range of temperature where sufficient ductility
of a material can be obtained to a zero ductility temperature. Thus, forging is preferable
to be performed in a temperature range of 950 to 1100°C.
[0073] In the solution treatment, a temperature range of 1050 to 1200°C is preferable to
be maintained for 1 to 24 hours. Here, the solution treatment is performed for the
purpose of solid dissolving an alloy element into a matrix sufficiently to obtain
an effect of solid solution strengthening sufficiently, and of enabling precipitation
control of a precipitate by a heat treatment thereafter. Further, the solution treatment
is sometimes performed for the purpose of adjusting a grain size.
[0074] When a temperature of the solution treatment is lower than 1050°C, an alloy element
is not solid dissolved into the matrix completely, and strengthening by a solid solution
strengthening element is not carried out sufficiently. Further, it also becomes difficult
to control a precipitated form of a precipitates by the heat treatment after the solution
treatment. On the other hand, when the temperature of the solution treatment is over
1200°C, coarsening of the grain size is brought about, and a mechanical strength is
reduced. Thus, the temperature of the solution treatment is set to be 1050 to 1200°C.
Further, it is further preferable that the temperature of the solution treatment is
set to be 1050 to 1150°C. Note that the Ni-based alloy and the turbine component having
been subjected to the solution treatment are cooled to a room temperature by water
cooling, forced air cooling, or the like, for example.
[0075] Next, there will be described an aging treatment applied to the Ni-based alloy and
the turbine component which have been cooled to the room temperature after the solution
treatment.
[0076] In the aging treatment, a temperature range of 700 to 800°C is preferable to be maintained
for 5 to 50 hours. The aging treatment can be performed in multiple stages. Note that
after the aging treatment the Ni-based alloy and the turbine component are cooled
to the room temperature by water cooling or furnace cooling, for example.
[0077] Here, the reason why the temperature and the time in the aging treatment are set
in the above-described range will be explained.
[0078] A major object of the aging treatment is control of a precipitation form of a γ'
phase precipitated in a grain. Further, the aging treatment influences a property
of a grain boundary. Therefore, for the aging treatment, the temperature and the time
condition are required to be determined in consideration of structures in the grain
and the grain boundary.
[0079] Fig. 2 and Fig. 3 are views showing electron micrographs of Ni-based alloys in order
to explain precipitation forms of carbides precipitated into grain boundaries by conditions
of the aging treatments. A composition of the Ni-based alloy shown here is 0.04% of
C, 18% of Cr, 12% of Co, 9% of Mo, 1.3% of Al, 1.4% of Ti, 0.003% of B, 0.1% of Ta,
0.3% of Nb, and the balance being Ni. Fig. 2 shows the microstructure having been
subjected to the aging treatment of 850°C for 10 hours, and Fig. 3 shows the microstructure
having been subjected to the aging treatment of 750°C for 10 hours. Further, the soaking
treatment and the solution treatment are performed in the above-described ranges.
Note that Fig. 2 and Fig. 3 also show precipitates 13 (γ' phase).
[0080] In an ordinary aging treatment, as shown in Fig. 2, a film-shaped carbide 11 is precipitated
in a manner to cover the grain boundary of the N-based alloy. The film-shaped carbide
11 is a brittle carbide (M
23C
6 type carbide) whose major constituents are Cr and Mo. The carbide 11 promotes destruction
of the grain boundary and substantially reduces a toughness of the material. Therefore,
it has been considered that it is necessary to perform the aging treatment preventing
precipitation of such a film-shaped carbide 11 covering the grain boundary.
[0081] However, as shown in Fig. 3, a thickness of the film-shaped carbide covering the
grain boundary becomes small, depending on the condition of the aging treatment. Note
that the carbide is continuously precipitated along the grain boundary. As a result
of a material testing, the inventors clarify that reduction of ductility/toughness
does not occur in a case where a thickness of a carbide is sufficiently small. The
above-described temperature and time are specified to a range satisfying both fine
precipitation of a γ' phase and suppression of coarsening of the carbide covering
the grain boundary.
[0082] When the temperature of the aging treatment is lower than 700°C, coarsening of the
carbide covering the grain boundary can be suppressed, but growth of the γ' phase
is quite slow. Thus, improvement of a mechanical strength by precipitation of the
γ' phase cannot be obtained. On the other hand, when the temperature of the aging
treatment is over 800°C, fine precipitation of the γ' phase is achieved, and a sufficient
strength is obtained. However, coarsening of the carbide covering the grain boundary
is significant, and a toughness is reduced.
[0083] Under the circumstances, the temperature of the aging treatment is set to be 700
to 800°C. Here, for the purpose of early precipitation of the γ' phase, heat treatments
of multiple stages, for example, two stages, can be performed in the aging treatment.
In also such a case, the temperature is set within the above-described temperature
range of the aging treatment. Further, the whole heat treatment time in multiple stages
is also set within the above-described time range of the aging treatment. For example,
there can be exemplified a treatment in which a temperature of 800°C is maintained
for 10 hours and thereafter a temperature of 750°C is maintained for 20 hours. Note
that temperature reduction from 800°C to 750°C is performed by furnace cooling, for
example.
[0084] Cooling after the aging treatment is performed by furnace cooling or air cooling,
for example. When the aging treatment is performed in multiple stages, cooling between
each aging treatment is performed by furnace cooling, for example, as described above.
Then, cooling is not performed to reach a room temperature but the aging treatment
is performed continuously.
[0085] Here, an intermediate heat treatment can be applied, before the aging treatment is
performed, to the Ni-based alloy and the turbine component which are cooled to the
room temperature after the solution treatment. An object of the intermediate heat
treatment is to form a block-shaped carbide intermittently along a grain boundary,
first, before the aging treatment, in order to suppress precipitation or coarsening
of a film-shaped carbide covering the grain boundary. This carbide is also a carbide
whose major constituents are Cr and Mo.
[0086] The intermediate heat treatment is preferable to be performed in a temperature range
of 1000 to 1050°C. In cases where the intermediate heat treatment temperature is lower
than 1000°C and where the intermediate heat treatment temperature is over 1050°C,
the block-shaped carbide is not precipitated. A time of the intermediate treatment
is set properly in correspondence with the Ni-based alloy and the turbine component
to be treated.
[0087] Note that in a case where a content ratio of C (carbon) is sufficiently small, precipitation
of a film-shaped carbide on a grain boundary is not prominent, and such an intermediate
heat treatment can be omitted. The case where the content ratio of C is sufficiently
small is, though it varies depending on a grain size or the like, a case where the
content ratio of C is 0.04% or less, for example. Cooling after the intermediate heat
treatment is performed by furnace cooling, water cooling, or forced air cooling, for
example. Then, the Ni-based alloy and the turbine component are cooled to the room
temperature.
(Influence of Chemical Composition)
[0088] Hereinafter, it will be described that the Ni-based alloy for forging of the embodiment
is excellent in a strength characteristic and forgeability.
[0089] Table 1 shows chemical compositions of a sample 1 to a sample 21 used for evaluation
of a strength characteristic, forgeability, and the like. Note that the samples 1,
4, 5, 8 and 13 shown in Table 1 are Ni-based alloys within a chemical composition
range of the Ni-based alloy for forging of the embodiment, and the samples 2, 3, 6,
7, 9 to 12 and 14 to 21 are Ni-based alloys whose compositions are not within the
chemical composition range of the Ni-based alloy for forging of the embodiment, and
are comparative examples.
[Table 1]
Mass% |
|
Ni |
C |
Cr |
Co |
Mo |
Al |
Ti |
B |
Ta |
Nb |
Mo+0.176Cr +0.037Co |
Sample 1 |
Balance |
0.04 |
18.5 |
12.3 |
9.0 |
1.3 |
1.4 |
0.003 |
0.11 |
0.3 |
12.7 |
Sample 2 |
Balance |
0.012 |
15.1 |
10.5 |
11.7 |
2.5 |
0.9 |
0.003 |
0.65 |
0.11 |
14.7 |
Sample 3 |
Balance |
0.03 |
22.1 |
14.5 |
5.8 |
1.4 |
2.5 |
0.003 |
0.06 |
0.66 |
10.2 |
Sample 4 |
Balance |
0.06 |
16.5 |
12.5 |
9.2 |
1.0 |
1.0 |
0.005 |
0.64 |
0.66 |
12.6 |
Sample 5 |
Balance |
0.02 |
18.1 |
10.2 |
9.5 |
1.4 |
1.4 |
0.004 |
0 |
0.12 |
13.1 |
Sample 6 |
Balance |
0.03 |
25.1 |
12.6 |
9.5 |
1.8 |
1.0 |
0.003 |
0 |
0.35 |
14.4 |
Sample 7 |
Balance |
0.05 |
24.1 |
14.3 |
10.0 |
1.3 |
1.3 |
0.003 |
0 |
0.65 |
14.8 |
Sample 8 |
Balance |
0.05 |
18.2 |
11.2 |
8.1 |
1.1 |
1.2 |
0.003 |
0.06 |
0 |
11.7 |
Sample 9 |
Balance |
0.02 |
20.1 |
12.6 |
6.3 |
2.0 |
0.86 |
0.003 |
0.1 |
0 |
10.3 |
Sample 10 |
Balance |
0.02 |
18.2 |
11.2 |
11.0 |
1.2 |
1.7 |
0.003 |
0.66 |
0 |
14.6 |
Sample 11 |
Balance |
0.05 |
19.0 |
10.8 |
11.0 |
0.85 |
1.5 |
0.003 |
0 |
0 |
14.7 |
Sample 12 |
Balance |
0.03 |
22.5 |
13.5 |
9.5 |
1.5 |
1.4 |
0.003 |
0 |
0 |
14.0 |
Sample 13 |
Balance |
0.05 |
24.1 |
12.6 |
7.1 |
1.5 |
1.0 |
0.003 |
0 |
0 |
11.8 |
Sample 14 |
Balance |
0.05 |
18.3 |
12.6 |
9.0 |
0.6 |
0.6 |
0.003 |
0.12 |
0.29 |
12.7 |
Sample 15 |
Balance |
0.05 |
26.5 |
15.3 |
7.8 |
1.3 |
1.4 |
0.003 |
0.11 |
0.3 |
13.0 |
Sample 16 |
Balance |
0.01 |
13.8 |
9.0 |
12.1 |
2.4 |
1.8 |
0.003 |
0.11 |
0.33 |
14.9 |
Sample 17 |
Balance |
0.01 |
19.5 |
10.0 |
11.1 |
3.1 |
0.8 |
0.003 |
0 |
0 |
14.9 |
Sample 18 |
Balance |
0.05 |
20.1 |
12.3 |
10.1 |
1.0 |
3.2 |
0.003 |
0 |
0 |
14.1 |
Sample 19 |
Balance |
0.05 |
24.1 |
14.3 |
11.5 |
1.3 |
1.3 |
0.003 |
0.12 |
0.31 |
16.3 |
Sample 20 |
Balance |
0.05 |
18.2 |
12.6 |
11.8 |
1.2 |
1.3 |
0.003 |
0.11 |
0.28 |
15.5 |
Sample 21 |
Balance |
0.05 |
17.0 |
12.6 |
6.3 |
1.1 |
0.9 |
0.003 |
0.1 |
0.31 |
9.8 |
[0090] A strength characteristic was evaluated by a tensile test, a toughness was evaluated
by a Charpy impact test, and forgeability was evaluated by visual observation. Further,
a thickness of a film-shaped carbide covering a grain boundary was measured by microstructure
observation.
[0091] A test piece used in each test was fabricated as follows.
[0092] Each of the Ni-based alloys of the sample 1 to the sample 21 having the chemical
compositions shown in Table 1 was melted in a vacuum induction melting furnace, to
fabricate an ingot.
[0093] Subsequently, a soaking treatment was applied to the ingot at 1050°C for 5 hours.
Thereafter, forging was performed in a temperature range of 950 to 1100°C (reheating
temperature was 1100°C) by a 500 kgf hammer forging machine. After forging, a solution
treatment was performed at a temperature of 1100°C for 4 hours, and thereafter, cooling
to a room temperature was carried out by air cooling. After cooling, an intermediate
heat treatment was performed at a temperature of 1025°C for 10 hours, and thereafter,
cooling to the room temperature was carried out by furnace cooling. After cooling,
a two-stage aging treatment at a temperature of 800°C for 10 hours and subsequently
at a temperature of 750°C for 20 hours was performed continuously. Thereafter, cooling
to the room temperature was performed by air cooling, so that a forged product was
made.
[0094] Then, from the above forged product, the test piece for the tensile test and the
Charpy impact test was obtained.
[0095] The tensile test was performed in accordance with JIS Z 2241, and measurement of
0.2% proof stress and a tensile strength at the room temperature was performed. The
Charpy impact test was performed in accordance with JIS Z 2242, and measurement of
a Charpy impact value was performed.
[0096] For evaluation of forgeability, the sample after the soaking treatment described
above was forged by a 500 kgf hammer forging machine, to fabricate a test piece of
a solid columnar shape with a diameter of 125 mm and a length of 210 mm. Further,
the forging treatment was performed until a forging ratio (a forging ratio based on
JIS G 0701 (representation of a forging ratio of a steel product forging operation)
became 3. Note that the forging treatment was performed in a range of 950 to 1100°C.
When a temperature of the test piece being an object to be forged was reduced, that
was, when the object to be forged was being cured, reheating to the reheating temperature
of 1100°C was performed and the forging treatment was repeated. The evaluation of
forgeability was performed by visual observation of existence/absence of a forging
crack after the test piece was cooled.
[0097] Here, the forging ratio is obtained by dividing a cross-sectional area of the object
to be forged before application of the forging treatment, the cross-sectional area
being vertical in a direction where the object to be forged is to be expanded, by
a cross-sectional area of the object to be forged after the forging treatment, the
cross-sectional area being vertical in a direction where the object to be forged has
been expanded.
[0098] In measurement of the thickness of the film-shaped carbide covering the grain boundary,
the forged product cooled to the room temperature after the aging treatment was used.
The thickness of the carbide was obtained by image-analyzing an electron micrograph
photographed at 20000 magnification by using a field-emission scanning electron microscope.
In each forged product, 5 representative grain boundaries were selected, and thicknesses
of 20 points of the carbide were measured per each grain boundary. Then, the above
thicknesses were arithmetically averaged, to obtain an average thickness of the carbide.
[0099] Test results and observation results are shown in Table 2. In Table 2, a case where
a forging crack does not exist is denoted as "No", and further, evaluation of forgeability
is denoted as "○" in order to indicate that forgeability is excellent. On the other
hand, a case where a forging crack exists is denoted as "Yes", and evaluation of forgeability
is denoted as "x" to indicate that forgeability is inferior.
[Table 2]
|
Forging state |
Room temperature tensile test |
Charpy impact test |
Structure observation |
|
Forging crack |
Forgeability |
0.2% proof stress, MPa |
Tensile strength, MPa |
Impact value, J/cm2 |
Average thickness of carbide, nm |
Sample 1 |
No |
○ |
585 |
1042 |
78 |
182 |
Sample 2 |
No |
○ |
668 |
1112 |
68 |
63 |
Sample 3 |
No |
○ |
704 |
1187 |
66 |
125 |
Sample 4 |
No |
○ |
570 |
1038 |
55 |
191 |
Sample 5 |
No |
○ |
588 |
1095 |
83 |
80 |
Sample 6 |
No |
○ |
632 |
1098 |
56 |
85 |
Sample 7 |
No |
○ |
648 |
1121 |
54 |
135 |
Sample 8 |
No |
○ |
591 |
1010 |
60 |
148 |
Sample 9 |
No |
○ |
641 |
1084 |
67 |
95 |
Sample 10 |
No |
○ |
549 |
1000 |
64 |
84 |
Sample 11 |
No |
○ |
555 |
1022 |
61 |
191 |
Sample 12 |
No |
○ |
591 |
1045 |
54 |
170 |
Sample 13 |
No |
○ |
625 |
1048 |
61 |
210 |
Sample 14 |
No |
○ |
460 |
906 |
112 |
189 |
Sample 15 |
Yes |
× |
584 |
1001 |
51 |
221 |
Sample 16 |
Yes |
× |
742 |
1167 |
44 |
58 |
Sample 17 |
Yes |
× |
768 |
1144 |
36 |
64 |
Sample 18 |
Yes |
× |
750 |
1058 |
30 |
204 |
Sample 19 |
Yes |
× |
741 |
1120 |
67 |
206 |
Sample 20 |
Yes |
× |
1017 |
1100 |
61 |
175 |
Sample 21 |
No |
○ |
482 |
950 |
56 |
168 |
[0100] As shown in Table 2, the sample 1 to the sample 13 are higher in both 0.2% proof
stresses and tensile strengthes compared with the sample 14. It is considered that
in the sample 1 to the sample 13, the 0.2% proof stresses and the tensile strengths
are high in values because sufficient solid solution strengthening and precipitation
strengthening are enhanced. Further, the sample 1 to the sample 13 are excellent in
forgeability, and thicknesses of the carbides are 250 nm or less. Further, results
of Charpy impact values of the sample 1 to the sample 13 each indicate a value of
50 J/cm
2 or more. Therefore, it is confirmed that the sample 1 to the sample 13 have practically
sufficient toughnesses.
[0101] On the other hand, when a value of "Mo + 0.176Cr + 0.037Co" is less than 11 mass%
as in the sample 21, a sufficient 0.2% proof stress or a tensile strength are not
obtained even in a case where each alloy component is within a chemical composition
range prescribed in the present embodiment. The sample 15 to the sample 20 indicate
high values in 0.2% proof stresses and the tensile strengths, but are inferior in
forgeability. This is considered to be a result of excessive addition of a strengthening
element.
[0102] As described above, in the Ni-based alloy departing from the chemical composition
range prescribed in the present embodiment or a range of "Mo + 0.176Cr + 0.037Co",
a result which is excellent in both the strength characteristic and the forgeability
is not obtained.
(Influence of Heat Treatment)
[0103] Here, in the sample 1, while conditions of the intermediate heat treatment and the
aging treatment were changed, there were performed tensile tests, Charpy impact tests,
evaluation of forgeability, and measurement of thicknesses of film-shaped carbides
covering grain boundaries. Note that methods of respective test, evaluation of forgeability,
measurement of the thickness of the carbide were the same as the aforementioned methods.
[0104] By using the sample 1 shown in Table 1, heat treatments were performed under respective
conditions of the intermediate heat treatment and the aging treatment shown in Table
3. Note that process steps other than the intermediate heat treatment and the aging
treatment are the same as those in the method for fabricating the test piece. In Table
3, for example, "800°C × 10h" means that the heat treatment is performed while a temperature
of 800°C is maintained for 10 hours. Further, in the aging treatment, when the two-stage
heat treatment is performed, heat treatment conditions are indicated in columns of
a first stage and a second stage.
[Table 3]
|
Intermediate heat treatment condition |
Aging treatment condition |
First stage |
Second stage |
Sample 1 |
1025°C×10h |
800°C×10h |
750°C×20h |
Sample 22 |
- |
750°C×10h |
- |
Sample 23 |
- |
750°C×30h |
- |
Sample 24 |
- |
750°C×48h |
- |
Sample 25 |
- |
800°C×10h |
- |
Sample 26 |
- |
800°C×30h |
- |
Sample 27 |
- |
800°C×40h |
- |
Sample 28 |
- |
800°C×10h |
780°C×20h |
Sample 29 |
- |
800°C×10h |
780°C×30h |
Sample 30 |
1025°C×10h |
750°C×10h |
- |
Sample 31 |
1025°C×10h |
800°C×10h |
- |
Sample 32 |
- |
675°C×10h |
- |
Sample 33 |
- |
675°C×30h |
- |
Sample 34 |
- |
850°C×10h |
- |
Sample 35 |
- |
850°C×30h |
- |
Sample 36 |
- |
850°C×10h |
780°C×20h |
Sample 37 |
- |
850°C×10h |
780°C×30h |
Sample 38 |
1025°C×10h |
675°C×30h |
- |
Sample 39 |
1025°C×10h |
850°C×30h |
- |
[0105] The sample 1, and a sample 22 to a sample 31 shown in Table 3 are heat-treated under
the heat treatment condition of the present embodiment, and the other samples are
comparative examples heat-treated under a condition departing from a range of the
heat treatment condition of the present embodiment. Test results and observation results
are shown in Table 4.
[Table 4]
|
Room temperature tensile test |
Charpy impact test |
Structure observation |
0.2% proof stress, MPa |
Tensile strength, MPa |
Impact value, J/cm2 |
Average thickness of carbide, nm |
Sample 1 |
585 |
1042 |
78 |
182 |
Sample 22 |
606 |
1052 |
100 |
87 |
Sample 23 |
629 |
1081 |
74 |
112 |
Sample 24 |
624 |
1082 |
68 |
138 |
Sample 25 |
581 |
1030 |
72 |
191 |
Sample 26 |
589 |
1049 |
62 |
215 |
Sample 27 |
578 |
1042 |
64 |
201 |
Sample 28 |
574 |
1025 |
60 |
180 |
Sample 29 |
566 |
1011 |
58 |
225 |
Sample 30 |
620 |
1047 |
62 |
99 |
Sample 31 |
571 |
1055 |
60 |
148 |
Sample 32 |
468 |
920 |
79 |
170 |
Sample 33 |
488 |
955 |
72 |
210 |
Sample 34 |
531 |
989 |
38 |
260 |
Sample 35 |
508 |
959 |
33 |
425 |
Sample 36 |
541 |
1000 |
35 |
320 |
Sample 37 |
513 |
968 |
33 |
468 |
Sample 38 |
495 |
940 |
61 |
204 |
Sample 39 |
524 |
978 |
28 |
380 |
[0106] The sample 1 and the sample 22 to the sample 31 are higher in both 0.2% proof stresses
and tensile strengths compared with samples 32 to samples 39. Average thicknesses
of film-shaped carbides covering grain boundaries in the sample 1 and the sample 22
to the sample 31 are each 250 nm or less. The sample 1 and the sample 22 to the sample
31, having thin average thicknesses of the carbides, exhibit higher Charpy impact
values compared with the sample 34 to the sample 37 and the sample 39.
[0107] As described above, under the aging treatment condition prescribed in the present
embodiment, it is possible to simultaneously attain fine precipitation of the γ' phase
in the grain and suppression of coarsening of the carbide covering the grain boundary.
Thereby, high values are obtained in both the tensile strength and the Charpy impact
value.
[0108] On the other hand, in the sample departing from the aging treatment condition prescribed
in the present embodiment, a result which is excellent in both the tensile strength
and the Charpy impact value is not obtained.
[0109] According to the embodiment described hereinabove, it becomes possible to have excellent
strength characteristic and forgeability.
[0110] While certain embodiments have been described, these embodiments have been presented
by way of example only, and are not intended to limit the scope of the inventions.