Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
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ALUMINIUM OXIDE FORMING NICKEL BASED ALLOY
The present invention relates to a nickel based alloy intended for use at high
temperatures, such as above 900 C. Specifically, the present invention
relates to
a dispersion strengthened nickel based alloy alloyed with aluminium which
enables formation of a stable aluminium oxide on the surface whereby the alloy
has a good oxidation resistance. Moreover, the present invention relates to a
powder of the nickel based alloy and to the use of the nickel based alloy.
Background art
Nickel based alloys alloyed with aluminium are used in a variety of high
temperature applications, such as in heat treatment furnaces, since they form
a
stable and protective aluminium oxide on the surface. The aluminium oxide
often
has a very good adhesion and does not tend to spall or fall off the surface.
Moreover, the aluminium oxide has a low growth rate even at high temperatures.
This type of alloys therefore often has a very good oxidation resistance.
Aluminium oxide forming nickel based alloys are known to be difficult to
manufacture, especially to hot-work. A strongly contributing factor to this is
the
intermetallic phase y' (Ni3A1) which is formed at temperatures below
approximately
900 C during slow cooling/heating, such as during heat treatments or during
hot
working. This intermetallic phase makes the alloy hard and brittle and
consequently difficult to work. The precipitation of y' also reduces the
activity of
aluminium in the alloy and thereby makes formation of the protective aluminium
oxide on the surface more difficult.
One example of an aluminium oxide forming nickel based alloy is
disclosed in US 4,882,125. The alloy comprises 27-35 % Cr, 2.5-5 % Al and 2.5-
6
% Fe. It is disclosed that high contents of aluminium reduces the toughness of
the
material and that the Al content should be at least 2.75 % in order to
generate a
good oxidation protection, but preferably not exceed 4 % in order not to
deteriorate
the ductility. The patent further teaches that high contents of Fe deteriorate
the
oxidation properties, for which reason the iron content should not exceed 6 %.
Another example of an aluminium oxide forming nickel based alloy is
disclosed in US 4,460,542. The alloy comprises 14-18 % Cr, 4-6% Al and 1.5-8 %
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Fe. This patent teaches that additions of 4-6 % of Al render superior
oxidation
properties compared to nickel based alloys which form chromium oxide on the
surface. Also in this patent it is disclosed that Fe has a negative effect on
the
oxidation properties, for which reason the iron content should be maximally 8
%.
WO 2004/067788 Al discloses yet another example of an aluminium
oxide forming nickel based alloy. In this case, the alloy comprises 15-40 %
Cr, 1.5-
7 % Al and 0.5-13% Fe. Best results are said to be accomplished when the alloy
comprises max 26.5 % Cr, max 11 % Fe and 3-6 % Al.
WO 00/34541 Al discloses a nickel based alloy comprising 19-23 % Cr, 3-
4.4 % Al and 18-22 % Fe. The alloy is intended for use at high temperatures.
WO
00/34541 Al discloses that the combination of 19-23 % Cr and 3-4 % Al is
critical
for formation of the protective A1203-Cr203 scale. The nickel based alloy is
strengthened by precipitation of 1 to 5 mole percent of granular Cr7C3 which
is said
to be accomplished by a 24 hour heat treatment. The alloy is produced by
melting
such as vacuum melting, casting and working into standard engineering shapes,
such as rod, bar etc.. This alloy shows good oxidation resistance up to 1000
C.
It is also previously known with iron based ferritic aluminium oxide forming
alloys. However, this type of alloys often has low mechanical strength at high
temperatures. Therefore, small particles are often added to increase the creep
strength of the material. This is described for example in Metals Handbook,
10th
edition, volume 2, page 943. Another problem with this type of alloys is that
their
ductility at room temperature often is very low which makes welding more
difficult.
In order to accomplish a reliable weld in a ferritic material a preheating of
the
material to be welded to at least 200 C is often required. In many cases, a
stress-
relieving anneal at 750-850 C is also required after welding.
Summary of the invention
The object of the present invention is to accomplish an alloy with excellent
oxidation resistance at high temperatures, specifically from about 900 C to at
least
about 1250 C, and which still has a good hot workability and good creep
strength.
The above identified object is achieved by a dispersion strengthened
nickel based alloy comprising in percent by weight (wt-%)
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C 0.05-0.2
Si max 1.5
Mn max 0.5
Cr 15-20
Al 4-6
Fe 15-25
Co max 1 0
N 0.03-0.15
0 max 0.5
one or more elements selected from the group consisting of Ta,
Zr, Hf, Ti and Nb 0.25-2.2
one or more elements selected from the group consisting of the
rare earth metals (REM) max 0.5,
balance Ni and normally occurring impurities.
The nickel based alloy in accordance with the present invention is austenitic
and
has a very good oxidation resistance, especially at high temperatures, such as
above 900 C. The oxidation resistance is high even at temperatures of about
1100 C. Since the present alloy forms a stable aluminium oxide on the surface,
it
can be used even at temperatures above those where chromium oxide forming
materials suffer from extensive oxidation, i.e. above approximately 1150 C.
It has been found that by adding relatively high contents of Fe to an
aluminium oxide forming nickel based alloy it is possible to reduce the
stability of
the intermetallic phase y', which in turn makes the alloy easier to
manufacture and
work. A reduced stability of y' renders a slower formation of such
precipitations for
a given cooling rate, which facilitates hot working of the alloy. This also
leads to a
reduced risk of reduced activity of Al, which in turn ensures that a stable
and
oxidation resistant aluminium oxide can be formed on the surface of the alloy.
The nickel based alloy according to the invention is more ductile at room
temperature than known ferritic aluminium oxide forming alloys. Therefore,
preheating or keeping the alloy warm before welding is unnecessary and
subsequent stress-relieving annealing can be avoided. The nickel based alloy
according to the invention consequently enables a facilitated welding
procedure
compared to ferritic aluminium oxide forming alloys.
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The nickel based alloy according to the invention is dispersion
strengthened. This is achieved by the addition of one or more elements
selected
from the group consisting of Ta, Zr, Hf, Ti and Nb. These elements form
dispersion
strengthening particles with C and/or N and optionally added 0. The dispersion
contributes to the mechanical strength and gives the alloy excellent creep
strength
even at high temperatures without impairing the hot-workability of the alloy.
The nickel based alloy is produced by means of powder metallurgy. The
powder metallurgical manufacturing process results in a rapidly solidified
material
wherein brittle phases do not have time to form and no great composition
variations are developed by segregation. A mixture of rapidly solidified
powder will
therefore render a metal body with essentially homogenous composition and an
essentially even distribution of very small dispersion particles.
A powder produced of the nickel based alloy will comprise dispersion
strengthened particles as described above, which will render a product
produced
of the powder excellent mechanical properties, especially at high
temperatures.
Furthermore, a powder of the nickel based alloy enables, in addition to
manufacturing of traditional forms such as tube, rod, wire, plate and strip,
also
manufacturing of solid components with complex geometry. Moreover, compound
materials wherein the nickel based alloy is incorporated can easily be
manufactured if desired, for example in order to produce a final product with
a first
load-bearing component and with a second corrosion resistant component.
The nickel based alloy according to the invention is especially suitable for
use at high temperatures, such as above 900 C and up to at least 1250 C, and
especially in applications wherein the mechanical load on the material can
become
high. Furthermore, the alloy according to the invention is suitable for use in
environments with high requirements for good oxidation resistance. Examples of
suitable applications are as construction materials for heat treatment
furnaces, in
rollers for roller hearth furnaces, as muffle tubes for annealing in
protective
atmosphere, as construction material for heating elements, combustion chamber
material in gas turbines, as gas-to-gas heat exchangers for example in the
glass
manufacturing industry or in gas turbines, as transportation belts woven from
wire
intended for heat treatment furnaces, in radiation tubes for heating in heat
treatment furnaces or as protective tubes for thermocouples.
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Brief description of the drawings
Figure 1a shows the result of a simulation of the effect of the Ni
content on
the phase stability at different temperatures.
Figure 1 b shows the influence of varying contents of Al and Fe on the
5 minimum stability of y'.
Figure 1c shows the influence of varying contents of AL and Cr on the
minimum stability of y'.
Figure 2 shows the result of a simulation of the effect of the Fe
content on
the stability of nickel aluminides.
Figure 3 shows the result of a simulation of the effect of the Al content
on
the stability of nickel aluminides.
Figure 4 shows the result of a simulation of the effect of Co on the
stability
of nickel aluminides.
Figure 5 shows result from tensile testing of examples of the alloy
according to the invention.
Figure 6 shows the yield strength of six different heats according to
the
invention at room temperature, 500 C and 600 C.
Figure 7 shows the tensile strength of six different heats according
to the
invention at room temperature, 500 C and 600 C.
Figure 8 shows the elongation to fracture of six different heats according
to
the invention at room temperature, 500 C and 600 C.
Figure 9 shows the result from oxidation testing in air at 1000 C of
eight
different heats according to the invention and two comparative
materials.
Figure 10 shows the result from oxidation testing in air at 1100 C of eight
different heats according to the invention and two comparative
materials.
Figure 11 shows a photograph of the microstructure of Heat A taken in
SEM.
Figure 12a shows the size distribution of carbonitrides precipitates in
Heat A.
Figure 12b shows the size distribution of precipitates in Heats A-D.
Figure 13 shows the result from creep testing of compositions which
are not
dispersion strengthened.
Figure 14 shows the result from oxidation testing in air at 1100 C of
four
compositions which are not dispersion strengthened.
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Detailed description
As mentioned above, nickel based alloys alloyed with aluminium are generally
considered difficult to hot-work. An important factor is that there is only a
limited
temperature window between melting of the alloy and precipitation of unwanted
intermetallic phases, such as nickel aluminides. The alloying elements Al and
Cr
are both beneficial for the oxidation resistance but makes a nickel based
alloy
difficult to work since they increase the stability of nickel aluminides and
therefore
reduces the temperature window for hot-working of the alloy. The hot
workability of
the alloy is a very important factor for enabling that products thereof can be
readily
and economically produced. It has been found that the alloy in accordance with
the present invention has an increased temperature window for hot-working as a
result of its composition which gives the alloy a good hot-workability.
The present invention is based on the discovery that relatively high
addition of Fe to a nickel based alloy with 4-6 % Al and high content of Cr
reduces
the stability of the intermetallic phase y'. Precipitations of the phase y'
improves
the creep strength at low temperatures but makes the production more difficult
since the alloy becomes hard and brittle at too high contents of y'. Moreover,
y'
reduces the activity of Al in the alloy which makes the formation of the
protective
aluminium oxide on the surface more difficult. For an alloy intended for use
at high
temperatures, such as above 900 C, it is consequently important to reduce the
content of y', which is achieved by the composition of the alloy in accordance
with
the present invention.
Moreover, precipitations of y' in previously known aluminium oxide forming
nickel based alloys are not stable above approximately 1000 C whereby its
influence on the creep strength ceases during use of such alloys above this
temperature. The alloy according to the present invention comprises a minimum
content of y' and is furthermore primarily intended for use at high
temperatures
where there consequently is a risk of dissolution of y'. In order to keep the
creep
strength the alloy is therefore dispersion strengthened. This is accomplished
above all by the selected contents of carbon and nitrogen and possibly oxygen
in
combination with the selected contents of Ta, Zr, Hf, Ti and Nb. It is
possible to
produce the alloy by conventional melting production process, but in that case
the
dispersion strengthening will be insufficient if even achieved. The alloy is
therefore
produced by way of powder metallurgy. Solid components can thereafter be
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manufactured from the produced powder by compaction in accordance with
previously known techniques such as hot isostatic pressing (HIP) or cold
isostatic
pressing (CIP). If needed the manufactured solid component can thereafter be
further worked, for example by rolling, extrusion or drawing in order to
achieve the
desired product form. It is also possible to produce complex geometries
directly
from the powder by means of sintering.
The composition of the present alloy and the fact that it is dispersion
strengthened has resulted in a nickel based alloy which has an excellent
oxidation
resistance even at temperatures as high as at least 1100 C, is relatively easy
to
hot-work and has good creep strength.
According to a preferred embodiment of the dispersion strengthened
nickel based alloy according to the invention, the dispersed particles have an
average diameter of less than lpm, preferably less than 500 nm. Best results
are
achieved when the dispersed particles have an average diameter of 50-200 nm.
According to yet a preferred embodiment of the dispersion strengthened
nickel based alloy according to the invention, more than 85 % of the dispersed
particles should be equal to or less than 300 nm in diameter.
The effect of the various elements on the properties of the alloy will now
be discussed below, wherein all given contents are in percent by weight.
Carbon
Carbon in free form will take interstitial locations in the crystal structure
and
thereby lock the mobility of dislocations at temperatures up to approximately
400-
500 C. Carbon also forms carbides with other elements in the alloy such as Ta,
Ti,
Hf, Zr and Nb. In a microstructure with finely dispersed carbides, these
carbides
provide obstacles for the dislocation movement and have effect even at higher
temperatures. Carbon is an essential element to improve the alloy's creep
strength
since the dislocation mobility is the mechanism that generates creep
elongation.
Too high contents of C will however lead to the alloy becoming difficult to
cold
work due to deteriorated ductility at lower temperatures, such as below 300 C.
The
alloy therefore comprises 0.05-0.2 % C.
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Silicon
Silicon can be present in the alloy in contents up to 1.5 %. Silicon in too
high
contents can in nickel based alloys lead to increased risk for precipitations
of
nickel silicides, which have an embritteling effect on this type of alloy.
Results from
creep testing of similar alloys have shown that the creep life time, i.e. the
time to
creep fracture, is reduced with Si contents close to 1.5 %. The reason for
this is
however not known. Because of this, the Si content should preferably be
maximally 1 %. According to a preferred embodiment, the alloy only comprises
impurity contents of Si, i.e. up to 0.3 %.
Manganese
Manganese is present in the alloy as an impurity. It is likely that up to 0.5
% can be
allowed without negatively influencing the properties of the alloy whereby the
alloy
comprises maximally 0.5 % Mn. According to a preferred embodiment, the alloy
only comprises impurity contents of Mn, i.e. up to 0.2 %.
Chromium
Chromium is an element which for a long period of time has been the leading
element when it comes to creating a dense and protective oxide scale. Less
than
15 % Cr in an austenitic structure tends to render an oxide which is not
entirely
covering the surface and which is not dense and consequently render an
insufficient oxidation resistance to the alloy. There is also a risk that the
material
closest to the oxide is depleted of Cr such that possible damages to the oxide
can
not heal since there is not sufficient Cr to form new oxide.
A nickel based alloy comprising 4 % Al should however not comprise more
than about 20 % Cr as higher contents increases the risk of formation of y'
and 13
phases. (This will for example be shown below with reference to Figure lc,
calculated for an alloy comprising approximately 19 % Fe.)
Therefore, in order to minimise the presence of the y' and 13 phases, the
alloy comprises max 20 % Cr. There may also be a risk of formation of other
unwanted phases, such as a-phase and chromium rich ferrite, at too high Cr
contents. Moreover, Cr may also at high contents stabilise nickel aluminides.
Therefore, the alloy comprises 15-20 % Cr, preferably 17-20 % Cr. Best
results are achieved when the alloy comprises 17-19 % Cr.
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Aluminium
Aluminium is an element that generates a much denser and more protective oxide
scale compared to Cr. Aluminium can however not replace Cr since the formation
of the aluminium oxide is slower than the chromium oxide at lower
temperatures.
The alloy comprises at least 4 % Al, preferably more than 4 % Al, which
ensures a
sufficient oxidation resistance at high temperatures and that the oxide covers
the
surface entirely. The relatively high content of Al provides excellent
oxidation
resistance even at temperatures of about 1100 C. At Al contents above 6 %
io there is a risk of formation of such an amount of intermetallic phases
in a nickel
based matrix that the ductility of the material is considerably deteriorated
(this will
also be discussed below with reference to Figure 3). The alloy should
therefore
comprise 4-6 % Al, preferably >4-5.5 %, more preferably >4 - 5.2 % Al.
Iron
It has been shown in accordance with the present invention that relatively
high
contents of Fe in an aluminium oxide forming nickel based alloy can have
positive
effects. Additions of Fe generate a metallic structure which is energetically
unfavourable for the formation of embritteling y', which in turn leads to the
risk of
the alloy becoming hard and brittle reducing considerably. Consequently, the
workability is improved. Therefore, the alloy comprises at least 15 % Fe. High
contents of iron may however lead to formation of unwanted phases. Therefore,
the alloy shall not comprise more than 25 % Fe.
Moreover, at Fe contents over approximately 21-22 % the risk of formation
of a 3-phase (NiAI), which in some cases can be embritteling, increases. (This
will
for example be shown below with reference to Figures lb and 2.)
Preferably, the alloy should therefore comprise 16-21.5% Fe. According
to a preferred embodiment, the alloy comprises 17-21 % Fe.
Nickel
The alloy according to the invention is a nickel based alloy. Nickel is an
element
which stabilises an austenitic structure in alloys and thereby counteracts
formation
of some brittle intermetallic phases, such as a-phase. The austenitic
structure of
the alloy is beneficial for example when it comes to welding. The austenitic
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structure also contributes to the good creep strength of the alloy at high
temperatures. This could be a result of that the diffusion rate is lower in an
austenitic structure than for example in a ferritic.
According to one embodiment, the alloy comprises 52-62 % Ni, preferably
5 52-60% Ni.
Cobalt
In some commercial alloys, a part of the Ni is substituted with Co in order to
increase the mechanical strength of the alloy which may also be done in the
alloy
10 according to the invention. A part of the Ni of the alloy can be
replaced with an
equal amount of Co. This increases the stability of the BCC-aluminide NiAl,
which
then grows at the expense of y', which can be advantageous in certain
temperature ranges. This Co addition must however be balanced against the
oxidation properties since the presence of NiAl will reduce the activity of Al
and
thereby deteriorate the ability to form aluminium oxide. The addition of Co
will also
affect the melting point of the alloy. For example, an addition of 10 % Co
will
render an alloy with precipitations of NiAl which are stable up to 950 C but
lower
the melting point with approximately 20 C. According to one embodiment of the
present invention, nickel is therefore partly substituted with Co. The Co
content
shall, however, not exceed 10 %.
Nitrogen
In the same way as C, free N takes interstitial locations in the crystal
structure and
thereby locks the dislocation mobility at temperatures up to approximately 400-
500 C. Nitrogen also forms nitrides and/or carbonitrides with other elements
in the
alloy such as Ta, Ti, Hf, Zr and Nb. In a microstructure where these particles
are
finely dispersed they confer obstacles for the dislocation mobility,
especially at
higher temperatures. Therefore, N is added in order to improve the creep
strength
of the alloy. However, when adding N to aluminium alloyed alloys the risk is
high
for formation of secondary aluminium nitrides and the present nickel based
alloy
therefore has a very limited N content. The alloy comprises 0.03-0.15 % N,
preferably 0.05-0.15 % N, more preferably 0.05-0.10 % N.
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Oxygen
Oxygen may be present in the present alloy either in the form of an impurity,
or as
an active addition up to 0.5 %. Oxygen may contribute to increasing the creep
strength of the alloy by forming small oxide dispersions together with Zr, Hf,
Ta
and Ti, which, when they are finely distributed in the alloy, improves its
creep
strength. These oxide dispersions have higher dissolution temperature than
corresponding carbides and nitrides, whereby oxygen is a preferred addition
for
use at high temperatures. Oxygen may also form dispersions with Al, the
elements
in group 3 of the periodic table, Sc, Y and La as well as the fourteen
lanthanides,
and in the same manner as with the above identified elements thereby
contribute
to higher creep strength of the alloy. According to a preferred embodiment,
the
alloy comprises 200-2000 ppm 0, preferably 400-1000 ppm 0.
Tantalum, Hafnium, Zirconium, Titanium and Niobium
The elements in the group consisting of Ta, Hf and Zr forms very small and
stable
particles with carbon and nitrogen. It is these particles which, if they are
finely
dispersed in the structure, help to lock dislocation movement and thereby
increase
the creep strength, i.e. provides the dispersion strengthening. It is also
possible to
accomplish this effect with addition of Ti. Additions of Ti can, however,
sometimes
lead to problems, especially during powder metallurgical production of the
alloy,
since it forms carbides and nitrides already in the melt before atomisation,
which in
turn may clog the orifice during the atomisation.
Niobium also forms stable dispersions with C and or N and can therefore
suitably be added to the alloy according to the invention.
The alloy comprises one or more elements selected from the group
consisting of Ta, Zr, Hf Ti and Nb in an amount of 0.25-2.2 %, preferably 0.3-
1.5
%, more preferably 0.6-1.5 %.
The alloy preferably comprises such an amount of the elements Ta, Zr, Hf,
Ti and Nb that essentially all C and N is bound to these elements. This
ensures
that for example the risk of formation of chromium carbides during high
temperature use of the alloy is significantly reduced.
According to a preferred embodiment, the alloy comprises 0.1-0.5% Hf.
According to another embodiment, the alloy comprises 0.05-0.35 % Zr. According
to yet another embodiment, the alloy comprises 0.05-0.5 % Ta. According to yet
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another embodiment, the alloy comprises 0.05-0.4 % Ti. According to yet
another
embodiment, the alloy comprises 0.1-0.8 % Nb.
Rare earth metals (REM)
Rare earth metals (REM) relates in this context to the elements of group three
of
the periodic table, Sc, Y, and La as well as the fourteen lanthanides. REM
affects
the oxidation properties by doping of the formed oxide. Excess alloying of
these
elements often gives an oxide which tends to spall of the surface and a too
low
addition of these elements tends to give an oxide with weaker adhesion to the
metal surface. The alloy may comprise one or more elements from the group
consisting of REM in a content of up to 0.5 % in total, preferably 0.05-0.25
%.
According to a preferred embodiment, yttrium is added to the alloy in an
amount of
0.05-0.25 %.
The nickel based alloy according to the invention may also comprise
normally occurring impurities as a result of the raw material used or the
selected
manufacturing process. Examples of impurities are Ca, S and P.
The dispersion strengthened nickel based alloy has a very good oxidation
resistance inter alia as a result of the Al and Cr contents. It also has very
good
mechanical properties, such as yield and tensile strength as well as
ductility. It has
very good workability, especially hot workability, which makes it easy to
manufacture products by for example hot extrusion or hot rolling.
The above identified nickel based alloy is foremost intended for use at
high temperatures. Examples of applications wherein the alloy is especially
suitable are construction materials for heat treatment furnaces, rollers for
roller
hearth furnaces, muffle tubes for annealing in protective atmosphere,
construction
material for heating elements, combustion chamber material in gas turbines,
gas-
to-gas heat exchangers for example in the glass manufacturing industry or in
gas
turbines, tubular reactors in high temperature processes, transportation belts
of
woven wires intended for heat treatment furnaces, radiation tubes for heating
of
heat treatment furnaces or protective tubes for thermocouples.
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Simulation
The stability of phases at different alloy compositions and temperature has
been
studied by thermodynamic simulations using the software Thermo-Calc. A
thermodynamic database for nickel based alloys called "NiFe-Super version 4"
-- was used for the simulations. It is commonly known that this type of
calculations in
most cases correspond well to the reality.
The influence of iron on the stability of the nickel aluminides 13 (NiAl) and
y'
(Ni3A1), and the stability of a (chromium rich ferrite) was studied. The
calculations
were made for the chromium content 18 wt-% and the aluminium content 4.5 wt-
-- %. The result for a simulation wherein temperature and nickel content have
been
varied is shown in Figure 1. Along the x-axis iron is replaced with nickel in
the
alloy.
These simulations have shown that there is an area for an alloy with 4.5
wt-% Al and 18 wt-% Cr where the stability of y' has a minimum. This minimum
is
-- at 58 wt-% Ni and with an iron content of approximately 19 wt-%, and is
marked in
the figure by the dotted circle. Lower contents of Fe increases the stability
of y'
whereas higher contents render formation of the nickel aluminide 13 (NiA1).
Compositions around this minimum give a wide temperature interval
between melting of the alloy and precipitation of nickel aluminides and
therefore
-- facilitate the hot-workability as explained above.
The influence of variations in Al and Cr content on the minimum identified
above has also been studied. By varying the Al content between 4 and 6 % and
at
the same time adjusting the Fe content such that the minimum in y'-stability
is
achieved, Figure lb can be calculated. Figure lb shows how the minimum is
-- moved when the contents of Fe and Al are varied. The minimum is moved along
the line in the figure at the same time as the temperature is changed. It is
clear
from the figure that increased Al content reduces the amount of Fe necessary
to
achieve the minimum. Moreover, the temperature rises for the minimum from 814
C at tic-mark 1 to 953 C at tic-mark 9.
Figure lc shows the same type of calculation as in Figure lb but wherein
the Cr and Al contents have been varied and the Fe content is kept at
approximately 19 %. It is clear from the figure that increased Al content
reduces
the amount of Cr necessary to achieve the minimum. Moreover, the temperature
rises from 815 C attic-mark 1 to 951 C attic-mark 10.
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In Figure 2, the influence of different iron contents on the stability of
nickel
aluminides, ferrite and austenite is shown. The composition was in this case
18 wt-
% Cr, 4.5 wt-% Al, balance Ni with three different iron contents 16 wt-%, 19
wt-%
and 22 wt-%, respectively. The lowest dissolution temperature for nickel
aluminides was obtained for the Fe content of 19 %. At the highest Fe content,
13 is
stable whereas the lowest Fe content increases the stability of y' which
results in a
higher dissolution temperature.
In Figure 3. the influence of different Al contents on the stability of nickel
aluminides and ferrite is shown. The composition was in this case 18 wt-% Cr,
19
wt-% Fe, balance Ni with four different Al contents 4 wt-%, 4.5 wt-%, 5 wt-%
and 6
wt-% respectively. Increasing Al contents increases the dissolution
temperature for
nickel aluminides. At the Al content of 6 % the intermetallic 13-phase is
stable up to
temperatures around 1100 C. Increasing Al contents increase the stability of
ferrite
at lower temperature ranges, below approximately 800 C.
Simulation of the effect of cobalt addition
In order to investigate which effect cobalt would have on the alloy
simulations were
made using the software Thermo-Calc. A thermodynamic database for nickel
based alloys called "NiFe-Super version 4" was used for the simulations. The
calculations were made with the starting composition 18 % Cr, 19 % Fe, 4.5 %
Al,
balance Ni. Nickel was substituted with 5, 10, and 15% Co in the starting
composition and the balance fraction of precipitations was calculated as a
function
of temperature. The influence of Co on the stability of the nickel aluminides
13
(NiAl) and y' (Ni3A1), a (chromium rich ferrite) as well as a-phase was
studied. The
result is shown in Figure 4.
The calculations show that additions of Co increase the dissolution
temperature for nickel aluminides. Additions of Co also increase the stability
of the
nickel aluminide 13 relative to y'. At the two highest contents of Co there is
also a
risk for precipitations of a-phase at temperatures below approximately 650 C,
Up
to 10 wt-% Co can be used in the alloy for uses at temperatures above 950 C.
Tensile testing
A number of compositions of the alloy according to the invention were produced
by
means of powder metallurgy and compacted by hot isostatic pressing followed by
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hot extrusion and subsequent water quenching. The compositions of the
different
heats are given in Table 1.
Table 1
Heat 1 2 A B C D E F
C 0.05 0.14 0.072 0.083 0.11 0.17 0.12 0.12
Si 0.05 0.10 0.09 0.05 0.04 0.04 0.08 0.09
Mn 0.06 0.10 0.10 0.10 0.06 0.06 0.10 0.09
Fe 19.0 18.8 18.8 18.8 19.1 19.1 19.8 21.0
Cr 18.2 17.7 17.7 17.8 17.8 18.0 18.0 18.1
Ni Bal. Bal. Bal. Bal. Bal. Bal. Bal. Bal.
Al 4.50 4.54 4.80 4.59 4.64 4.66 4.59 4.65
Co 0.02 0.02 0.03 0.04 0.18 0.35 2.56 6.93
Nb <0.01 0.01 <0.01 <0.01 0.24 0.57 <0.01 <0.01
Ti <0.01 <0.01 <0.01 <0.01 0.14 0.33 <0.01 0.02
Zr <0.01 0.33 0.39 0.41 0.18 0.40 0.40 0.41
Ta <0.001 0.353 0.42 0.46 0.22 0.48 0.46 0.44
Hf 0.395 0.455 0.56 0.48 0.36 0.41 0.45 0.50
Y 0.273 0.285 0.26 0.21 0.26 0.21 0.19 0.13
N 0.071 0.072 0.068 0.071 0.068 0.067 0.077 0.077
0 0.025 0.034 0.0394 0.0345 0.0349 0.0238 0.0368 0.0290
Ta+Zr+ 0.395 1.148 1.37 1.35 1.14 2.19 1.31 1.37
Hf+Ti+
Nb
5
Tensile testing of the compositions was performed according to standard
SS-EN 10002-1 at room temperature. Three samples of each composition were
tested and the results from the tensile testing in the form of the average of
the
three samples are shown in Table 2. Moreover, Heat 1 was also tested directly
10 after HIP (i.e. prior to extrusion).
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Table 2
Sample Yield strength Yield strength Tensile Elongation to
Rpo 2 RP10 strength fracture A
[N/mm2] [N/mm2] Rm [A]
[N/mm2]
1 (after HIP) 400 not analysed 826 37
1 (extruded) 389 not analysed 814 34
2 430 not analysed 877 32
A 492 548 963 35
501 557 971 35
530 587 989 34
538 602 1002 32
488 553 964 31
441 510 932 35
The results show that the alloy according to the invention has a good
elongation to fracture at room temperature which reduces the risk for crack
formation during cold working. Furthermore, the alloy has a yield strength
which is
higher than many austenitic steels and nickel based alloys, which generally
have a
yield strength of approximately 200-300 MPa. The results can for example be
compared with an austenitic chromium-nickel steel with a nominal composition
of
0.07 wt-`)/0 C, 1.6 wt-`)/0 Si, 1.5 wt-`)/0 Mn, 25 wt-`)/0 Cr, 35 wt-`)/0 Ni,
0.16 wt-`)/0 N, 0.05
wt-% Ce and balance Fe (corresponding to UNS S35315), which has a yield
strength Rpo 2 of about 260 MPa, a tensile strength Rm of about 600 MPa and an
elongation to fracture of about 35 %. The results could also be compared to a
dispersion strengthened aluminium oxide forming ferritic steel known under the
trade name KANTHAL APMT which has a nominal composition comprising 21 wt-
(:)/0 Cr, 5 wt-% Al, 3 wt-% Mo, max 0.7 (:)/0 Si, max 0.4 wt-% Mn, max 0.08 wt-
% C,
and which has a yield strength Rpo 2 of about 550 MPa, a tensile strength Rm
of
about 750 MPa and an elongation to fracture of about 25 %.
Moreover, tensile testing at 500 C and 600 C respectively of Heats A-F
given in Table 1 was made in accordance with standard SS-EN 10002-5. Three
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samples of each composition were tested and the results from the tensile
testing in
the form of the average of the three samples are shown in Table 3.
The results from the tensile testing at 500 C and 600 C indicates that the
alloy according to the invention has good high temperature mechanical
properties
and has good elongation to fracture at these temperatures. This, together with
successful results from hot extrusion and hot rolling, indicates that the
alloy has
good hot-workability.
The results from the tensile testing of Heats 1 and 2 are shown in Figure 5
and the results from the tensile testing of Heats A-F are shown in Figures 6
to 8.
Table 3
Sample Yield Tensile Elongation Yield Tensile
Elongation
strength strength to fracture strength strength to
fracture
Rpo 2 at Rm at A at Rpo 2 at Rm at A at
500 C 500 C 500 C 600 C 600 C 600 C
[N/mm2] [N/mm2] [A] [N/mm2] [N/mm2] [A]
A 423 845 28 486 817 19
B 428 864 28 499 827 19
C 467 889 29 503 820 19
D 461 892 27 577 848 23
E 417 857 24 472 824 19
F 373 812 25 430 834 28
Impact testing
Impact testing was performed on material produced from the metal powders from
the heats given in Table 1. Samples were produced by hot isostatic pressing
(HIP)
and subsequent hot extrusion with water quenching. Testing according to SS-EN
10045-1 was performed at room temperature and was performed on three
samples each of the compositions. The results are shown in Table 4.
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Table 4
Heat Sample Impact strength [J] Average value [J]
1 1 133 119
2 106
3 117
2 1 50 45
2 42
3 44
A 1 84 92
2 91
3 102
B 1 82 81
2 92
3 68
C 1 48 47
2 46
3 48
D 1 64 64
2 64
3 64
E 1 50 49
2 50
3 46
F 1 68 65
2 64
3 62
The impact strength for all heats is well above the 27 Joule which is
generally used as a limit value between ductile and brittle material.
Oxidation test at 1000 C
Samples in the form of coupons were produced from the heats given in Table 1.
The coupons were grid with 220 pm paper. Furthermore, one sample of a nickel
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19
based alloy known under the trade name SANDVIK SANICRO 80 (corresponding
to UNS N06003) and one sample of the dispersion strengthened aluminium oxide
forming ferritic steel known under the trade name KANTHAL APMT (which has a
nominal composition comprising 21 wt-% Cr, 5 wt-% Al, 3 wt-% Mo, max 0.7 % Si,
max 0.4 wt-% Mn, max 0.08 wt-% C), were produced for comparison.
Oxidation test was performed at 1000 C in air. The samples were
removed from the furnace and cooled to room temperature after 24, 48, 95, 186,
500 and 1005 hours respectively and weighed. After weighing, the samples were
inserted into the furnace for continued heating and oxidation. The results
from the
oxidation test are shown in Figure 9.
The results show that the alloy according to the invention has a very good
oxidation resistance at 1000 C. All heats except D have considerably better
oxidation resistance than SANDVIK SAN ICRO 80. Furthermore, the alloys
according to the invention have an oxidation resistance at this temperature
which
is comparable to that of KANTHAL APMT, which is an alloy that is considered to
have an excellent oxidation resistance.
The alloys according to the present invention quickly form a protective
oxide which after formation has a very slow growth rate. No negative effects
of the
high iron content, which have previously been reported in patents US 4,882,125
and US 4,460,542 could be observed. It can be noted that most chromium oxide
forming austenitic alloys commonly used at high temperatures have an oxide
growth rate which is more than 4-8 times as high at this temperature.
Oxidation test at 1100 C
Samples were produced from the same compositions and in the same manner as
in the case of the oxidation test at 1000 C. An oxidation test was performed
at
1100 C in air. Samples were removed after 24, 48, 95, 186, 500 and 1005 hours
respectively and weighed. The results from the oxidation test are shown in
Figure
10.
The results show that the alloy according to the invention has very good
oxidation resistance at 1100 C. The reference alloys used in this work,
SANDVIK
SAN ICRO 80 and KANTHAL APMT, are known to have excellent oxidation
resistance for chromia formers and for ferritic alumina formers, respectively.
The
oxidation test of the alloys according to the present invention shows, in
general,
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better oxidation resistance than that of SAN DVIK SAN ICRO 80 and some even
better than that of KANTHAL APMT. All tested alloys show a substantially
better
oxidation resistance than that of the alloy presented in WO 00/34541.
Tentative
oxidation studies at 1200 C indicate that the alloy according to the present
5 invention shows an even higher degree of oxidation resistance compared to
the
chromia forming alloys SANDVIK SANICRO 80 and the previously mentioned
UNS S35315. This shows that the aluminium addition in the developed alloy
increases the oxidation resistance, especially at temperatures above 1100 C.
10 Microstructure
An example of the microstructure in a material, with the composition according
to
Heat A, produced from metal powder which was compacted by HIP, hot extruded
and water quenched is shown in Figure 11. The photograph was taken in a
scanning electron microscope (SEM) with a 30 000x magnitude. The light
15 precipitates seen in the microstructure are carbonitrides containing
mainly Hf, Ta
and Zr.
An image analysis of close to 10000 carbonitride precipitates of the
material in Figure 11 was performed using SEM. The average diameter of the
precipitates was about 130 nm. The frequency of carbon itride precipitates in
20 different size ranges from the image analysis is shown in Figure 12a.
Furthermore, the size of the dispersion strengthening precipitates in Heats
B to D was investigated. Figure 12b shows the relative frequency of the
particle
diameter of Heats A to D. It is clear that the dispersions in all heats
generally have
a diameter of less than 300 nm.
Creep testing of Heats 1 and 2
The creep strength for Heat 1 and Heat 2 given in Table 1 was performed. Test
samples were produced from metal powder which was compacted by HIP. During
the creep testing threaded samples with a length of 35 mm and a 5 mm diameter
at the waist were used. The testing was performed at the temperature 1200 C
and
4 MPa load. The test was performed for double samples. Heat 1, which comprises
only a small content of dispersion strengthened particles due to the low
content of
C (0.05 %) and only 0.395 % Hf (no additions of Nb, Ti, Zr and Ta), showed a
time
to fracture of 358 and 387 hours respectively, for the samples. However, Heat
2
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21
which has a relatively high content of dispersion strengthened particles due
to the
relatively high content of C (0.14 %) and 1.148% in total of Zr, Ta and Hf,
showed
a time to fracture of 3064 and 4576 hours respectively. The beneficial effect
of the
dispersion strengthening is thus clear from these results.
Creep testing of Heats A-F
Test samples for creep testing were produced from metal powder which was
compacted by HIP and thereafter hot extruded from 77 mm diameter to 25 mm
diameter followed by water quenching. During the creep testing threaded
samples
with a length of 35 mm and a 5 mm diameter at the waist were used. The testing
was performed at the temperature 1200 C with 5 MPa load and at the temperature
1000 C with 15MPa load. The time to rupture for the different materials is
shown in
the Table 5.
Table 5.
Time to rupture tr Time to rupture tr
(h) (h)
Heat 1200 C/5M Pa 1000 C/15M Pa
A 571 337
B 689 1629
C 780 496
D 4041W 4007W
E 223 286*
F 327 263
,.
Ongoing test
The results show how material according to the invention has a creep
strength superior to commercially available wrought nickel base alloys. It
also
shows how material according to the invention has sufficient creep strength
and
oxidation resistance for practical usage at temperatures exceeding 1200 C in
contrast to a vast majority of commercially available nickel base alloys.
The high creep strength of Heat D is believed to be a result of the high
content of carbon as well as the high contents of Ti, Nb, Ta, HF and Zr.
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Creep testing of heats which are not dispersion strengthened
A number of experimental heats of approximately 1 kg size were produced by
induction melting and casting under a protective argon atmosphere for sake of
comparison. The heats were not dispersion strengthened since they were not
powder metallurgically produced. The compositions are given in Table 6.
The produced materials were then turned to rods with a diameter of 15
mm and thereafter hot rolled at 1200 C. Test samples for creep testing were
produced from work pieces which had been hot rolled to 10 mm square cross
section. During the creep testing threaded samples with a length of 35 mm and
a
diameter of 5 mm at the waist were used.
Table 6
Heat 4249 4250 4251 4252 4253 4254 4256 4257 4258
C 0.13 0.13 0.05 0.09 0.12 0.06 0.13 0.09 0.05
Si n.a. n.a. n.a. n.a. n.a. n.a. n.a. n.a.
n.a.
Cr 18.8 18.7 18.3 17.3 18.0 18.0 17.9 17.9 17.8
Ni 55.5 52.0 52.2 58.3 58.2 58.1 57.8 58.3 58.9
Al 5.2 4.2 4.1 4.3 4.3 4.2 4.2 4.3 4.1
Nb n.a. n.a. n.a. n.a. n.a. n.a. n.a. n.a.
n.a.
Ti n.a. n.a. n.a. n.a. n.a. n.a. n.a. n.a.
n.a.
Ta 0.35 0.14 0.13 0.32 0.14 0.14 0.14 0.35 0.32
Zr 0.37 0.49 0.47 0.51 0.50 0.47 0.16 0.13 0.03
N 0.048 0.057 0.06 0.012 0.061 0.058 0.054 0.014 0.054
Y 0.01 0.01 0.006 0.005 0.009 0.011 0.006 0.019 0.012
Fe 19.3 26.0 25.2 18.1 18.6 18.6 18.5 18.5 18.7
Hf 0.24 0.17 0.16 0.35 0.15 0.15 0.24 0.20 0.16
Ta+Zr+ 0.96 0.80 0.76 1.18 0.79 0.76 0.54 0.68 0.51
Hf+Nb+Ti
n.a.= not added
The creep testing was performed at the temperature of 1200 C and 4 MPa
load. The result is shown in Figure 13.
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A comparison of the times to fracture in Figure 13 with those of the tests of
Heat 2 above shows the beneficial effect on the creep strength when the
material
has been produced by powder metallurgy. Heat 2 was tested under the same load
and temperature as the comparative melts given in Table 6, and showed times to
fracture above 3000 hours, whereas the comparative melts all fractured well
under
500 hours.
Heat 4249, which has a high content of C (0.13 %) and a relatively high
content of Ta+Zr+Hf (0.96 %), still has a creep strength below 500 hours to
fracture whereas Heat 2 comprising approximately the same amount of C (0.14 %)
and a slightly higher content of the dispersion strengthening elements
(1.148%)
showed more than 6 times the time to fracture.
Oxidation test at 1100 C of heats which are not dispersion strengthened
Samples in the form of coupons were produced from heats 4249, 4251, 4257, and
4258 and grid with 220 pm paper. The samples were oxidation tested at 1100 C
in
air. The samples were removed after 24, 48, 96, 186, 500, and 1000 hours
respectively and weighed. Results from the oxidation test are shown in Figure
14.
The results show that the alloy has a very good oxidation resistance at
1100 C. Since the oxidation properties of the material should be independent
of
the dispersion strengthening, the results indicate that powder metallurgically
produced dispersion strengthened alloys with the same compositions, that is,
the
alloy according to the invention, should exhibit equally good oxidation
resistance at
this temperature.
