Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
CA 02508698 2011-06-03
1
SPECIFICATION
Ni-BASED SINGLE CRYSTAL SUPER ALLOY
TECHNICAL FIELD
The present invention relates to a Ni-based single crystal super alloy, and
more particularly,
to a technology employed for improving the creep characteristics of Ni-based
single crystal super
alloy.
BACKGROUND ART
An example of the typical composition of Ni-based single crystal super alloy
developed for
use as a material for moving and stationary blades subject to high
temperatures such as those in
aircraft and gas turbines is shown in Table 1.
Table 1
Alloy Elements wt o)
name Al Ti Ta Nb Mo W Re C Zr Hf Cr Co Ru Ni
CMSX-2 6.0 1.0 6.0 - 1.0 8.0 - - - - 8.0 5.0 - Rem
CMSX-4 5.6 1.0 6.5 - 0.6 6.0 3.0 - - - 6.5 9.0 - Rem
ReneN6 6.0 - 7.0 03 1.0 6.0 5.0 - - 0.2 4.0 13.0 - Rem
CMSX-10 5.7 03 8.4 0.1 0.4 5.5 6.3 - - 0.03 2.3 3.3 - Rem
K
3B 5.7 0.5 8.0 - - 5.5 6.0 0.05 - 0.15 5.0 12.5 3.0 Rem
In the above-mentioned Ni-based single crystal super alloys, after performing
solution
treatment at a prescribed temperature, aging treatment is performed to obtain
an Ni-based single
crystal super alloy. This alloy is referred to as a so-called precipitation
hardened alloy, and has a
from in which the precipitation phase in the form of a y' phase is
precipitated in a matrix in the form
of a y phase.
TM
Among the alloys listed in Table 1, CMSX-2 (Cannon-Muskegon, US Patent No.
4,582,548) is a first generation alloy, CMSX-4 (Cannon Muskegon, US Patent No.
4,643,782) is a
CA 02508698 2011-06-03
2
second-generation alloy, ReneN6 (General Electric, US Patent No. 5,455,120)
and CMSX-10K
(Canon-Muskegon, US Patent No. 5,366,695) are third-generation alloys, and 3B
(General Electric,
US Patent No. 5,151,249) is a fourth-generation alloy.
Although the above-mentioned CMSX-2, which is a first-generation alloy, and
CMSX-4,
which is a second-generation alloy, have comparable creep strength at low
temperatures, since a
large amount of the eutectic y' phase remains following high-temperature
solution treatment, their
creep strength is inferior to third-generation alloys.
In addition, although the third-generation alloys of ReneN6 and CMSX 10 are
alloys
designed to have improved creep strength at high temperatures in comparison
with
second generation alloys, since the composite ratio of Re (5 wt% or more)
exceeds the amount of
Re that dissolves into the matrix (y phase), the excess Re compounds with
other elements and as a
result, a so-called TCP (topologically close packed) phase precipitates at
high temperatures causing
the problem of decreased creep strength.
In addition, making the lattice constant of the precipitation phase (y' phase)
slightly smaller
than the lattice constant of the matrix (y phase) is effective in improving
the creep strength of
Ni-based single crystal super alloys. However, since the lattice constant of
each phase fluctuates
greatly fluctuated according to the composite ratios of the composite elements
of the alloy, it is
difficult to make fine adjustments in the lattice constant and as a result,
there is the problem of
considerable difficulty in improving creep strength.
In consideration of the above circumstances, the object of the present
invention is to provide
a Ni-based single crystal super alloy that makes it possible to improve
strength by preventing
precipitation of the TCP phase at high temperatures.
DISCLOSURE OF INVENTION
The following constitution is employed in the present invention in order to
achieve the above
object,
The Ni-based single crystal super alloy of the present invention is
characterized by having a
composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta,1.1-4.5 wt~/o
ofMo, 4.0-10.0 wt /a
of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-9.9 wt% of Co
and 4.1-14.0 wt%
CA 02508698 2011-06-03
3
of Ru in terms of its weight ratio, with the remainder consisting of Ni and
unavoidable impurities.
In addition, the Ni-based single crystal super alloy of the present invention
is characterized
by having a composition comprising 5.0-7.0 wt% of Al, 4.0-6.0 wt% of Ta,1.1-
4.5 wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
and 4.1-14.0 wt% of Ru in terms of weight ratio, with the remainder consisting
of Ni and
unavoidable impurities.
In addition, the Ni-based single crystal super alloy of the present invention
is characterized
by having a composition comprising 5.0-7.0 wt% of Al, 4.0-6.0 wt% of Ta,
2.94.5 wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co and
4.1-14.0 wt% of Ru in terms of weight ratio, with the remainder consisting of
Ni and unavoidable
impurities.
According to the above Ni-based single crystal super alloy, precipitation of
the TCP phase,
which causes a decrease in creep strength, during use at high temperatures is
inhibited by the
addition of Ru. In addition, by setting the composite ratios of other
composite elements within
their optimum ranges, the lattice constant of the matrix (y phase) and the
lattice constant of the
precipitation phase (y' phase) can be made to have optimum values.
Consequently, strength at
high temperatures can be enhanced. Furthermore, since the composition of Ru is
4.1-14.0 wt%,
precipitation of the TCP phase, which causes a decrease in creep strength,
during use at high
temperatures, is inhibited.
In addition, the Ni-based single crystal super alloy of the present invention
is preferably
having a composition comprising 5.9 wt% of Al, 5.9 wt% of Ta, 3.9 wt% of Mo,
5.9 wt% of W,
4.9wt%ofRe, 0.10wt%ofElf, 2.9wt%ofCr, 5.9 wt% of Co and 5.0 wt% ofRu interms
of
weight ratio, with the remainder consisting of Ni and unavoidable impurities,
in the Ni-based single
crystal super alloys previously described
According to an Ni-based single crystal super alloy having this composition,
the creep
endurance temperature at 137 MPa and 1000 hours can be made to be 1344 K (1071
C).
In addition, the Ni-based single crystal super alloy of the present invention
is preferably
having a composition comprising 5.8 wt% of Co, 2.9 wt% of Cr, 3.1 wt% of Mo,
5.8 wt% of W,
5.8wt%ofAl,5.6wt%ofTa,5.0wt%ofRu,4.9wt%ofReand0.10wt%ofHfinterms of
CA 02508698 2011-06-03
4
weight ratio, with the remainder consisting of Ni and unavoidable impurities,
in the Ni-based single
crystal super alloys previously described.
According to an Ni-based single crystal super alloy having this composition,
the creep
endurance temperature at 137 MPa and 1000 hours can be made to be 1366 K (1093
C).
In addition, the Ni-based single crystal super alloy of the present invention
is preferably
having a composition comprising 5.8 wt% of Co, 2.9 wt% of Cr, 3.9 wt% of Mo,
5.8 wt% of W,
5.8wt%ofAl,5.8wt%(5.82wt%)or5.6wt%ofTa,6.0wt%ofRu,4.9wt%of,Reand0.10
wt% of Hf in terms of weight ratio, with the remainder consisting of Ni and
unavoidable impurities,
in the Ni-based single crystal super alloys previously described.
According to an Ni-based single crystal super alloy having this composition,
the creep
endurance temperature at 137 MPa and 1000 hours can be made to be 1375 K (1102
C) or 1379
K (1106 C).
Furthermore, 0-2.0 wt% of Ti in temps of weight ratio can be included in the
Ni-based single
crystal super alloys previously described.
Furthermore, 0-4.0 wt% ofNb in terms of weight ratio can be included in the Ni-
based
single crystal super alloys previously described.
Furthermore, at least one of elements selected from B, C, Si, Y, La, Ce, V and
Zr can be
included in the Ni-based single crystal super alloys previously described.
In this case, it is preferable that 0.05 wt% or less of B, 0.15 wt% or less of
C, 0.1 wt% or less
of Si, 0.1 wt% or less of Y, 0.1 wt% or less of La, 0.1 wt% or less of Ce, l
wt% or less of V and 0.1
wt% or less of Zr in terms of weight ratio are included in the alloys.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta,1.1-4.5
wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
10.0-14.0 wt% of Ru, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15
wt% or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less
of La, 0.1 wt% or less
of Ce, l wt% or less of V and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.8-7.0 wt% of Al, 4.0-5.6 wt% of Ta, 3.3-4.5
wt% of Mo,
CA 02508698 2011-06-03
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.9-4.3 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt% of Ru, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15 wt%
or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce, l wt% or less of V and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta,1.1-4.5
wt% ofMo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.9-5.0 wt% of Cr, 0-
9.9 wt% of Co,
6.5-14.0 wt% of Ru, 4.0 wt% or less of Nb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15 wt%
or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce, l wt% or less of V and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-6.0 wt% of Ta, 3.3-4.5
wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt% of Ru, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt%o or
less of B, 0.15 wt%
or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce, l wt% or less of V and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-5.6 wt% of Ta, 3.34.5
wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt% of Ru, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15 wt%
or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce,1 wt% or less of V and 0.1 wt% or less ofZr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta, 3.1-4.5
wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt o of Ru, 4.0 wt% or less ofNb, 0.05 wt% or less of B, 0.15 wt% or
less of C, 0.1 wt%
or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of La, 0.1 wt% or less of
Ce, l wt% or less of V
and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
CA 02508698 2011-06-03
6
having a composition comprising 5.8-7.0 wt~/o of Al, 4.0-10.0 wt% of Ta, 3.1-
4.5 wt% of Mo,
4.0-10.0 wM/o of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt% of Ru, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15 wo
or less of C,' 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce, l wt/o or less of V and 0.1 wt% or less of Zr.
Furthermore, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta, 3.1-4.5
wt% of Mo,
4.0-10.0 wt% of W, 3.1-8.0 wt% of Re, 0-0.50 wt% of Hf, 2.9-4.3 wt% of Cr, 0-
9.9 wt% of Co,
4.1-14.0 wt%. ofRu, 4.0 wt% or less ofNb, 2.0 wt% or less of Ti, 0.05 wt% or
less of B, 0.15 wt%
or less of C, 0.1 wt% or less of Si, 0.1 wt% or less of Y, 0.1 wt% or less of
La, 0.1 wt% or less of
Ce, l wt% or less of V and 0.1 wt% or less of Zr.
In addition, the above described Ni-based single crystal super alloy is more
preferably
having a composition comprising 5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta+Nb+Ti,
3.3-4.5 wt% of
Mo, 4.0-10.0 wt% of W, 3.1-8.0 wt% ofRe, 0-0.50 wt% of Hf, 2.0-5.0 wt%o of Cr,
0-9.9 wt% of
Co, 4.1-14.0 wt% of Ru, 0.05 wt% or less of B, 0.15 wt% or less of C, 0.1 wt%
or less of Si; 0.1
wt% or less off, 0.1 wt% or less of La, 0.1 wt% or less of Ce,1 wt% or less of
V and 0.1 wt% or
less of Zr.
Moreover, the Ni-based single crystal super alloy of the present invention is
characterized by
a2 <_ 0.999a1 when the lattice constant of the matrix is taken to be al and
the lattice constant of the
precipitation phase is taken to be a2 in the Ni-based single crystal super
alloys previously described
According to this Ni-based single crystal super alloy, the relationship
between al and a2 is
such that a2!5 0.999a1 when the lattice constant ofthe matrix is taken to be
al and the lattice
constant of the precipitation phase is taken to be a2, and since the lattice
constant a2 of the
precipitation phase is -0.1% or less of the lattice constant al of the matrix,
the precipitation phase
that precipitates in the matrix precipitates so as to extend continuously in
the direction
perpendicular to the direction of the load. Asa result, strength at high
temperatures can be
enhanced without dislocation defects moving within the alloy structure under
stress.
In this case, it is more preferable that the lattice constant of the crystals
of the precipitation
phase a2 is 0.9965 or less ofthe lattice constant of the crystals of the
matrix al
CA 02508698 2011-06-03
7
Furthermore, the Ni-based single crystal super alloy of the present invention
is characterized
by comprising the feature that the dislocation space of the alloy is 40 nm or
less.
BRIEF DESCRIPTION OF DRAWINGS
FIG. 1 is a diagram showing a relationship between change of lattice misfit of
the alloy and
creep rupture life of the alloy.
FIG. 2 is a diagram showing a relationship between dislocation space of the
alloy and creep
rupture life of the alloy.
FIG. 3 is a transmission electron microgram of the Ni-based single crystal
super alloy
showing an embodiment of the dislocation networks and dislocation space of the
Ni-based single
crystal super alloy of the present invention.
BEST MODE FOR CARRYING OUT THE INVENTION
The following provides a detailed explanation for carrying out the present
invention.
The Ni-based single crystal super alloy of the present invention is an alloy
comprised of Al,
Ta, Mo, W, Re, Hf, Cr, Co, Ru, Ni (remainder) and unavoidable impurities.
The above Ni-based single crystal super alloy is an alloy having a composition
comprising
5.0-7.0 wt% of Al, 4.0-10.0 wt% of Ta,1.1-4.5 wt% of Mo, 4.0-10.0 wt% of W,
3.1-8.0 Wt% of
Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-9.9 wt% of Co and 4.1-14.0 wt%
ofRu, with the
remainder consisting of Ni and unavoidable impurities.
In addition, the above Ni-based single crystal super alloy is an alloy having
a composition
comprising 5.0-7.0 w0/6 of Al, 4.0-6.0 wt% of Ta,1.1-4.5 wt% of Mo, 4.0-10.0
wt% of W, 3.1-8.0
wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-9.9 wt% of Co and 4.1-14.0
wt% of Ru, with
the remainder consisting ofNi and unavoidable impurities.
Moreover, the above Ni-based single crystal super alloy is an alloy having a
composition
comprising 5.0-7.0 wt% of Al, 4.0-6.0 wt% of Ta, 2.9-4.5 wt% ofMo, 4.0-10.0
wt% of W, 3.1-8.0
wt% of Re, 0-0.50 wt% of Hf, 2.0-5.0 wt% of Cr, 0-9.9 wt% of Co and 4.1-14.0
wt% ofRu, with
the remainder consisting of Ni and unavoidable impurities.
All of the above alloys have an austenite phase in the form of a y phase
(matrix) and an
CA 02508698 2011-06-03
8
intermediate regular phase in the form of a y' phase (precipitation phase)
that is dispersed and
precipitated in the matrix. They' phase is mainly composed of an intermetallic
compound
represented by Ni3AI, and the strength of the Ni-based single crystal super
alloy at high
temperatures is improved by this y' phase.
Cr is an element that has superior oxidation resistance and improves the high
temperature
corrosion resistance of the Ni-based single crystal super alloy. The composite
ratio of Cr is
preferably within the range of 2.0 wt% or more to 5.0 wt% or less, and more
preferably 2.9 wt%.
This ratio is more preferably within the range of 2.9 wt% or more to 5.0 wt%
or less, more
preferably within the range of 2.9 wt% or more to 4.3 wt% or less, and most
preferably 2.9 wt%.
If the composite ratio of Cr is less than 2.0 wt%, the desired high
temperature corrosion resistance
cannot be secured, thereby making this undesirable. If the composite ratio of
Cr exceeds 5.0 wt%,
in addition to precipitation of the y' phase being inhibited, harmful phases
such as a a phase or p
phase form that cause a decrease in strength at high temperatures, thereby
making this undesirable.
In addition to improving strength at high temperatures by dissolving into the
matrix in the
form of they phase in the presence of W and Ta, Mo also improves strength at
high temperatures
due to precipitation hardening. Furthermore, Mo also improves the
aftermentioned lattice misfit
and dislocation networks of the alloy which relate characteristics of this
alloy.
The composite ratio of Mo is preferably within the range of 1.1 wt% or more to
4.5 wt% or
less, more preferably within the range of 2.9 wt% or more to 4.5 wt% or less.
This ratio is more
preferably within the range of 3.1 wt% or more to 4.5 wt% or less, more
preferably within the
range of 3.3wt% or more to 4.5 wt/o or less, and most preferably 3.1 wt% or
3.9 wt%. Ifthe
composite ratio of Mo is less than 1.1 wt%, strength at high temperatures
cannot be maintained at
the desired level, thereby making this undesirable. If the composite ratio of
Mo exceeds 4.5 wt%,
strength at high temperatures decreases and corrosion resistance at high
temperatures also
decreases, thereby making this undesirable.
W improves strength at high temperatures due to the actions of solution
hardening and
precipitation hardening in the presence of Mo and Ta as previously mentioned.
The composite
ratio of W is preferably within the range of 4.0 wt% or more to 10.0 wt% or
less, and most
preferably 5.9 wt% or 5.8 wt%. If the composite ratio of W is less than 4.0
wt%, strength at high
CA 02508698 2011-06-03
9
temperatures cannot be maintained at the desired level, thereby making this
undesirable. If the
composite ratio of W exceeds 10.0 wt/o, high-temperature corrosion resistance
decreases, thereby
making this undesirable.
Ta improves strength at high temperatures due to the actions of solution
hardening and
precipitation hardening in the presence of Mo and W as previously mentioned,
and also improves
strength at high temperatures as a result of a portion of the Ta undergoing
precipitation hardening
relative to the y' phase. The composite ratio of Ta is preferably within the
range of 4.0 wt% or
more to 10.0 wt% or less, more preferably within the range of 4.0 wt% or more
to 6.0 wt% or less.
This ratio is more preferably within the range of 4.0 wt% or more to 5.6 wt%
or less, and most
preferably 5.6 wt% or 5.82 wt%. If the composite ratio of Ta is less than 4.0
wt%, strength at
high temperatures cannot be maintained at the desired level, thereby making
this undesirable. If
the composite ratio of Ta exceeds 10.0 wt%, the u phase and phase form that
cause a decrease in
strength at high temperatures, thereby making this undesirable.
Al improves strength at high temperatures by compounding with Ni to form an
intermetallic
compound represented by Ni3AI, which composes they' phase that finely and
uniformly disperses
and precipitates in the matrix, at a ratio of 60-70% in terms of volume
percent. The composite
ratio of Al is preferably within the range of 5.0 wt% or more to 7.0 wt% or
less. This ratio is more
preferably within the range of 5.8 wt% or more to 7.0 wt% or less, and most
preferably 5.9 wt% or
5.8 wt%. If the composite ratio of Al is less than 5.0 wt%, the precipitated
amount of they' phase
becomes insufficient, and strength at high temperatures cannot be maintained
at the desired level,
thereby making this undesirable. If the composite ratio of Al exceeds 7.0 wt%,
a large amount of
a coarse y phase referred to as the eutectic y' phase is formed, and this
eutectic y' phase prevents
solution treatment and makes it impossible to maintain strength at high
temperatures at a high level,
thereby making this undesirable.
Hf is an element that segregates at the grain boundary and improves strength
at high
temperatures by strengthening the grain boundary as a result of being
segregated at the grain
boundary between they phase and y' phase. The composite ratio of Hf is
preferably within the
range of 0.01 wt% or more to 0.50 wt% or less, and most preferably 0.10 wt%.
If the composite
ratio of Hf is less than 0.01 wt'/o, the precipitated amount of the y' phase
becomes insufficient and
CA 02508698 2011-06-03
strength at high temperatures cannot be maintained at the desired level,
thereby making this
undesirable. However, the composite ratio of Hf may be within the range of 0
wt% or more to
less than 0.01 wt%, if necessary. Furthermore, if the composite ratio of Hf
exceeds 0.50 wt%,
local melting is induced which results in the risk of decreased strength at
high temperatures, thereby
making this undesirable.
Co improves strength at high temperatures by increasing the solution limit at
high
temperatures relative to the matrix such as Al and Ta, and dispersing and
precipitating a fine /
phase by heat treatment The composite ratio of Co is preferably within the
range of 0.1 wt% or
more to 9.9 wt% or less, and most preferably 5.8 wt%. If the composite ratio
of Co is less than
0.1 wt%, the precipitated amount of the l phase becomes insufficient and the
strength at high
temperatures cannot be maintained, thereby making this undesirable. However,
the composite
ratio of Co may be within the range of 0 wt% or more to less than 0.1 wt%, if
necessary.
Furthermore, if the composite ratio of Co exceeds 9.9 wt%, the balance with
other elements such as
Al, Ta, Mo, W, Hf and Cr is disturbed resulting in the precipitation of
harmful phases that cause a
decrease in strength at high temperatures, thereby making this undesirable.
Re improves strength at high temperatures due to solution strengthening as a
result of
dissolving in the matrix in the form of they phase. On the other hand, if a
large amount of Re is
added, the harmful TCP phase precipitates at high temperatures, resulting in
the risk of decreased
strength at high temperatures. Thus, the composite ratio of Re is preferably
within the range of
3.1 wt% or more to 8.0 wt% or less, and most preferably 4.9 wt%. If the
composite ratio of Re is
less than 3.1 wt%, solution strengthening of they phase becomes insufficient
and strength at high
temperatures cannot be maintained at the desired level, thereby making this
undesirable. If the
composite ratio of Re exceeds 8.0 wt%, the TCP phase precipitates at high
temperatures and
strength at high temperatures cannot be maintained at a high level, thereby
making this undesirable.
Ru improves strength at high temperatures by inhibiting precipitation of the
TCP phase.
The composite ratio of Ru is preferably within the range of 4.1 wt% or more to
14.0 wt% or less.
This ratio is more preferably within the range of 10.0 wt% or more to 14.0 wt%
or less, or
preferably within the range of 6.5 wo or more to 14.0 wt% or less, and most
preferably 5.0 wt%,
6.0 wt% or 7.0 wt%. If the composite ratio of Ru is less than 1.0 wt%, the TCP
phase precipitates
CA 02508698 2011-06-03
11
at high temperatures and strength at high temperatures cannot be maintained at
a high level, thereby
making this undesirable. If the composite ratio of Ru is less than 4.1 wt%,
strength at high
temperatures decreases compared to the case when the composite ratio of Ru is
4.1 wt% or more.
Furthermore, if the composite ratio of Ru exceeds 14.0 wt%, the s phase
precipitates and strength at
high temperatures deceases which is also undesirable.
Particularly in the present invention, by adjusting the composite ratios of
Al, Ta, Mo, W, Hf,
Cr, Co and Ni to the optimum ratios, together with improving strength at high
temperatures by
setting the aftermentioned lattice misfit and dislocation networks of the
alloy which are calculated
from the lattice constant of they phase and' the lattice constant of the y'
phase within their optimum
ranges, and precipitation of the TCP phase can be inhibited by adding Ru.
Furthermore, by
adjusting the composite ratios of Al, Cr, Ta and Mo to the aforementioned
ratios, the production
cost for the alloy can be decreased. In' addition, relative strength of the
alloy can be increased and
the lattice misfit and dislocation networks of the alloy can be adjusted to
the optimum value.
In addition, in usage environments at a high temperature from 1273 K (1000 C)
to 1373K
(1100 C), when the lattice constant of the crystals that compose the matrix in
the form of the y
phase is taken to be al, and the lattice constant of the crystals that compose
the precipitation phase
in the form of the y' phase is taken to be a2, then the relationship between
al and a2 is preferably
such that a2 S 0.999a1. Namely, lattice constant a2 of the crystals of the
precipitation phase is
preferably -0.1% or less lattice constant al of the crystals of the matrix.
Furthermore, it is more
preferable that the lattice constant of the crystals of the precipitation
phase a2 is 0.9965 or less of the
lattice constant of the crystals of the matrix al. In this case, the above-
described relationship
between al and a2 becomes a2 _< 0.9965a1. In the following descriptions, the
percentage of the
lattice constant a2 relative to the lattice constant al is called "lattice
misfit".
In addition, in the case both of the lattice constants are in the above
relationship, since the
precipitation phase precipitates so as to extend continuously in the direction
perpendicular to the
direction of the load when the precipitation phase precipitates in the matrix
due to heat treatment,
creep strength can be enhanced without movement of dislocation defects in the
alloy structure in
the presence of stress.
In order to make the relationship between lattice constant al and lattice
constant a2 such that
CA 02508698 2011-06-03
12
a2< _ 0.999a1, the composition ofthe composite elements that compose the Ni-
based single crystal
super alloy is suitably adjusted.
FIG.1 shows a relationship between the lattice misfit of the alloy and the
time until the alloy
demonstrates creep rupture (creep rupture life).
In FIG. 1, when the lattice misfit is approximately - 0.35 or lower, the creep
rupture life is
approximately higher than the required value (the value shown by a dotted line
in a vertical axis of
the figure). Therefore, in the present invention, the preferable value of the
lattice misfit is
determined to - 0.35 or lower. In order to maintain the lattice misfit to -
0.35 or lower, the
composition of Mo is maintained to a high level, and the composition of the
other composite
elements is suitably adjusted.
According to the above Ni-based super crystal super alloy, precipitation of
the TCP phase,
which causes decreased creep strength, during use at high temperatures is
inhibited by addition of
Ru. In addition, by setting the composite ratios of other composite elements
to their optimum
ranges, the lattice constant of the matrix (y phase) and the lattice constant
of the precipitation phase
(y' phase) can be made to have optimum values. As a result, creep strength at
high temperatures
can be improved.
Ti can be further included in the above Ni-based super crystal super alloy.
The composite
ratio of Ta is preferably within the range of 0 wt% or more to 2.0 wt% or
less. If the composite
ratio of Ti exceeds 2.0 wt%, the harmful phase precipitates and the strength
at high temperatures
cannot be maintained, thereby making this undesirable.
Furthermore, Nb can be further included in the above Ni-based super crystal
super alloy.
The composite ratio of Nb is preferably within the range of 0 wt% or more to
4.0 wt% or less. If
the composite ratio of Nb exceeds 4.0 wt%, the harmful phase precipitates and
the strength at high
temperatures cannot be maintained, thereby making this undesirable.
Alternatively, strength at high temperatures can be improved by adjusting the
total
composite ratio of Ta, Nb and Ti (Ta+Nb+Ti) within the range of 4.0 wt% or
more to 10.0 wt~/o or
less.
Furthermore, in addition to the unavoidable impurities, B, C, Si, Y, La, Ce, V
and Zr and the
like can be included in the above Ni-based super crystal super alloy, for
example. When the alloy
CA 02508698 2011-06-03
13
includes at least one of elements selected from B, C, Si, Y, La, Ce, V and Zr,
the composite ratio of
each element is preferably 0.05 wt% or less of B, 0.15 wt% or less of C, 0.1
wt% or less of Si, 0.1
wt% or less of Y, 0.1 wt% or less of La, 0.1 wt% or less of Ce, 1 wt%o or less
of V and 0.1 wt% or
less of Zr. If the composite ratio of each element exceeds the above range,
the hannfiul phase
precipitates and the strength at high temperatures cannot be maintained,
thereby making this
undesirable.
Furthermore, in the above Ni-based single crystal super alloy, it is
preferable that a
dislocation space of the alloy is 40 nm or less. The reticulated dislocation
(displacement of atoms
which are connected as a line) in the alloy is called dislocation networks,
and a space between
adjacent reticulations is called "dislocation space". FIG. 2 shows a
relationship between the
dislocation space of the alloy and the time until the alloy demonstrates creep
rupture (creep rupture
life).
In FIG. 2, when the dislocation space is approximately 40 nm or lower, the
creep rupture life
is approximately higher than the required value (the value shown by a dotted
he in a vertical axis
of the figure). Therefore, in the present invention, the preferable value of
the dislocation space is
determined to 40 nm or lower. In order to maintain the dislocation space to 40
rim or lower, the
composition of Mo is maintained to a high level, and the composition of the
other composite
elements is suitably adjusted.
FIG. 3 is a transmission electron microgram of the Ni-based single crystal
super alloy
showing an embodiment (aftennentioned embodiment 3) of the dislocation
networks and
dislocation space of the Ni-based single crystal super alloy of the present
invention. As shown in
FIG. 3, in case of the Ni-based single crystal super alloy of the present
invention, the dislocation
space is 40 rim or lower.
In addition, some of the conventional Ni-based single crystal super alloys may
cause reverse
partitioning, however, in Ni-based single crystal super alloy of the present
invention does not cause
reverse partitioning.
Embodiments
The effect of the present invention is shown using following embodiments.
Melts of various Ni-based single crystal super alloys were prepared using a
vacuum melting
CA 02508698 2011-06-03
14
furnace, and alloy ingots were cast using the alloy melts. The composite ratio
of each of the alloy
ingots (reference examples 1-6, embodiments 1-14) is shown in Table 2.
Table 2
Sample Elements (wt%)
(alloy name) Al Ta Nb Mo W Re Hf Cr Co Ru Ni
Reference 6.0 5.8 3.2 6.0 5.0 0.1 3.0 6.0 2.0 Rem
Example 1
Reference 5.9 5.7 3.2 5.9 5.0 0.1 3.0 5.9 3.0 Rem
Example 2
Reference 6.0 6.0 4.0 6.0 5Ø 0.1 3.0 6.0 3.0 Rem
Example 3
Reference 5.9 5.9 4.0 5.9 5.0 0.1 3.0 5.9 4.0 Rem
Example 4
Reference 5.9 5.7 3.1 5.9 4.9 0.1 2.9 5.9 4.0 Rem
Example 5
Reference
5.7 5.7 2.9 7.7 4.8 0.1 2.9 5.7 3.0 Rem
Exam le 6
Embodiment 1 5.9 5.9 3.9 5.9 4.9 0.1 2.9 5.9 5Ø Rem
Embodiment 2 5.8 5.6 3.1 5.8 4.9 0.1 2.9 5.8 5.0 Rem
Embodiment 3 5.8 5.8 3.9 5.8 4.9 0.1 2.9 5.8 6.0 Rem
Embodiment 4 5.6 5.6 2.8 5.6 6.9 0.1 2.9 5.6 5.0 Rem
Embodiment 5 5.6 5.0 0.5 2.8 5.6 6.9 0.1 2.9 5.6 5.0 Rem
Embodiment 6 5.6 5.6 1.0 2.8 5.6 4.7 0.1 2.9 5.6 5.0 Rem
Embodiment 7 5.8 5.6 3.9 5.8 4.9 0.1 2.9 5.8 6.0 Rem
Embodiment 8 5.7 5.5 1.0 3.8 5.7 4.8 0.1 2.8 5.5 5.9 Rem
Embodiment 9 5.8 5.6 3.1 6.0 5.0 0.1 2.9 5.8 4.6 Rem
Embodiment 10 5.8 5.6 3.1 6.0 5.0 0.1 2.9 5.8 5.2 Rem
Embodiment 11 5.8 5.6 3.3 6.0 5.0 0.1 2.9 5.8 5.2 Rem
Embodiment 12 5.8 5.6 3.3 6.0 5.0 0.1 2.9 5.8 6.0 Rem
Embodiment 13 5.9 2.9 1.5 3.9 5.9- 4.9 0.1 2.9 5.9 6.1 Rem
Embodiment 14 5.7 5.52 3.1 5.7 4.8 0.1 2.9 5.7 7.0 Rem
CA 02508698 2011-06-03
Next, solution treatment and aging treatment were performed on the alloy
ingots followed by
observation of the state of the alloy structure with a scanning electron
microscope (SEM).
Solution treatment consisted of holding for 1 hour at 1573K (1300 C) followed
by heating to
1613K (1340 C) and holding for 5 hours. In addition, aging treatment consisted
of consecutively
performing primary aging treatment consisting of holding for 4 hours at 1273K
1423K (1000 C
-1150 C) and secondary aging treatment consisting of holding for 20 hours at
1143K (870 C).
As a result, a TCP phase was unable to be confirmed in the structure of each
sample.
Next, a creep test was performed on each sample that underwent solution
treatment and
aging treatment. The creep test consisted of measuring the time until each
sample (reference
examples 1-6 and embodiments 1-14) demonstrated creep rupture as the sample
life under each of
the temperature and stress conditions shown in Table 3. Furthermore, the value
of the lattice
misfit of each sample was also measured, and the result thereof is disclosed
in Table 3. In
addition, the value of the lattice misfit of each of the conventional alloys
shown in Table 1
(comparative examples 1-5) was also measured, and the result thereof is
disclosed in Table 4.
CA 02508698 2011-06-03
16
Table 3
Sample Creep test conditions/rupture life (h)
(alloy name) 1273K (1000 C) 1373K (1100 C) Lattice Misfit
245 MPa 137 MPa
Reference Example 1 209.35 105.67 -0.39
Reference Example 2 283.20 158.75 -0.40
Reference Example 3 219.37 135.85 -0.56
Reference Example 4 274.38 153.15 -0.58
Reference Example 5 328.00 487.75 -0.58
Reference Example 6 203.15 -0.41
Embodiment 1 509.95 326.50 -0.60
Embodiment 2 420.60 753.95 -0.42
Embodiment 3 1062.50 -0.62
Embodiment 4 966.00 -0.44
Embodiment 5 1256.00 -0.48
Embodiment 6 400.00 -0.45
Embodiment 7 1254.00 -0.60
Embodiment 8 682.00 -0.63
Embodiment 9 550.00 -0.42
Embodiment 10 658.50 -0.45
Embodiment 11 622.00 -0.48
Embodiment 12 683.50 -0.51
Embodiment 13 412.7 766.35 -0.62
Embodiment 14 1524.00 -0.45
Table 4
Sample (alloy name) Lattice Misfit
Comparative Example 1(CMSX 2 -0.26
Comparative Example 2 (CMSX-4) -0.14
Comparative Exam le 3 (ReneN -0.22
Comparative Example 4 (CMSX 1 OK) -0.14
Comparative Example 5 (3B) -0.25
As is clear from Table 3, the samples of the reference examples 1-6 and
embodiments 1-14
CA 02508698 2011-06-03
17
were determined to have high strength even under high temperature conditions
of 1273K (1000 C).
In particular, reference example 5 having a composition of 4.0 wt% of Ru,
embodiments 1,2,4,9,
and 11 having a composition approximately 5.0 wt% of Ru, embodiments 3, 12 and
13 having a
composition of 6.0 wt% of Ru, and embodiment 14 having a composition of 7.0
wt% of Ru, were
determined to have high strength at high temperature.
Furthermore, as is=clear from Tables 3 and 4, the lattice misfit of
comparative examples
were -0.35 and more, whereas those of reference examples 1-6 and embodiments 1-
14 were -0.35
or less.
In addition, the creep rupture characteristics (withstand temperate e) were
compared for the
alloys of the prior art shown in Table 1(Comparative Examples 1 through 5) and
the sample
shown in Table 2 (reference examples 1-6 and embodiments 1-14). The result
thereof is
disclosed in Table 5. Creep rupture characteristics were determined either as
a result of measuring
the temperature until the sample ruptured under conditions of applying stress
of 137 MPa for 1000
hours, or converting the rupture temperature of the sample under those
conditions.
CA 02508698 2011-06-03
18
Table 5
Sample (alloy name) Withstand temperature ( C)
Reference Example 1 1315K (1042 C)
Reference Example 2 1325K 1052 C)
Reference Example 3 1321K (1048 C)
Reference Example 4 1324K (1051 C)
Reference Example 5 1354K (1081 C)
Reference Example 6 1332K (1059 C)
Embodiment 1 1344K (1071 C)
Embodiment 2 1366K 1093 C)
Embodiment 3 1375K (1 102C)
Embodiment 4 1372K 1099 C)
Embodiment 5 1379K (1106 C)
Embodiment 6 1379K (1106 C)
Embodiment 7 1379K (1106 C)
Embodiment 8 1363K 1090 C)
Embodiment 9 1358K (1085 C)
Embodiment 10 1362K (1089 C)
Embodiment 11 1361K (1088 C)
Embodiment 12 1363K (1090 C)
Embodiment 13 1366K (1093 C)
Embodiment 14 1384K (1111 C)
Comparative Example 1(CMSX-2) 1289K (1016 C)
Comparative Example 2 (CMSX-4) 1306K (1033 C)
Comparative Example 3 (ReneN6) 1320K (1047 C)
Comparative Example 4 (CMSX-1 OK) 1345K (1072 C)
Comparative Example 5 (3B) 1353K (1080 C)
(Converted to 137 MPa,1000 hours)
As is clear from Table 5, the samples of reference examples 1-6 and
embodiments 1-14
were determined to have a high withstand temperature (1356K (1083 C)) equal to
or greater than
the alloys of the prior art (comparative Examples 1-5). In particular, samples
of reference
examples 1-6 and embodiments 1-14 were determined to have a high withstand
temperature
(embodiment 1: 1344K (1071 C), embodiment 2:1366K (1093 C), embodiment 3:1375K
(1102 C), embodiment 4:1372K (1099 C), embodiment 5:1379K (1106 C), embodiment
6:
1379K (1106 C), embodiment 7:1379K (1106 C), embodiment 8:1363K (1090 C),
embodiment
9:1358K (1085 C), embodiment 10: 1362K (1089 C), embodiment 11: 1361K (1088
C),
CA 02508698 2011-06-03
19
embodiment 12:1363K (1090 C), embodiment 13:1366K (1093 C) and embodiment
14:1384K
(1111 C)).
Thus, this alloy has a higher heat resistance temperature than Ni-based single
crystal super
alloys of the prior art, and was determined to have high strength even at high
temperatures.
Furthermore, in the Ni-based single crystal super alloy, if the composite
ratio of Ru
excessively increases, the s phase precipitates and strength at high
temperatures deceases.
Therefore, the composite ratio of Ru is preferably be determined to a range so
as to keep the
balance against the composition of the other composite elements is suitably
adjusted (4.1 wt% or
more to 14.0 wt% or less, for example).
