EP0709477B1 - Heat-resistant nickel-based alloy excellent in weldability - Google Patents
Heat-resistant nickel-based alloy excellent in weldability Download PDFInfo
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- EP0709477B1 EP0709477B1 EP19950114242 EP95114242A EP0709477B1 EP 0709477 B1 EP0709477 B1 EP 0709477B1 EP 19950114242 EP19950114242 EP 19950114242 EP 95114242 A EP95114242 A EP 95114242A EP 0709477 B1 EP0709477 B1 EP 0709477B1
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- alloy
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
- C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
- C22C19/056—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 10% but less than 20%
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- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C19/00—Alloys based on nickel or cobalt
- C22C19/03—Alloys based on nickel or cobalt based on nickel
- C22C19/05—Alloys based on nickel or cobalt based on nickel with chromium
- C22C19/051—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W
- C22C19/055—Alloys based on nickel or cobalt based on nickel with chromium and Mo or W with the maximum Cr content being at least 20% but less than 30%
Definitions
- the present invention relates to a heat-resistant nickel-based alloy that can be used as a material for forming the stationary turbine vane of a gas turbine and other parts to be exposed to high temperatures.
- Heat-resistant alloys heretofore used as materials for parts to be exposed to high temperatures include an Ni-based alloy which enjoys the effects of both strengthening through precipitation of an intermetallic compound Ni 3 (Al,Ti), i.e., a ⁇ ' phase, and strengthening through solid solution with Mo, W, etc., and a Co-based alloy strengthened through precipitation of a carbide.
- an increase in the amount of precipitation of the ⁇ ' phase generally tends to lower the weldability of the alloy, though it improves the high-temperature strength of the alloy.
- this is clear from the fact that an alloy increased in the amount of precipitation of the ⁇ ' phase to improve the high-temperature strength thereof (Japanese Patent Publication No. 6,968/1979) is very poor in weldability, while an alloy decreased in the amount of precipitation of the ⁇ ' phase to improve the weldability thereof (Japanese Patent Laid-Open No. 104,738/1989) is very low in high-temperature strength.
- the Co-based alloy though generally good in weldability, is low in high-temperature strength, and no remarkable improvement can be expected.
- the Ni-based alloy since the high-temperature strength of the Co-based alloy is limited, the Ni-based alloy must be improved in weldability without detriment to the high-temperature strength thereof.
- the contents of ⁇ ' phase-forming elements such as Al and Ti should not be lowered, but the contents of other elements such as W, C, and Zr must be adjusted for the desired purpose of obtaining an alloy which can be used to produce, for example, welded structures to be used at high temperatures, such as the stationary vane of a gas turbine and apparatuses having a welded structure.
- the performance of such an alloy is characterized by a creep rupture life of at least 110 hours as measured under 20 kgf/mm 2 at 900°C, and a maximum crack length of at most 0.8 mm, as measured using 5x60x100 mm test pieces TIG-welded to each other under welding conditions involving a welding current of 100 A, a welding voltage of 12 V and a welding speed of 1.67 mm/sec, according to a varestraint test wherein the added strain (total strain) is 0.25% or 0.77%.
- an alloy having an excellent high-temperature strength and a good weldability can be obtained by increasing the high-temperature strength through the addition of Cr and Co within such respective ranges of contents as not to form deleterious phases such as a ⁇ phase and a ⁇ phase and further addition of ⁇ ' phase-forming elements such as Al, Ti, Nb and Ta as well as solid solution strengthening elements, such as W and Mo, while at the same time improving the weldability through the addition of suitable amounts of C, Zr and B liable to segregate in grain boundaries, as corresponds to an alloy composition which will be described later; and that a Ni-based alloy usable as a material for parts to be exposed to high temperatures and used in a low-grade fuel such as heavy oil, i.e., having excellent oxidation resistance and corrosion resistance as well, can be prepared.
- the present invention has been completed based on these findings.
- Fig. 1 is a diagram showing the scope of the alloy of the present invention and the test results with respect to creep rupture life.
- Fig. 2 is a diagram showing a comparison of alloys tested for creep rupture strength.
- Fig. 3 is a diagram showing the relationship between the maximum varestraint crack length and the creep rupture life.
- Fig. 4 is a perspective view of the stationary vane of a gas turbine produced using the alloy of the present invention and subjected to a weldability test.
- Fig. 5 is an illustration of the welded portion in the weldability test.
- Figs. 6A and 6B are illustrations of the essentials of the varestraint test carried out for the evaluation of weldabilities of alloys according to the present invention and comparative alloys.
- C forms a carbide which precipitates particularly in crystal grain boundaries and in dendrite boundaries to strengthen the grain boundaries and the dendrite boundaries.
- the strengthening effect thereof is negligible.
- it exceeds 0.25% the ductility and creep strength of the alloy are lowered. It is especially preferably in the range of 0.09 to 0.23%.
- the Cr content is specified to be 18 to 25% in the foregoing nickel-based nickel alloy (1) of the first type and 10 to less than 20% in the nickel-based nickel alloy (2) of the second type.
- Cr is an element capable of imparting an oxidation resistance and a corrosion resistance at high temperatures to the alloy.
- the Cr content is lower than the above-specified lower limits, the effect thereof is poor.
- it exceeds the above-specified upper limits it involves a fear of forming the a phase when the alloy is used at a high temperature for a long period of time.
- the nickel-based nickel alloy (1) is provided having particular regard to the corrosion resistance and oxidation resistance thereof, while the nickel-based nickel alloy (2) is provided having particular regard to the high-temperature strength thereof.
- Co has a function of increasing the limit of solid solution (solid solution limit) of ⁇ ' phase-forming elements such as Ti and Al into the matrix at a high temperature.
- ⁇ ' phase-forming elements such as Ti and Al into the matrix at a high temperature.
- a Co content'of at least 15.0% must be adopted.
- the Co content is specified to be at most 25.0% in order to avoid a fear of forming the ⁇ phase.
- Ti is an element required for precipitation of the ⁇ ' phase to increase the high-temperature strength of the alloy.
- the Ti content is lower than 1.0%, the desired strength cannot be secured.
- it is specified to be at most 5.0% because too much addition of Ti spoils the ductility and weldability of the alloy.
- Al forms the ⁇ ' phase like Ti to increase the high-temperature strength of the alloy while contributing to imparting to the alloy an oxidation resistance and a corrosion resistance at high temperatures.
- the Al content must be at least 1.0%, while it is specified to be atmost 4.0% because too much addition of Al spoils the ductility and weldability of the alloy.
- the (Al + Ti) content is especially preferably in the range of 3.0 to 7.0%.
- W and Mo have a function of solid solution strengthening and,weak precipitation strengthening to contribute to imparting a high-temperature strength to the alloy.
- the (W + 1/2Mo) content must be at least 0.5%. Since too much addition of these elements spoils the ductility of the alloy, the W content, the Mo content, and the (W + 1/2Mo) content are specified to be at most 10%, at most 3.5%, and at most 10%, respectively.
- Ta and Nb contribute to an improvement in high-temperature strength through solid solution strengthening and ⁇ ' phase precipitation strengthening. This effect is exhibited when the Ta content is at least 0.5% and when the Nb content is at least 0.2%. On the other hand, since too much addition of these elements lowers the ductility of the alloy, the Ta content and the Nb content are specified to be at most 4.5% and at most 3.0%, respectively.
- the Ta content and the Nb content are especially preferably in the range of 1.0 to 4.2% and in the range of 0.5 to 1.5%, respectively.
- Zr exhibits the effect of increasing the bonding strength in crystal grain boundaries to strengthen the grain boundaries.
- the Zr content is lower than 0.005%, no improvement in creep strength can be observed.
- it exceeds 0.10% the weldability of the alloy is unfavorably lowered.
- it must be in the range of 0.005 to 0.10%, and is especially preferably in the range of 0.01 to 0.10%.
- B increases the bonding strength in crystal grain boundaries like Zr to strengthen the grain boundaries.
- the B content is lower than 0.001%, no improvement in creep strength can be observed.
- it exceeds 0.01% the weldability of the alloy is unfavorably lowered.
- the B content is specified to be in the range of 0.001 to 0.01%.
- the lower limit of the (Al + Ti) content is specified, with taking also into account the Cr content, to be at least 4%, as shown in the same figure.
- W and Mo have a function of solid solution strengthening and carbide precipitation strengthening to exhibit the effect of increasing the high-temperature strength of the alloy.
- the (W + 1/2Mo) content must be at least 0.5%.
- the upper limit of the (W + 1/2Mo) content is specified to be 10%.
- Table 1 shows the chemical compositions (by wt.%) of representative alloys invented for the stationary vane of a gas turbine.
- Table 2 shows the chemical compositions of comparative alloys as conventional alloys.
- Each composition was melted in a high-frequency vacuum melting furnace to prepare 20 kg of an ingot. This sample was precision-cast as the master ingot according to a lost wax process, and then heat-treated at 1,160°C for 4 hours, at 1,000°C for 6 hours, and at 800°C for 4 hours. Thereafter, it was machined into creep rupture test pieces of 6.25 mm ⁇ x 25 mm in parallel portion size, 5x60x100 mm varestraint test pieces, etc. Alloys Nos.
- alloys according to the present invention are alloys according to the present invention, while Alloys Nos. X, Y, Z, and 19 to 36 are comparative alloys. Additionally stated, the Alloys Nos. X and Y are examples of the aforementioned alloy of Japanese Patent Publication No. 6,968/1979, while the Alloy No. Z is an example of the aforementioned alloy of Japanese Patent Laid-Open No. 104,738/1989.
- Fig. 1 shows the relationship between the (Al + Ti) content and the (W + 1/2Mo) content for every sample as well as the creep rupture life under 20 kgf/mm 2 at 900°C in ( ) accompanying every sample No. Additionally stated, in Fig. 1, the alloys according to the present invention were indicated by the open symbol (o), while the comparative alloys are indicated by the solid symbol ( ⁇ ).
- Fig. 2 shows a comparison of Alloys Nos. 9 and 11 of the present invention in Table 1 with the Comparative Alloys Nos. Y, Z, and 20 in Table 2 with respect to creep rupture strength under 20 kgf/mm 2 at 900°C and under 10 kgf/mm 2 at 980°C.
- the test results at 900°C and 980°C correspond to the points of 20 kgf/mm 2 and 10 kgf/mm 2 , respectively, in terms of stress represented by the ordinate.
- Alloys Nos. 9 and 11 of the present invention are higher in Larson-Miller parameter under the same test stress than the Comparative Alloys Nos. Y, Z, and 20. This is the effect of increasing the (Al + Ti) content and the (W + 1/2Mo) content while decreasing the Cr content (No.11).
- Comparative Alloy No. Y which was slightly higher in (Al + Ti) content than the Alloy No. 9 and high also in Cr content
- Comparative Alloy No. 20 which was low in (Al + Ti) content but high in (W + 1/2Mo) content
- Comparative Alloy No. Z which was low in both of (Al + Ti) content and (W + 1/2Mo) content, etc. are lower in Larson-Miller parameter under the same test stress than the alloys of the present invention.
- varestraint test As shown in Figs. 6A and 6B.
- reference numerals are as follows: 12: varestraint test piece (before application of flexural stress), 13: yoke, 14: bead, 15: welding torch, 16: varestraint test piece (after application of flexural stress), and 17: bending block.
- test pieces were TIG-welded to each other under welding conditions involving a welding current of 100 A, a welding voltage of 12 V, and a welding speed of 1.67 mm/sec, and then loaded with a total strain of 0.25% or 0.77%.
- the resulting maximum crack length as a yardstick of the zone turned brittle when welded was measured.
- Fig. 3 shows the relationship between the maximum crack length and the creep rupture life (900°C x 20 kgf/mm 2 ).
- the ordinate in the same figure demonstrates that the smaller the maximum crack length, the better the weldability. Accordingly, as the point is located at the furtherest right side and at the lower side, the alloy is higher in high-temperature strength and better in weldability, respectively.
- the resulting product was subjected to a solution heat treatment at 1,160°C for 4 hours, and then subjected to a weldability test.
- the stationary vane had a profile portion width of about 200 mm and a height of about 200 mm, and was a cast article having a hollow structure provided with an internal air path for cooling the same.
- build-up welding, or padding was carried out in ventral places 1, 2, 3, and 4 of a vane portion, places 5 and 6 of the leading edge, and a place 7 of the trailing edge.
- Reference numeral 9 represents an outer shroud.
- the shroud portion 8 (Alloy No. 11 of the present invention) of the inner shroud 8 was welded with a cover plate 10 (Hastelloy X alloy) with a fillet welding of Hastelloy W alloy 11 according to the TIG welding method.
- a visual inspection, a fluorescence penetrant inspection, an observation of the microstructure of the cross section at the position as shown in Fig. 5, etc. were carried out and no cracks were found in any place.
- substantially the same stationary vane of a gas turbine as described above was produced using Comparative Alloy No. Y (Japanese Patent Publication No.
- a heat-resistant Ni-based alloy can be obtained, which has a higher high-temperature strength and a better weldability than conventional heat-resistant Ni-based alloys.
- This heat-resistant Ni-based alloy is especially suitable as a material for the stationary vane of a gas turbine required to be reliable in keeping with an increase in the service temperature of the gas turbine.
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Description
The present invention relates to a heat-resistant
nickel-based alloy that can be used as a
material for forming the stationary turbine vane of a
gas turbine and other parts to be exposed to high
temperatures.
Heat-resistant alloys heretofore
used as materials for parts to be exposed to high
temperatures, such as the stationary turbine vane of a
gas turbine, include an Ni-based alloy which enjoys the effects of both
strengthening through precipitation of an
intermetallic compound Ni3(Al,Ti), i.e., a γ' phase,
and strengthening through solid solution with Mo, W,
etc., and a Co-based alloy strengthened through
precipitation of a carbide.
In the Ni-based alloy, an increase in the amount
of precipitation of the γ' phase generally tends to
lower the weldability of the alloy, though it improves
the high-temperature strength of the alloy. For
example, this is clear from the fact that an alloy
increased in the amount of precipitation of the γ'
phase to improve the high-temperature strength thereof
(Japanese Patent Publication No. 6,968/1979) is very
poor in weldability, while an alloy decreased in the
amount of precipitation of the γ' phase to improve the
weldability thereof (Japanese Patent Laid-Open No.
104,738/1989) is very low in high-temperature
strength. On the other hand, the Co-based alloy,
though generally good in weldability, is low in high-temperature
strength, and no remarkable
improvement can be expected.
As is apparent from the foregoing, since the
high-temperature strength of the Co-based alloy is
limited, the Ni-based alloy must be improved in
weldability without detriment to the high-temperature
strength thereof.
In order to improve the weldability of the Ni-based
alloy without detriment to the high-temperature
strength thereof, the contents of γ' phase-forming
elements such as Al and Ti should not be lowered, but
the contents of other elements such as W, C, and Zr
must be adjusted for the desired purpose of obtaining
an alloy which can be used to produce, for example,
welded structures to be used at high temperatures,
such as the stationary vane of a gas turbine and apparatuses
having a welded structure. The performance of
such an alloy is characterized by a creep rupture life
of at least 110 hours as measured under 20 kgf/mm2 at
900°C, and a maximum crack length of at most 0.8 mm, as
measured using 5x60x100 mm test pieces TIG-welded to
each other under welding conditions involving a
welding current of 100 A, a welding voltage of 12 V
and a welding speed of 1.67 mm/sec, according to a
varestraint test wherein the added strain (total
strain) is 0.25% or 0.77%.
As a result of intensive investigations, the
inventors of the present invention have found out that
an alloy having an excellent high-temperature strength
and a good weldability can be obtained by increasing
the high-temperature strength through the addition of Cr
and Co within such respective ranges of contents as
not to form deleterious phases such as a σ phase and a
µ phase and further addition of γ' phase-forming
elements such as Al, Ti, Nb and Ta as well as solid
solution strengthening elements, such as W and Mo, while
at the same time improving the weldability through the
addition of suitable amounts of C, Zr and B liable to
segregate in grain boundaries, as corresponds to an
alloy composition which will be described later; and
that a Ni-based alloy usable as a material for
parts to be exposed to high temperatures and used in a
low-grade fuel such as heavy oil, i.e., having excellent
oxidation resistance and corrosion resistance as well,
can be prepared. The present invention has been
completed based on these findings.
Specifically, in accordance with the present
invention, there are provided:
Fig. 1 is a diagram showing the scope of the
alloy of the present invention and the test results
with respect to creep rupture life.
Fig. 2 is a diagram showing a comparison of
alloys tested for creep rupture strength.
Fig. 3 is a diagram showing the relationship
between the maximum varestraint crack length and the
creep rupture life.
Fig. 4 is a perspective view of the stationary
vane of a gas turbine produced using the alloy of the
present invention and subjected to a weldability test.
Fig. 5 is an illustration of the welded portion
in the weldability test.
Figs. 6A and 6B are illustrations of the
essentials of the varestraint test carried out for the
evaluation of weldabilities of alloys according to the
present invention and comparative alloys.
The functions of elements in the alloy
composition of the heat-resistant Ni-based alloy of
the present invention will now be described together
with the reasons for specifying the contents (by
weight) of the elements added thereto.
C forms a carbide which precipitates
particularly in crystal grain boundaries and in
dendrite boundaries to strengthen the grain boundaries
and the dendrite boundaries. When the C content is
lower than 0.05%, the strengthening effect thereof is
negligible. When it exceeds 0.25%, the ductility and creep
strength of the alloy are lowered. It is especially
preferably in the range of 0.09 to 0.23%.
The Cr content is specified to be 18 to 25% in
the foregoing nickel-based nickel alloy (1) of the
first type and 10 to less than 20% in the nickel-based nickel
alloy (2) of the second type. Cr is an element
capable of imparting an oxidation resistance and a
corrosion resistance at high temperatures to the
alloy. When the Cr content is lower than the above-specified
lower limits, the effect thereof is poor.
When it exceeds the above-specified upper limits, it
involves a fear of forming the a phase when the alloy
is used at a high temperature for a long period of
time. Additionally stated, the nickel-based nickel
alloy (1) ) is provided having particular regard to the
corrosion resistance and oxidation resistance thereof,
while the nickel-based nickel alloy (2) is provided
having particular regard to the high-temperature
strength thereof.
Co has a function of increasing the limit of
solid solution (solid solution limit) of γ' phase-forming
elements such as Ti and Al into the matrix at
a high temperature. With the Al and Ti contents of
the alloy according to the present invention, a Co
content'of at least 15.0% must be adopted. On the
other hand, the Co content is specified to be at most
25.0% in order to avoid a fear of forming the σ phase.
Ti is an element required for precipitation of
the γ' phase to increase the high-temperature strength
of the alloy. When the Ti content is lower than 1.0%,
the desired strength cannot be secured. On the other
hand, it is specified to be at most 5.0% because too
much addition of Ti spoils the ductility and
weldability of the alloy.
Al forms the γ' phase like Ti to increase the
high-temperature strength of the alloy while
contributing to imparting to the alloy an
oxidation resistance and a corrosion resistance at
high temperatures. The Al content must be at least
1.0%, while it is specified to be atmost 4.0% because
too much addition of Al spoils the ductility and
weldability of the alloy. The (Al + Ti) content is
especially preferably in the range of 3.0 to 7.0%.
W and Mo have a function of solid solution
strengthening and,weak precipitation strengthening to
contribute to imparting a high-temperature
strength to the alloy. In order to secure the
foregoing effect, the (W + 1/2Mo) content must be at
least 0.5%. Since too much addition of these elements
spoils the ductility of the alloy, the W content, the
Mo content, and the (W + 1/2Mo) content are specified
to be at most 10%, at most 3.5%, and at most 10%,
respectively.
Ta and Nb contribute to an improvement in high-temperature
strength through solid solution
strengthening and γ' phase precipitation
strengthening. This effect is exhibited when the Ta
content is at least 0.5% and when the Nb content is at
least 0.2%. On the other hand, since too much
addition of these elements lowers the ductility of the
alloy, the Ta content and the Nb content are specified
to be at most 4.5% and at most 3.0%, respectively.
The Ta content and the Nb content are especially
preferably in the range of 1.0 to 4.2% and in the
range of 0.5 to 1.5%, respectively.
Zr exhibits the effect of increasing the bonding
strength in crystal grain boundaries to strengthen the
grain boundaries. When the Zr content is lower than
0.005%, no improvement in creep strength can be
observed. On the other hand, when it exceeds 0.10%,
the weldability of the alloy is unfavorably lowered.
Thus, it must be in the range of 0.005 to 0.10%, and
is especially preferably in the range of 0.01 to
0.10%.
B increases the bonding strength in crystal
grain boundaries like Zr to strengthen the grain
boundaries. When the B content is lower than 0.001%,
no improvement in creep strength can be observed. On
the other hand, when it exceeds 0.01%, the weldability
of the alloy is unfavorably lowered. Thus, the B
content is specified to be in the range of 0.001 to
0.01%.
The reasons why limitations are made within the
ranges surrounded by the lines in Fig. 1 are as
follows. Al and Ti precipitate the γ' phase, i.e.,
Ni3(Al, Ti), as a factor in strengthening the Ni-based
alloy to increase the high-temperature strength
thereof. Since too much addition of these elements
lowers the weldability and ductility of the alloy,
however, the (Al + Ti) content is specified to be at
most 7%. When it is too low, the effect of increasing
the high-temperature strength of the alloy is
decreased. Thus, it is specified to be at least 3% as
shown in the same figure. Additionally stated, since
the Cr content also exerts an influence on the high-temperature
strength of the alloy, the lower limit of
the (Al + Ti) content is specified, with taking also
into account the Cr content, to be at least 4%, as
shown in the same figure. W and Mo have a function of
solid solution strengthening and carbide precipitation
strengthening to exhibit the effect of increasing the
high-temperature strength of the alloy. In order to
secure this effect, the (W + 1/2Mo) content must be at
least 0.5%. On the other hand, since too much
addition of these elements fosters precipitation of
deleterious phases such as the a phase to lower the
ductility and strength of the alloy, the upper limit
of the (W + 1/2Mo) content is specified to be 10%.
The following specific Examples will illustrate
the present invention in more detail.
Table 1 shows the chemical compositions (by
wt.%) of representative alloys invented for the
stationary vane of a gas turbine. On the other hand,
Table 2 shows the chemical compositions of comparative
alloys as conventional alloys. Each composition was
melted in a high-frequency vacuum melting furnace to
prepare 20 kg of an ingot. This sample was
precision-cast as the master ingot according to a lost
wax process, and then heat-treated at 1,160°C for 4
hours, at 1,000°C for 6 hours, and at 800°C for 4
hours. Thereafter, it was machined into creep rupture
test pieces of 6.25 mm x 25 mm in parallel portion
size, 5x60x100 mm varestraint test pieces, etc.
Alloys Nos. 1 to 18 in Table 1 are alloys according to
the present invention, while Alloys Nos. X, Y, Z, and
19 to 36 are comparative alloys. Additionally stated,
the Alloys Nos. X and Y are examples of the
aforementioned alloy of Japanese Patent Publication
No. 6,968/1979, while the Alloy No. Z is an example of
the aforementioned alloy of Japanese Patent Laid-Open
No. 104,738/1989.
Fig. 1 shows the relationship between the (Al +
Ti) content and the (W + 1/2Mo) content for every
sample as well as the creep rupture life under 20
kgf/mm2 at 900°C in ( ) accompanying every sample No.
Additionally stated, in Fig. 1, the alloys according
to the present invention were indicated by the open
symbol (o), while the comparative alloys are indicated
by the solid symbol ().
Alloys of the present invention with high (Al +
Ti) and (W + 1/2Mo) contents which are in the range
surrounded by the line connecting points A, B, C,
D, and E (1, 4, 11, 12, 13, 14, 15, and 16) all
exhibit a high strength, and Alloy No. 11 in
particular exhibits an especially high strength.
Alloys of the present invention with a low Cr content
and with (Al + Ti) and (W + 1/2Mo) contents which are
in the range surrounded by the line connecting
points F, G, H, and D (2, 3, 5, 6, 7, 8, 9, 10, 17,
and 18) exhibit an especially high strength.
Fig. 2 shows a comparison of Alloys Nos. 9
and 11 of the present invention in Table 1 with the
Comparative Alloys Nos. Y, Z, and 20 in Table 2 with
respect to creep rupture strength under 20 kgf/mm2 at
900°C and under 10 kgf/mm2 at 980°C. The abscissa
represents the Larson-Miller parameter: P = Tkx:(20 +
log t)x10-3 [Tk : test temperature (°K), t: rupture
life (hr)]. The test results at 900°C and 980°C
correspond to the points of 20 kgf/mm2 and 10 kgf/mm2,
respectively, in terms of stress represented by the
ordinate. It is demonstrated that the higher the
parameter P in the abscissa, the higher the strength.
Alloys Nos. 9 and 11 of the present invention are
higher in Larson-Miller parameter under the same test
stress than the Comparative Alloys Nos. Y, Z, and 20.
This is the effect of increasing the (Al + Ti) content
and the (W + 1/2Mo) content while decreasing the Cr
content (No.11). On the other hand, Comparative
Alloy No. Y which was slightly higher in (Al + Ti) content than
the Alloy No. 9 and high also in Cr content,
Comparative Alloy No. 20 which was low in (Al + Ti) content but
high in (W + 1/2Mo) content, Comparative Alloy No.
Z which was low in both of (Al + Ti) content and (W + 1/2Mo)
content, etc. are lower in Larson-Miller parameter
under the same test stress than the alloys of the
present invention.
The weldability was evaluated according to a
varestraint test, as shown in Figs. 6A and 6B. In the
figures, reference numerals are as follows:
12: varestraint test piece (before application of
flexural stress), 13: yoke, 14: bead, 15: welding
torch, 16: varestraint test piece (after application
of flexural stress), and 17: bending block.
Specifically, test pieces were TIG-welded to
each other under welding conditions involving a
welding current of 100 A, a welding voltage of 12 V,
and a welding speed of 1.67 mm/sec, and then loaded
with a total strain of 0.25% or 0.77%. The resulting
maximum crack length as a yardstick of the zone turned
brittle when welded was measured. Fig. 3 shows the
relationship between the maximum crack length and the
creep rupture life (900°C x 20 kgf/mm2). The ordinate
in the same figure demonstrates that the smaller the
maximum crack length, the better the weldability.
Accordingly, as the point is located at the furtherest right
side and at the lower side, the alloy is higher in
high-temperature strength and better in weldability,
respectively. The Alloys Nos. 3, 7, 9, 10, 11, 12,
and 15 with a Zr content of at most 0.1% and a B
content of at most 0.01 according to the present
invention are all small in maximum crack length in the
varestraint test. Alloys Nos. 9, 11, and 12 in
particular showed a maximum crack length of at most
0.3 mm as the target and a creep rupture life of at
least 185 hours, and hence have excellent properties.
On the other hand, Comparative Alloys Nos. X, Y,
25, 27, 28, 33, and 35 all showed a maximum crack
length in the varestraint test of at least 0.8 mm to
miss the target, though they showed a creep rupture
life of at least 110 hours. As is apparent from the
foregoing results, a good weldability and a high creep
strength can be secured either if the relationship
between the (Al + Ti) content and the (W + 1/2Mo)
content are specified to be in the range A-B-C-D-E
even though the Zr content and the B content are
lowered, or if the relationship between the (Al + Ti)
content and the (W + 1/2Mo) content are specified to
be in the range A-B-C-F-G-H-E with a decrease in Cr
content.
Alloy No. 11 of Example 1, as shown in Table
1, was used to produce the stationary vane of a gas
turbine as shown in Fig. 4 according to the lost wax
precision-casting process. The resulting product was
subjected to a solution heat treatment at 1,160°C for
4 hours, and then subjected to a weldability test.
The stationary vane had a profile portion width of
about 200 mm and a height of about 200 mm, and was a
cast article having a hollow structure provided with
an internal air path for cooling the same. As shown
in Fig. 4, build-up welding, or padding, was carried
out in ventral places 1, 2, 3, and 4 of a vane
portion, places 5 and 6 of the leading edge, and a
place 7 of the trailing edge. Reference numeral 9
represents an outer shroud. As shown in Fig. 5, the
shroud portion 8 (Alloy No. 11 of the present
invention) of the inner shroud 8 was welded with a
cover plate 10 (Hastelloy X alloy) with a fillet
welding of Hastelloy W alloy 11 according to the TIG
welding method. After the welding, a visual
inspection, a fluorescence penetrant inspection, an
observation of the microstructure of the cross section
at the position as shown in Fig. 5, etc. were carried
out and no cracks were found in any place. Additionally
stated, substantially the same stationary vane of a
gas turbine as described above was produced using
Comparative Alloy No. Y (Japanese Patent Publication
No. 6,968/1979), and subjected to a weldability test.
As a result, many cracks were found by a
fluorescence penetrant inspection, while cracks of
about 1 mm in length were found by an observation
of the microstructure of the cross section.
As described hereinbefore, according to the
present invention, a heat-resistant Ni-based alloy can
be obtained, which has a higher high-temperature
strength and a better weldability than conventional
heat-resistant Ni-based alloys. This heat-resistant
Ni-based alloy is especially suitable as a material
for the stationary vane of a gas turbine required to
be reliable in keeping with an increase in the service
temperature of the gas turbine.
Claims (2)
- A heat-resistant nickel-based alloy excellent in weldability, said nickel-based alloy comprising, in terms of wt. %, 0.05 to 0.25% of C, 18 to 25% of Cr, 15 to 25% of Co, up to 3.5% of Mo and 5 to 10% of W with the proviso to exclude an amount of W of 5 % with the content of one or both of Mo and W being 5 to 10% in terms of W + 1/2Mo excluding an amount of 5 % of W + 1/2Mo, 1.0 to 5.0% of Ti, 1.0 to 4.0% of Al, 0.5 to 4.5% of Ta, 0.2 to 3.0% of Nb, 0.005 to 0.10% of Zr, 0.001 to 0.01% of B, and the balance of Ni and unavoidable impurity elements, wherein the (Al + Ti) content and the (W + 1/2Mo) content are within the range surrounded by the line connecting the point A (Al + Ti: 3%, W + 1/2Mo: 10%), the point B (Al + Ti: 5%, W + 1/2Mo: 7.5%), the point C (Al + Ti: 5%, W + 1/2Mo: 5%), the point D (Al + Ti: 7%, W + 1/2Mo: 5%), and the point E (Al + Ti: 7%, W + 1/2Mo: 10%) in this sequence in Fig. 1.
- A heat-resistant nickel-based alloy excellent in weldability, said nickel-based alloy comprising, in terms of wt. %, 0.05 to 0.25% of C, 10 to 20% of Cr with the proviso that the amount of Cr of 20 % is excluded, 15 to 25% of Co, up to 3.5% of Mo and 0.5 to 10% of W with the content of one or both of Mo and W being 0.5 to 10% in terms of W + 1/2Mo, 1.0 to 5.0% of Ti, 1.0 to 4.0% of Al, 0.5 to 4.5% of Ta, 0.2 to 3.0% of Nb, 0.005 to 0.10% of Zr, 0.001 to 0.01% of B, and the balance of Ni and unavoidable impurity elements, wherein the (Al + Ti) content and the (W + 1/2Mo) content are within the range surrounded by the line connecting the point A (Al + Ti: 3%, W + 1/2Mo: 10%), the point B (Al + Ti: 5%, W + 1/2Mo: 7.5%), the point C (Al + Ti: 5%, W + 1/2Mo: 5%), the point F (Al + Ti: 4%, W + 1/2Mo: 5%), the point G (Al + Ti: 4%, W + 1/2Mo: 0.5%), the point H (Al + Ti: 7%, W + 1/2Mo: 0.5%), and the point E (Al + Ti: 7%, W + 1/2Mo: 10%) in this sequence in Fig. 1.
Applications Claiming Priority (2)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP6267111A JP2862487B2 (en) | 1994-10-31 | 1994-10-31 | Nickel-base heat-resistant alloy with excellent weldability |
| JP267111/94 | 1994-10-31 |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| EP0709477A1 EP0709477A1 (en) | 1996-05-01 |
| EP0709477B1 true EP0709477B1 (en) | 1998-05-27 |
Family
ID=17440220
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| EP19950114242 Expired - Lifetime EP0709477B1 (en) | 1994-10-31 | 1995-09-11 | Heat-resistant nickel-based alloy excellent in weldability |
Country Status (4)
| Country | Link |
|---|---|
| EP (1) | EP0709477B1 (en) |
| JP (1) | JP2862487B2 (en) |
| CA (1) | CA2146534C (en) |
| DE (1) | DE69502680T2 (en) |
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US12319985B2 (en) | 2019-10-02 | 2025-06-03 | Alloyed Limited | Nickel-based alloy |
Families Citing this family (18)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US6120624A (en) * | 1998-06-30 | 2000-09-19 | Howmet Research Corporation | Nickel base superalloy preweld heat treatment |
| US6284392B1 (en) * | 1999-08-11 | 2001-09-04 | Siemens Westinghouse Power Corporation | Superalloys with improved weldability for high temperature applications |
| CA2287116C (en) * | 1999-10-25 | 2003-02-18 | Mitsubishi Heavy Industries, Ltd. | Process for the heat treatment of a ni-base heat-resisting alloy |
| JP4382244B2 (en) | 2000-04-11 | 2009-12-09 | 日立金属株式会社 | Method for producing Ni-base alloy having excellent resistance to high-temperature sulfidation corrosion |
| US6696176B2 (en) * | 2002-03-06 | 2004-02-24 | Siemens Westinghouse Power Corporation | Superalloy material with improved weldability |
| JP3842717B2 (en) * | 2002-10-16 | 2006-11-08 | 株式会社日立製作所 | Welding material, welded structure, gas turbine rotor blade, and gas turbine rotor blade or stationary blade repair method |
| US6866727B1 (en) * | 2003-08-29 | 2005-03-15 | Honeywell International, Inc. | High temperature powder metallurgy superalloy with enhanced fatigue and creep resistance |
| JP4906611B2 (en) * | 2007-07-03 | 2012-03-28 | 株式会社日立製作所 | Ni-based alloy |
| JP5078537B2 (en) | 2007-10-15 | 2012-11-21 | 三菱重工業株式会社 | Repair method |
| JP5201334B2 (en) * | 2008-03-19 | 2013-06-05 | 大同特殊鋼株式会社 | Co-based alloy |
| US20110062220A1 (en) * | 2009-09-15 | 2011-03-17 | General Electric Company | Superalloy composition and method of forming a turbine engine component |
| JP5657964B2 (en) * | 2009-09-15 | 2015-01-21 | 三菱日立パワーシステムズ株式会社 | High-strength Ni-base forged superalloy and manufacturing method thereof |
| JP5296046B2 (en) | 2010-12-28 | 2013-09-25 | 株式会社日立製作所 | Ni-based alloy and turbine moving / stator blade of gas turbine using the same |
| US20160326613A1 (en) * | 2015-05-07 | 2016-11-10 | General Electric Company | Article and method for forming an article |
| GB2565063B (en) | 2017-07-28 | 2020-05-27 | Oxmet Tech Limited | A nickel-based alloy |
| EP3572541B1 (en) | 2018-05-23 | 2023-05-17 | Rolls-Royce plc | Nickel-base superalloy |
| GB2584654B (en) | 2019-06-07 | 2022-10-12 | Alloyed Ltd | A nickel-based alloy |
| CN112981186B (en) * | 2021-04-22 | 2021-08-24 | 北京钢研高纳科技股份有限公司 | Low stacking fault energy superalloys, structural parts and their applications |
Family Cites Families (6)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| GB1367661A (en) * | 1971-04-07 | 1974-09-18 | Int Nickel Ltd | Nickel-chromium-cobalt alloys |
| US4039330A (en) * | 1971-04-07 | 1977-08-02 | The International Nickel Company, Inc. | Nickel-chromium-cobalt alloys |
| JPS546968A (en) | 1977-06-13 | 1979-01-19 | Unitika Ltd | Sewing process |
| US4810467A (en) | 1987-08-06 | 1989-03-07 | General Electric Company | Nickel-base alloy |
| JP2778705B2 (en) * | 1988-09-30 | 1998-07-23 | 日立金属株式会社 | Ni-based super heat-resistant alloy and method for producing the same |
| EP0561179A3 (en) * | 1992-03-18 | 1993-11-10 | Westinghouse Electric Corp | Gas turbine blade alloy |
-
1994
- 1994-10-31 JP JP6267111A patent/JP2862487B2/en not_active Expired - Lifetime
-
1995
- 1995-04-06 CA CA 2146534 patent/CA2146534C/en not_active Expired - Lifetime
- 1995-09-11 EP EP19950114242 patent/EP0709477B1/en not_active Expired - Lifetime
- 1995-09-11 DE DE1995602680 patent/DE69502680T2/en not_active Expired - Lifetime
Cited By (1)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US12319985B2 (en) | 2019-10-02 | 2025-06-03 | Alloyed Limited | Nickel-based alloy |
Also Published As
| Publication number | Publication date |
|---|---|
| CA2146534C (en) | 2001-10-02 |
| DE69502680D1 (en) | 1998-07-02 |
| DE69502680T2 (en) | 1998-09-24 |
| JP2862487B2 (en) | 1999-03-03 |
| CA2146534A1 (en) | 1996-05-01 |
| JPH08127833A (en) | 1996-05-21 |
| EP0709477A1 (en) | 1996-05-01 |
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