Note: Descriptions are shown in the official language in which they were submitted.
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METHOD FOR PRODUCTION OF AN IRON-CHROMIUM ALLOY
The invention relates to a ferritic iron-chromium alloy produced
by melting metallurgy.
DE 100 25 108 Al discloses a high-temperature material
comprising an iron alloy forming chromium oxide with up to 2% by
weight of at least one oxygen-affine element from the group Y,
Ce, Zr, Hf and Al, up to 2% by weight of an element M from the
group Mn, Ni and Co, which forms a spinel phase of MCr204 type at
high temperatures, up to 2% by weight of a further element from
the group Ti, Hf, Sr, Ca and Zr, which increases the electrical
conductivity of Cr-based oxides. The chromium content should lie
within a concentration range between 12 and 28%. Areas of use
for this high-temperature material are bipolar plates in a high-
temperature fuel cell.
.EP 1 298 228 Al relates to a steel for a high-temperature fuel
cell that has the following composition: not more than 0.2% C,
not more than 1% Si, not more than 1% Mn, not more than 2% Ni,
15 - 30% Cr, not more than 1% Al, not more than 0.5% Y, not more
than 0.2% REM and not more than 1% Zr, the remainder being iron
and production-related impurities.
Low hot strength and inadequate creep strength at temperatures
of 700 C and above are common to these two alloys. Precisely in
the range above 700 C up to approximately 900 C, however, these
alloys have excellent oxidation and corrosion resistance.
A creep-resistant ferritic steel, comprising precipitates of an
intermetallic phase of Fe2(M, Si) or Fe7(M, Si)6 type with at
least one metallic alloying element M, which may be formed by
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the elements niobium, molybdenum, tungsten or tantalum, has
become known from DE 10 2006 007 598 Al. The steel is preferably
intended to be used for a bipolar plate in a fuel-cell stack.
EP 1 536 031 Al discloses a metallic material for fuel cells,
containing C .1C_ 0.2%, 0.02 to 1% Si, 2% Mn, 10 to 40% Cr, 0.03
to 5% No, 0.1 to 3% Nb, at least one of the elements from the
group Sc, Y, La, Ce, Pr, Nd, Pm, Sn, Zr and Hf
1, the remainder
being iron and unavoidable impurities, wherein the composition
is supposed to satisfy the following equation: 0.1 Mo / Nb
30.
EP 1 882 756 Al describes a ferritic chromium steel, especially
usable in fuel cells. The chromium steel has the following
composition: C max. 0.1%, Si 0.1 - 1%, Mn max. 0.6%, Cr 15 -
25%, Ni max. 2%, No 0.5 - 2%, Nb 0.2 - 1.5%, Ti max. 0.5%, Zr
max. 0.5%, REM max. 0.3%, Al max. 0.1%, N max. 0.07%, the
remainder being Fe and melting-related impurities, wherein the
content of Zr + Ti is at least 0.2%.
In comparison with DE 100 25 108 Al and EP 1 298 228 A2, all of
these alloys have better hot strength and elevated creep
strength at temperatures of 700 C and above, specifically due to
formation of precipitates, which hinder the dislocation
movements and thus plastic deformation of the material. In the
case of DE 10 2006 007 598 Al, for example, these precipitates
consist of a Laves phase, an intermetallic compound with the
composition Fe2(M, Si) or Fe7(M, Si)6, wherein M may be niobium,
molybdenum, tungsten or tantalum. Therein a proportion by volume
of 1 to 8%, preferably 2.5 to 5%, should be reached. However,
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there may also be other precipitates such as Fe-containing
particles and/or Cr-containing particles and/or Si-containing
,particles, such as described, for example, in EP 1 536 031 Al,
or there may be carbides containing Nb, W, No. It is common to
all of these particles that they make deformation of the
material difficult.
From the state of the art described above, it is known that
small additions of Y, Zr, Ti, Hf, Ce, La and similar reactive
elements can influence the oxidation resistance of Fe-Cr alloys
very positively.
The alloys cited in DE 10 2006 007 598 Al, EP 1 536 031 Al and
EP 1 882 756 Al are optimized for the application as
interconnector plates for the high-temperature fuel cells: By
use of a ferritic alloy containing 10 to 40% chromium, they have
an expansion coefficient adapted as well as possible to the
ceramic components anode and electrolyte.
Further requirements on the interconnector steel of a high-
temperature fuel cell are, besides the creep strength already
mentioned above, very good corrosion resistance, good
conductivity of the oxide layer and little chromium
volatilization.
The requirements on the reformers and the heat exchangers for
the high-temperature fuel cell are the best possible creep
strength, very good corrosion resistance and little chromium
volatilization. The oxide for these components does not have to
be conductive.
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The requirements for components for the exhaust-gas line of a
combustion engine, for example, or for steam boilers,
superheaters, turbines and other parts of a power plant, are
best possible creep strength and very good corrosion resistance.
In these cases chromium volatilization does not cause any
poisoning phenomena as in the fuel cell, and the protecting
oxide does not have to be conductive for such components.
In DE 10 2006 007 598 Al, for example, the excellent corrosion
resistance is achieved by formation of a chromium oxide top
layer. By the fact that a spinel containing Mn, Ni, Co or Cu is
additionally formed on the chromium oxide top layer, fewer
volatile chromium oxides or chromium oxyhydroxides that poison
the cathode are formed. By the fact that Si is stably bound in
the Fe2(M, Si) or Fe7(M, Si)6 Laves phase, a nonconductive
subsurface layer of silicon oxide is also not formed under the
chromium oxide top layer. The corrosion resistance is further
improved by the fact that the Al content is kept low and so the
increase of the corrosion due to the internal oxidation of the
aluminum is avoided. A small Ti addition additionally favors
strengthening of the surface and thus prevents swelling of the
oxide layer and the inclusion of metallic zones in the oxide
layer, which increases the oxidation. In addition, the addition
of oxygen-affine elements such as La, Ce, Y, Zr or the like
further increases the corrosion resistance.
From the market, increased requirements are being imposed on
products, necessitating elevated hot strength and creep strength
together with an elongation of at least 18% at application
temperature for avoidance of brittle failure together with at
least equally good oxidation or corrosion resistance and a
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higher service temperature of the alloy, specifically while
retaining acceptable deformability, measured as plastic
deformation in the tension test with an elongation of > 13% at
room temperature.
Furthermore, the following investigation methods are used.
In a creep test, a specimen is subjected to a constant static
tensile force at a constant temperature. For the purpose of
comparability, this tensile force is expressed as an initial
tensile stress relative to the initial cross section of the
specimen. In the creep test, the time tB until break - the time
to break - of the specimen is measured in the simplest case. The
test can then be performed without measurement of the elongation
of the specimen in the course of the test. The elongation at
break is then measured after the end of the test.
The specimen is mounted at room temperature in the creep-testing
machine and heated to the desired temperature without loading by
a tensile force. After reaching the test temperature, the
specimen is maintained for one hour without loading for
temperature equilibration. Thereafter the specimen is loaded
with the tensile force and the test time begins.
The time to break can be taken as a measure of the creep
strength. The longer the time to break is at a specified
temperature and initial tensile stress, the greater the creep
strength of the material is. The time to break and the creep
strength decrease with increasing temperature and increasing
initial tensile stress (see, for example, "Burgel", page 100).
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The deformability is determined in a tension test according to
DIN 50145 at room temperature. In the process, the offset yield
strength Rr0.2, the tensile strength Rm and the elongation at
break are determined. The elongation A is determined on the
broken specimen from the elongation of the original gauge length
Do:
A - (Lu-L0)/L0 100% = AL/L0 100%
Where Lu = gauge length after break.
Depending on gauge length, the elongation at break is denoted by
subscripts:
A5r gauge length Lo = 5=d0 or Lo = 5.65 = -\/S
A10, gauge length Lo = 10.d0 or Lo = 11.3 = -\/S0
or, for example, AL=100r for the freely chosen gauge length L =
100 mm.
(do initial diameter, Se initial cross section of the flat
specimen)
The magnitude of the elongation A in the tension test at room
temperature can be taken as a measure of the deformability.
The Laves phase(s) or the Fe-containing particles and/or Cr-
containing particles and/or Si-containing particles and/or
carbides can be made visible on a metallographic ground section
by etching with V2A pickling fluid or electrolytic etching with
oxalic acid. During etching with V2A pickling fluid, the grains
or grain boundaries also are additionally etched visibly. Only
particles with a size of approximately 0.5 m and larger are
visible by viewing in an optical microscope. Smaller particles
may not be recognized, but are definitely present. Therefore
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metallography is used only in support of the explanation, while
the efficacy of a measure is assessed more practically by the
time to rupture or creep strength.
In the Manual of High-Temperature Materials Technology, Ralf
Burgel, 3rd Revised Edition, Viehweg Verlag, December 2006,
hereinafter referred to as "Btirgel", the "Possible measures for
increasing the creep strength of metallic materials" are
presented on pages 196 to 199 and in Table 3.7.
The measures
- "High melting point, face-centered cubic material",
- "High modulus of elasticity",
- "Material with low stacking fault energy",
cannot be used for improvement of the cited parameters, since
they necessitate a change of material type, which is not
possible here and also is not the task.
The measures
- "Solid solution hardening"
- "Particle hardening"
- "High particle volume fraction"
- "Particles with small diffusion coefficients of the
alloying element in question"
have already been employed in DE 10 2006 007 598 Al and/or in EP
1 536 031 Al and/or in EP 1 882 756 Al.
The measures
- "Particles with low solubility in the matrix"
- "Coherent particles with low interfacial enthalpy relative
to the matrix",
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are not applicable for the precipitates under consideration.
Likewise, the measures
- "Carbides or borides as grain-boundary precipitates; avoid
oxides and sulfides",
- "Add positively active grain-boundary elements in precisely
controlled dosage, for example, B, C, Zr, Ce",
- "Higher purity of the alloy"
- "Add getter elements (for example, for S)",
- "High corrosion resistance"
have already been described in DE 10 2006 007 598 Al and/or in
EP 1 536 031 Al and/or in EP 1 882 756 Al.
The measures
- "Small dendrite arm spacings in cast microstructures",
- "Small grain structure in the main loading direction",
- "Single crystal"
- "Low density of components loaded by their own weight and
rotating"
cannot be applied to this alloy type, or to the production
route, or the use.
For the task of improving the creep strength of the
precipitation-hardened iron-chromium alloy, the measures
1) "Coarse grain microstructure",
2) "Jaggedness of the grain boundaries due to precipitates",
3) "Optimized heat treatment (adjust optimum particle
diameter, eliminate segregations in cast microstructure,
purposefully adjust possible grain boundary roughness)",
4) "Avoid cold working",
are to be considered.
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The invention relates to a method for production
of a component made from a precipitation-hardened iron-chromium
alloy, by means of which the high hot strength or creep strength
of a precipitation-hardened ferritic alloy can be further
increased compared with the state of the art while retaining
acceptable deformability at room temperature.
It is also intended to provide a thermomechanically treated
component/semifinished product consisting of an iron-chromium
alloy, which can be used for achievement of high hot strength.or
creep, strength while retaining acceptable deformability at room
temperature.
Finally, it is intended that the component/semifinished product
produced in this way can be used for specific technical
applications in the temperature 'range above 550 C.
This is accomplished on the one hand by a method for
production of a component from an iron-chromium alloy
precipitating Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides, in that a semifinished product produced from the alloy
is subjected to a thermomechanical treatment, wherein in a first
step the alloy is solution annealed at temperatures the
=
solution-annealing temperature, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water, in a second step mechanical working of the
semifinished product in the range from 0.05 to 99% is performed
and in a subsequent step Fe2(M, Si) or Fe7(M, Si)6 Laves phases
and/or Fe-containing particles and/or Cr-containing particles
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and/or Si-containing particles and/or carbides are precipitated
purposefully and in finely dispersed form by the fact that the
component made from the worked semifinished product is brought
to an application temperature between 550 C and 1000 C by
heating at 0.1 C/min to 1000 C/min.
This is accomplished on the other hand by a method for
production of a component from an iron-chromium alloy
precipitating Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides, in that a semifinished product produced from the alloy
is subjected to a thermomechanical treatment, wherein in a first
step the alloy is solution annealed at temperatures the
solution-annealing temperature, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water, in a second step mechanical working of the
semifinished product in the range from 0.05 to 99% is performed
and in a subsequent step Fe2(M, Si) or Fe7(M, Si)6 Laves phases
and/or Fe-containing particles and/or Cr-containing particles
and/or Si-containing particles and/or carbides are precipitated
purposefully and in finely dispersed form by the fact that the
worked semifinished product is subjected for a time between tmin
and tmax to a heat treatment in the temperature range between
550 C and 1060 C under protective gas or air, followed by
quenching in stationary protective gas or air, moving (blown)
protective gas or air or in water or for heat treatments up to
800 C is quenched in the oven, wherein tmth and tinax are
calculated according to the following formulas:
tmin. = Ta = 10(6740/T9.216)
and tam( = Ta = 10 (17960/Ta-15.2) where Ta T
+273.15,
and wherein the desired component is made before or after this
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heat treatment.
The times tmin and trnaõ are expressed in minutes and the heat
treatment temperature T in C.
For the first step, the following temperature ranges and times
are practical for solution annealing:
> 1050 C for longer than 6 minutes
> 1060 C for longer than 1 minute
The resulting changes of the material characteristics are
explained in more detail in the course of the further
description.
Furthermore, there is also provided a metallic
component or semifinished product consisting of the following
chemical composition (in % by weight)
Cr 12 - 30%
Mn 0.001 - 2.5%
Nb 0.1 - 2%
= 0.1 - 5%
Si. 0.05 1- 1%
= 0.002 - 0.1%
= 0.002 - 0.1%
max. 0.01% S=
Fe remainder
as well as the usual. melting-related impurities,
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which at the end of a thermomechanical treatment has a deformed
microstructure, to the effect that Laves phase(s) is or are
embedded in finely dispersed form in the microstructural
dislocations of the microstructure, wherein, in a creep test
with, for example, 35 MPa at /50 C and an elongation of at least
18%, a time to break that exceeds the time to break of a coarse-
grained, completely recrystallized microstructure by a factor of
at least 1.5 is established in the microstructure.
A comparable result is achieved for creep tests with different
stresses and temperatures, wherein the temperatures for the
creep test preferably lie in the range between 500 and 1000 C.
Measures 1 to 4 described above will now be considered.
Surprisingly it has been found in this connection that, in
contrast to measure 4 "cold deformation", preworking followed by
an adapted annealing treatment can bring about prolongations of
the times to break of the specimen in the creep test that go
more than 1.5 times, preferably more than 3 times beyond the
times to break for a coarse-grained microstructure (measure 1).
Furthermore, it is proposed that, for the third step - the
precipitation of the Laves phase(s) - the worked semifinished
product or if applicable the component made therefrom, by a
combination of heating at 0.1 C/min to 1000 C/min is subjected
to a heat-treatment temperature between 550 C and 1060 C with
subsequent heat treatment for a time between tmin and tmõ at this
temperature under protective gas or air, followed by quenching
in stationary protective gas or air, moving (blown) protective
gas or air or in water or for heat treatments up to 800 C is
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quenched in the oven, after which the desired component is made
if applicable, wherein trtan and tmax are calculated according to
the following formulas: tõõ = Ta = 0(6740/Ta-9.216)and tõ = Ta =
(17960/Ta-15.72) where T, = T +273.15.
In addition, the possibility exists that, for the third step -
the precipitation of the Laves phase(s) - the worked
semifinished product or the component made therefrom is
subjected for a time between train and tõ,ax to a heat treatment in
the temperature range between 550 C and 1060 C under protective
gas or air, followed by quenching in stationary protective gas
or air, moving (blown) protective gas or air or in water or for
heat treatments up to 800 C is quenched in the oven, after which
the desired component is made if applicable, wherein trõn and tmax
are calculated according to the following formulas: tmin - 'a =
lo (67"/Ta-9.216)0 (17960/Ta-15.72)
and L./flax Ta =
where Tõ = T +273.15 and
then the component that has been made is brought by heating at
0.1 C/min to 1000 C/min to an application temperature between
550 C and 1000 C.
According to a further concept of the invention, a semifinished
product from an alloy of the following composition (in % by
weight) is treated thermomechanically:
Cr 12 to 30%
Mn 0.001 to 2.5%
Nb 0.1 to 2%
W 0.1 to 5%
Si 0.05 to 1%
= 0.002 to 0.03%
= 0.002 to 0.03%
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max. 0.01%
Fe remainder
as well as the usual melting-related impurities.
With the inventive method it is possible to produce semifinished
= products in the form of sheets, strips, bars, forgings, pipes or
wire and to make components in the most diverse forms needed for
the respective application.
It is of special advantage that only little or even no Fe2(M, Si)
or Fe7(M, Si)6 Laves phases and/or Fe-containing particles and/or
Cr-containing particles and/or Si-containing particles and/or
carbides are still present in the semifinished product after
solution annealing at temperatures the solution-annealing
temperature, preferably 1050 C for longer than 6 minutes or >
1060 C for longer than 1 minute, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water, in the initial state before deformation.
The working of the semifinished product can take place by hot
working. Alternatively, however, the forming can also be brought
about by cold working.
In the first case, the semifinished product is hot-worked with a
starting temperature > 1070 C, wherein the last 0.05 to 95% of
mechanical deformation is applied between 1000 and 500 C,
advantageously the last 0.5 to 90% between 1000 C and 500 C.
In the second case, the degree of cold working of the
semifinished product is 0.05 to 99%, advantageously 0.05 to 95%
or 0.05 to 90%.
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As a further concept according to the invention, it is proposed
that the mechanical working of the semifinished product be 20 to
99% and then the worked semifinished product be subjected for a
time between traifl and tmax to a heat treatment in the temperature
range between 950 C and 1060 C under protective gas or air,
followed by quenching in stationary protective gas or air,
moving (blown) protective gas or air or in water, after which
the desired component be made, wherein tmin and tmax are
calculated according to the following formulas:
train T a = 10 (6740/Ta-9.216) and
tmax = Ta = 1 0 (17960/Ta-15.72) where Ta = T +273.15
with
tmin and tmax in minutes
and the heat-treatment temperature T in degrees Celsius.
If the already indicated alloy is used as an interconnector for
a solid oxide fuel cell, then a content of 0.001 - 0.5% aluminum
is advantageous.
For other areas of use, such as, for example, in the reformer or
heat exchanger for the fuel cell, for which no conductive oxide
layer is necessary, a content of 2 to 6% aluminum is
advantageous, since then a closed aluminum oxide layer can form,
which once again has a much slower growth rate compared with a
chromium oxide layer and additionally has much less chromium
oxide volatilization than a chromium-manganese spinel.
For the areas of use that neither need a conducive oxide nor
have special requirements on chromium volatilization, both
variants may be considered. In this connection, it is to be kept
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in mind in particular that the processability and weldability of
the alloy deteriorate with increasing aluminum content and so
higher costs are incurred. Therefore, when an oxide layer.
consists of a chromium oxide and a chromium-manganese spinel,
adequate oxidation resistance can be assured by use of 0.001 -
0.5% aluminum. If greater oxidation resistance is necessary, as
is assured, for example, by the formation of an aluminum oxide
layer, a content of 2.0 - 6.0% aluminum is advantageous. These
two alloy variants can be used, for example, as components for
the exhaust-gas line of a combustion engine or for steam
boilers, superheaters, turbines and other parts of a power
plant.
A preferred aluminum range is in particular the range from 2.5%
to 5.0%, which is still characterized by good processability.
In the already indicated alloy, the following elements may be
additionally used individually or in combination: .
La 0.02 to 0.3%
Ti 0.01 to 0.5%
Mg 0.0001 to 0.07%
Ca 0.0001 to 0.07%
0.002 to 0.03%
Ni/Co/Cu 0.01 to 3%
up.to 0.005%.
The contents of the elements that can be additionally introduced
in the alloy may be adjusted as follows: Mg 0.0001 to 0.05%, Ca
0.0001 to 0.03%, P 0.002 to 0.03%.
Furthermore, the alloy (in % by weight) may contain one or more
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of the elements Ce, La, Pr, Ne, Sc, Y, Zr or Hf in contents of
0.02 - 0.3%.
If necessary, the alloy (in % by weight) may contain one or more
of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf in contents of 0.02
- 0.2%.
To achieve the desired effects, the Nb content is 0.3 to 1.0%
and the Si content 0.15 to 0.5%.
If necessary, the element tungsten may be replaced entirely or
partly by at least one of the elements Mo or Ta.
If necessary, the alloy may also even contain max. 0.2% V and/or
max. 0.005% S. In this case the oxygen content should not be
greater than 0.01%.
If necessary, the alloy may also even contain max. 0.003% boron.
Furthermore, the alloy should have a maximum of 0.01% of the
following elements respectively: Zn, Sn, Pb, Se, Te, Bi, Sb.
Components/semifinished products that on the one hand consist of
the cited alloy composition and on the other hand have been
produced by the Inventive method may preferably be used as
interconnector in a fuel cell or as material in a component,
such as a reformer or a heat exchanger in an ancillary aggregate
of the fuel cell.
Alternatively, the possibility also exists of using the
component/semifinished product produced according to the
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inventive method or the alloy itself as a structural element in
the exhaust-gas line of a combustion engine or for steam
boilers, superheaters, turbines and other parts of a power plant
or in the chemical process industry.
By means of the inventive method, Laves phases, by virtue of the
thermomechanical treatment, can be precipitated purposefully and
in fine dispersion at the dislocations of the microstructure in
alloys produced by melting metallurgy.
The details and the advantages of the invention will be
explained in more detail in the following examples.
In the following, the inventive method steps will be subjected
to closer examination.
The first step for the thermomechanical treatment of an iron-
chromium alloy precipitating Laves phases and/or Fe-containing
particles and/or Cr-containing particles and/or Si-containing
particles and/or carbides must be annealing above the solution
annealing temperature, so that the Laves phases and/or Fe-
containing particles and/or Cr-containing particles and/or Si-
containing particles and/or carbides are dissolved and are
available for precipitation ,in the subsequent thermomechanical
treatment. The solution annealing temperature is alloy-specific,
but preferably lies above 1050 C for a period of longer than 6
minutes or above 1060 C for longer than 1 minute, followed by
quenching in. stationary protective gas or air, moving (blown)
protective gas or air or in water. The exact temperature control
above this solution annealing temperature is not determining for
the characteristics. The annealing may be carried out in air or
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under protective gas. It should lie under the melting
temperature, preferably < 1350 C. Also, for cost reasons, the
annealing times should preferably be < 24 hours, but may also be
longer depending on performance. The solution annealing follows
quenching in stationary protective gas or air, moving (blown)
protective gas or air or in"water, during which only little
Laves phase is newly formed.
In addition, care is to be taken that, especially for thicker-
walled components, all parts of the component reach the required
minimum annealing time at the specified temperature. This is to
be considered in the determination of the starting point of the
annealing time.
In a second step, an elevated dislocation density must be
introduced into the material. Elevated dislocation densities
have worked microstructure or recovered microstructure, wherein
the dislocations there are arranged at small-angle grain
boundaries.
The second step must therefore be working, so that the
dislocations are introduced into the material, which then, in
the subsequent annealing treatment, ensure a homogeneous
dispersion of the Laves phases and/or Fe-containing particles
and/or Cr-containing particles and/or Si-containing particles
and/or carbides.
This deformation may be cold working, but also hot working,
wherein it must be ensured during hot working that the
microstructure is not already recrystallized during rolling.
This is achieved by restricting the deformation range for the
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last working and the temperature at which this is carried out.
For deformations above 1000 C, the material already tends to
recrystallization or recovery during working, so that the
working must preferably be carried out below 1000 C. At
temperatures below 500 C, exist in the range of the
embrittlement that occurs in ferrites at 475 C. There this has a
smaller elongation and an elevated working resistance, which
makes working less advantageous and reduces the economic
benefit.
Precipitates smaller than a certain size are less effective
(see, for example, "Bilrgel", page 141). Therefore the
dislocation density generated by the deformation should not be
too high, since then very many precipitates are indeed formed
but are too fine, and the excess dislocations can move freely
and in this way the preworking becomes harmful. This means that
preferably the greatest deformation is 90% for the part of the
hot working 1000 C and 90% for the cold working.
During working in the range of 20 to 99%,annealing between 950 C
and 1050 C may cause recovery of the microstructure. Thereby the
dislocation density is reduced, so that the positive effect on
the dispersion of the Laves phases and/or Fe-containing
particles and/or Cr-containing particles and/or Si-containing
particles and/or carbides is established again.
The one possibility of introducing the Laves phases and/or Fe-
containing particles and/or Cr-containing particles and/or Si-
containing particles and/or carbides into the worked material is
to make the needed components from the semifinished product and
then to bring the component that has been made to the
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application temperature between 550 C and 1000 C by heating at
0.1 C/min to 1000 C/rain. During heating, the Laves phases and/or
Fe-containing particles and/or Cr-containing particles and/or
Si-containing particles and/or carbides are precipitated as a
fine dispersion in the microstructure. The fine dispersion is
generated by nucleation in the lower temperature range, followed
by some growth of the nuclei at the higher temperatures.
Therefore the heating rate should not be slower than 1000 C/min,
because otherwise the time for this process is too short.
Heating rates slower than 0.1 C/min are uneconomical.
A second possibility is a separate heat treatment of the
material. For this purpose, the worked semifinished
product/component is subjected for a time between train and tin, to
a heat treatment in the temperature range between 550 C and
1060 C under protective gas or air, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water or for heat treatments up to 800 C is
quenched in the oven, wherein
0 (6740/Ta-9.216)
Ta = and
1 0 (17960/Ta-15.72)
tmax = Ta = where Ta = T +273.15,
with indication of tra,n and tmax in minutes and heat treatment
temperature T in C. In this connection, the desired component
can be made before or after this heat treatment.
In the annealing steps, care is to be taken that, especially for
thicker-walled semifinished products/components, all parts of
the component reach the required minimum annealing time at the
specified temperature. This is to be considered in determination
of the starting point of the annealing time. Likewise, care is
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to be taken that no region of the semifinished product/component
exceeds the required maximum annealing time.
Times shorter than tmin are not sufficient for formation of the
Laves phases and/or Fe-containing particles and/or Cr-containing
particles and/or Si-containing particles and/or carbides. For
times longer than tõx, the danger exists of too great coarseness
of the precipitates, whereby the particles can no longer
contribute markedly to the creep strength. For times longer than
tõ, the possibility exists in the upper temperature range of
550 C and 1060 C that a recovered microstructure will be formed,
which certainly can still be effective. However, with increasing
recovery, the dislocation density is further reduced, so that
the dispersion of the precipitates becomes increasingly
inhomogeneous and the positive effect on the creep strength
ultimately vanishes. In addition, at the lower temperatures in
the range of 550 C and 1060 C, times longer than tiriax are
uneconomical.
The annealing step may be carried out under protective gas
(argon, hydrogen and similar atmospheres with reduced oxygen
partial pressure). For economic reasons, the quenching step is
carried out in stationary protective gas or air, moving (blown)
protective gas or air or in water, while oven quenching should
be avoided in particular for temperatures above 800 C but is
also possible at temperatures < 800 C.
The oxidation resistance and the thermal expansion coefficient
of the material are determined via the chromium content. The
oxidation resistance of the material is based on the formation
of a closed chromium oxide layer. Below 12%, iron-containing
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oxides, which impair the oxidation resistance, are formed
increasingly, especially at higher operating temperatures. The
chromium content is therefore adjusted to. 12%. Above 30%
chromium, the processability of the material and its usability
are impaired by increased formation of embrittling phases,
especially the sigma phase. The chromium content is therefore
limited to ._30%. The expansion coefficient decreases with
increasing chromium content.
Especially for use in a fuel cell, therefore, the expansion
coefficient can be adjusted in a range that matches the ceramics
in the fuel cell. These are chromium contents around 22 to 23%.
=
For other applications, however, for example for reformers or in
power plants, this restriction does not exist.
The addition of manganese brings about formation of a chromium-
manganese spinel on the chromium oxide layer, which is formed on
the material at low aluminum contents below 2%. This chromium-
manganese spinel reduces the chromium volatilization and
improves the contact resistance. A manganese content of at least
0.001% is necessary for this. More than 2.5% manganese impairs
the oxidation resistance by formation of a very thick chromium-
manganese spinel layer.
Niobium, molybdenum, tungsten or tantalum can participate in the
formation of precipitates in iron-containing alloys, such as,
for example, carbides and/or the M in the Fe2(M, Si) or Fe7(M,
Si)6 Laves phases. Molybdenum, tungsten or tantalum are also good
solid-solution hardeners and thus contribute to the improvement
of the creep strength. In this connection the lower limit is
determined in each case by the fact that a certain content must
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be present in order to be effective, while the upper limit is
determined by the processability. Thus the preferred ranges are,
for
Nb 0.1 - 2%
W: 0.1 - 5%
W may also be replaced entirely or partly by Mo and/or Ta: 0.1 -
5%.
Silicon can participate in the formation of precipitates in
iron-containing alloys, for example in the Fe2(M, Si) or Fe7(M,
Si)6 Laves phases. It favors the increased precipitation and
stability of these Laves phases and in this way contributes to
the creep strength. During formation of the Laves phases, it is
completely bound in these. Thus the formation of a silicon
dioxide layer no longer takes place under the chromium oxide
layer. At the same time, the incorporation of M in the oxide
layer is reduced, whereby the negative influence of M on the
oxidation resistance is prevented. At least 0.05% Si must be
present for the desired effect to occur. If the content of Si is
too high, the negative effect of the Si may reappear. The Si
content is therefore limited to 1%.
Aluminum in contents below 1% impairs the oxidation resistance,
since it leads to internal oxidation. However, an aluminum
content higher than 1% leads to formation, under the chromium
oxide layer, of an aluminum oxide layer, which is not
electrically conductive and thus reduces the contact resistance.
Therefore the aluminum content is limited to 0.5% when a
chromium oxide former is desired or its oxidation resistance is
sufficient. An example of this is, for example, for the
application as interconnector plate. However, a certain aluminum
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content of at least 0.001% is necessary for deoxidation of the
melt. If no conductive oxide is necessary and at the same time
the requirement of much higher oxidation resistance than is
given by a chromium oxide layer is still required, the alloy may
form a closed aluminum oxide layer by a content of aluminum of
at least 2% (DE 101 20 561). Aluminum contents above 6.0% lead
to processing problems and thus to increased costs.
Carbon leads to carbide precipitates and thus contributes to the
creep strength. The carbon content should be < 0.1%, in order
not to impair the processability. However, it should be >
0.002%, so that an effect can occur.
The nitrogen content should be 0.1% maximum, in order to avoid
formation of nitrides, which impair the processability. It
should be higher than 0.002%, in order to assure the
processability of the material.
The contents of sulfur should be made as low as possible, since
this interface-active element impairs the oxidation resistance.
A maximum of 0.01% S is therefore stipulated.
Oxygen-affine elements such as Ce, La, Pr, Ne, Sc, Y, Zr, Hf
improve the oxidation resistance by reducing the oxide growth
and improving the adherence of the oxide layer. A minimum
content of 0.02% of one or more of the elements Ce, La, Pr, Ne,
Sc, Y, Zr, Hf is practical in order to obtain the oxidation-
resistance-increasing effect of the Y. For cost reasons, the
upper limit is set at 0.3% by weight.
As with every oxygen-affine element, titanium is bound in the
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oxide layer during oxidation. In addition, it also causes
internal oxidation. However, the resulting oxides are so small
and finely dispersed that they cause hardening of the surface
and thus prevent swelling of the oxide layer and inclusion of
metallic zones during the oxidation (see DE 10 2006 007 598 Al).
-These swellings are unfavorable, since the resulting cracks
cause an increase of the oxidation rate. Thus Ti contributes to
the improvement of the oxidation resistance. For effectiveness
of the Ti content, at least 0.01% Ti must be present, but not
more than 0.5%, since this does not improve the effect further
but increases the costs.
The content of phosphorus should be lower than 0.030%, since
this interface-active element impairs the oxidation resistance.
Too low P content increases the costs. The P content is
therefore 0.002%.
The contents of magnesium and calcium are adjusted in the spread
.ranges of 0.0001 to 0.05% by weight and 0.0001 to 0.03% by
weight respectively.
It has been found that cobalt contents of 3% and higher impair
the oxidation resistance. For cost reasons, =the lower limit is
set at 0.01% by weight. For nickel and copper, the same applies
as for cobalt.
Boron is limited to max. 0.005%, since this element reduces the
oxidation resistance.
2
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Specific aspects of the invention relate to:
(1) a method for producing a component, from an iron-
chromium alloy precipitating at least one of Laves phases,
Fe-containing particles, Cr-containing particles, Si-containing
particles and carbides, consisting of (in weight -%):
Cr 12 - 30%
Mn 0.001 - 2.5%
Nb 0.1 - 2%
W. 0.1 - 5%
Si 0.05 - 1%
0.002 - 0.1%
0.002 - 0.1%
max. 0.01%
Fe remainder
as well as the usual melting-related impurities,
wherein a mechanical deformability at room temperature of >13%
is obtained, measured as plastic elongation in the tension test
in that a semifinished product produced from the alloy is
subjected to a thermomechanical treatment, wherein in a first.
step the alloy is solution annealed at temperatures the
solution-annealing temperature which are 1050 C for longer
than 6 minutes or 1060 C for longer than 1 minute, followed
by quenching in stationary protective gas or air, moving
(blown) protective gas or air or in water, in a second step
mechanical working of the semifinished product in the range
from 0.05 to 99% is performed and in a subsequent step at least
one of Fe2(M, Si) Laves phases, Fel(M, Si)6 Laves phases,
Fe-containing particles, Cr-containing particles, Si-containing
particles and carbides are precipitated purposefully and in
finely dispersed form, by the fact that the component made from
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the worked semifinished product is brought to an application
temperature between 550 C and 1000 C by heating at 0.1 C/min to
1000 C/min;
(2) a method for producing of a component from an iron-
chromium alloy precipitating at least one of Laves phases,
Fe-containing particles, Cr-containing particles, Si-containing
particles and carbides, consisting of (in weight -%):
Cr 12 - 30%
Mn 0.001 - 2.5%
Nb 0.1 - 2%
0.1 - 5%
Si 0.05 - 1%
0.002 - 0.1%
0.002 - 0.1%
S max. 0.01%
Fe remainder
as well as the usual melting-related impurities,
wherein a mechanical deformability at room temperature of >13%
is obtained, measured as plastic elongation in the tension test
in that a semifinished product produced from the alloy is
subjected to a thermomechanical treatment, wherein in a first
step the alloy is solution annealed at temperatures the
solution-annealing temperature which are 1050 C for longer
than 6 minutes or 1060 C for longer than 1 minute, followed
by quenching in stationary protective gas or air, moving
(blown) protective gas or air or in water, in a second step
mechanical working of the semifinished product in the range
from 0.05 to 99% is performed and in a subsequent step at least
one of Fe2(M, Si) Laves phases, Fe7(M, Si)6 Laves phases,
Fe-containing particles, Cr-containing particles, Si-containing
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particles and carbides are precipitated purposefully and in
finely dispersed form by the fact that the worked semifinished
product is subjected for a time between tgar, and tõx to a heat
treatment in the temperature range between 550 C and 1060 C
under protective gas or air, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water or for heat treatments up to 800 C is
quenched in the oven, wherein tiõ,, and tmax are calculated
according to the following formulas:
0 (6740/M-9.216) 10 (17960/M-15.72)
tzun = Ta = and tõ. = Ta = where
Ta = T +273.15,
and wherein the desired component is made before or after this
heat treatment;
(3) the method according to (1) or (2), wherein only
little or even no Fe2(M, Si) Laves phases, Fe7(M, Si)6 Laves
phases, Fe-containing particles, Cr-containing particles,
Si-containing particles, carbides are still present in the
semifinished product after solution annealing at temperatures
the solution-annealing temperature, followed by quenching in
stationary protective gas or air, moving (blown) protective gas
or air or in water, in the initial state before deformation; .
(4) the method according to any one of (1) to (3),
wherein the working of the semifinished product takes place by
hot working;
(5) the method according to any one of (1) to (4),
wherein the hot working of the semifinished product begins with
a starting temperature > 1070 C, wherein the last 0.05 to 90%
of mechanical deformation is applied between 1000 C and 500 C;
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(6) the method according to any one of (1) to (5),
wherein the hot working of the semifinished product begins with
a starting temperature > 1070 C, wherein the last 0.05 to 95%
of mechanical deformation is applied between 1000 C and 500 C;
(7) the method according to any one of (1) to (6),
wherein the hot working of the semifinished product begins with
a starting temperature > 1070 C, wherein the last 0.05 to 90%
of mechanical deformation is applied between 1000 C and 500 C;
(8) the method according to any one of claims (1) to (7),
wherein the hot working of the semifinished product is followed
by cold working;
(9) the method according to any one of (1) to (8),
wherein the working of the semifinished product is carried out
by cold working;
(10) the method according to (9), wherein the degree of .
cold working of the semifinished product is 0.05 to 99%;
(11) the method according to (9) or (10), wherein the cold
working of the semifinished product is 0.05 to 95%;
(12) the method according to any one of (9) to (11),
wherein the cold working of the semifinished product is 0.05
to 90%;
(13) the method according to any one of (1) to (12),
wherein the mechanical working of the semifinished product
is 20 to 99% and then the worked semifinished product is
subjected for a time between t.,õ and tiaax to a heat treatment in
the temperature range between 950 C and 1060 C under protective
gas or air, followed by quenching in stationary protective gas
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or air, moving (blown) protective gas or air or in water and
after this the desired component is made with
(6740/Ta-9.216)
tmln Ta = and
10 (17960/Ta-15.72)
tmax = Ta where Ta = T +273.15
5 and indication of trlin and tmax in minutes and of heat-treatment
temperature T in C;
(14) the method according to any one of (1) to (13),
wherein the alloy additionally contains (in % by weight) 0.02
to 0.3% La;
10 (15) the method according to any one of (1) to (14),
wherein the alloy additionally contains (in % by weight) 0.01
to 0.5% Ti;
(16) the method according to any one of (1) to (15),
wherein the alloy additionally contains 0.02 to 0.3% of one or
more of the elements Ce, Pr, Ne, Sc, Y, Zr or Hf;
(17) the method according to any one of (1) to (16),
wherein the alloy additionally contains (in % by weight) 0.001
to 0.5% Al;
(18) the method according to any one of Cl) to (17),
wherein the alloy additionally contains (in % by weight) 2.0
to 6.0% Al;
(19) the method according to (18), wherein the alloy
additionally contains (in % by weight) 2.5 to 5.0% Al;
(20) the method according to any one of (1) to (19),
wherein the alloy additionally contains one or more of the
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=
elements 0.0001 to 0.07% Mg, 0.0001 to 0.07% Ca,
0.002 - 0.03% P;
(21) the method according to any one of (1) to (20),
wherein the alloy further contains 0.01 to 3.0% of one or more
of the elements Ni, Co or Cu;
(22) the method according to any one of (1) to (21),
wherein the alloy further contains up to 0.005% B;
(23) the method according to any one of (1) to (22),
wherein the iron-chromium alloy, thermomechanically treated and
precipitating Laves phases in finely dispersed form, of the
following composition containing (in % by weight);
Cr 12 - 30%
Mn 0.001 - 2.5%
Nb 0.1 - 2%
W 0.1 - 5%
Si 0.05 - 1%
0.002 - 0.03%
0.002 - 0.03%
max. 0.005%
Fe remainder
as well as the usual melting-related impurities;
(24) the method according to any one of (1) to (23),
wherein the alloy additionally contains (in % by weight) 0.02.
to 0.2% of the element La;
(25) the method according to any one of (1) to (24),
wherein the alloy additionally contains (in % by weight) 0.02
to 0.2% Ti;
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(26) the method according
to any one of (1) to (25),
wherein the alloy additionally contains (in % by weight) 0.02
to 0.2% of one or more of the elements Ce, Pr, Ne, Sc, Y, Zr
or Hf;
(27) the method according to
any one of (1) to (26),
wherein the alloy additionally contains (in % by weight) one or
more of the elements 0.0001 - 0.05% Mg, 0.0001 - 0.03% Ca,
0.002 - 0.03% P;
(28) the method according to any one of (1) to (27),
wherein the alloy further contains (in % by weight) up
to 0.003% B;
(29) the method according to any one of (1) to (28),
wherein (in % by weight) the Nb content is 0.3 to 1.0% and the
Si content is 0.15 - 0.5%;
(30) the method according to
any one of (1) to (29),
wherein the alloy contains (in % by weight) max. 0.2% V and/or
max. 0.005% S;
(31) the method according
to any one of (1) to (30),
wherein the alloy contains (in % by weight) max. 0.01% 0;
(32) the method according to
any one of (1) to (31),
wherein the alloy contains (in % by weight) max. 0.01% of each
of the elements Zn, Sn, Pb, Se, Te, Bi and Sb respectively;
(33) the method
according to any one of (1) to (32),
wherein the semifinished product is formed by sheet, strip,
bar, forging, pipe or wire;
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(34) the method according to any one of (1) to (33),
wherein the heat treatment is carried out only after finishing
of the component;
(35) use of a component produced according to any one of
(1) to (34), as interconnector in a fuel cell;
(36) use of a component produced according to any one of
(1) to (3)4, as material in a component, or in an ancillary
aggregate of a fuel cell;
(37) the use according to (36), wherein the component is a
reformer or a heat exchanger;
(38) use of a component produced according to any one of
(1) to (34), in the exhaust-gas line of a combustion engine;
and
(39) use of a component produced according to any one of
(1) to (34), for steam boilers, superheaters, turbines and
other parts of a power plant or in the chemical process
industry.
The subject matter of the invention will now be explained in
more details on the basis of exemplary embodiments.
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The analyses of the batches used for the following examples are
presented in Table 1. These batches were melted in the arc
furnace in an amount of approximately 30 metric tons, thereafter
cast in a ladle and subjected to a decarburization and
deoxidation treatment as well as to a vacuum treatment in a VOD
system and cast to ingots. These were then hot-rolled and,
depending on final thickness, cold-rolled with intermediate
annealing steps. After the hot-rolling, the oxide layer was
removed by pickling.
A material with an analysis as indicated in Table 1 precipitates
mainly Fe2(M, Si) or Fe7(M, Si)6 Laves phases and, in much
smaller contents, carbides.
Example 1
In this example, material from the batch 161061 listed in Table
1 was hot-rolled to 12 mm thick sheet after solution annealing
above 1070 C for a period of longer than 7 minutes followed by
quenching in stationary air, wherein the mechanical working was
begun with a start temperature > 1070 C and the last 78% of
mechanical deformation was applied by rolling between 500 C and
1000 C.
Figure 1 shows the typical appearance of a microstructure
deformed in this way. In the microsections etched by means of
electrolytic etching with oxalic acid, it can be clearly seen
that only little Laves phase has been precipitated in
microscopically visible form.
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When the material formed in this way is then annealed at 1075 C
for 20 minutes with quenching in stationary air, a
microstructure is obtained with only few precipitates of Laves
phase and a grain size of approximately 137 m (Figure 2), which
is a typical coarse-grained microstructure.
When a creep test as described above is performed on this
material with an initial stress of 35 MPa at a temperature of
750 C, the specimen breaks after 12.8 hours at an elongation A
of 69.8%. (Table 2). At room temperature, the material annealed
in this way has an elongation of 35%, which is a very good value
for a ferrite.
If, in contrast, from the hot-rolled material, which is
synonymous with preworking, a specimen for a creep test is made
as simulation for a component and this is then heated at
approximately 60 C/minute to an application temperature of 750 C
and then a creep test is performed with an initial stress of 35
MPa at a temperature of 750 C, the specimen surprisingly breaks
only after 255 hours at an elongation A of 29%, which means a
prolongation of the time to break by a factor of 20. Making the
component is very easily possible, since the hot-worked
condition, as was described above, has an elongation of 19% in
the tension test at room temperature, which is a good value and
makes the material readily processable.
This example clearly shows that the microstructure with the
preworking and the coarse-grained microstructure is superior
with respect to time to break or creep strength, which
contradicts the state of the art as described in "BUrgel" pages
196 to 199 Table 3.7.
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Example 2
In this example, annealing steps were performed on the hot-
rolled material from Example 1 for 20 minutes each between 600 C
and 1000 C or for some temperatures also for 240 or 1440 minutes
(see Table 3 for trnin and tmaõ according to Equation 1 and 2) in
air, followed by quenching in stationary air. After the heat
treatment, specimens were made from the sheet and then the creep
test was performed with a stress of 35 MPa at 750 C as described
above. The results are compiled in Table 3.
After 20 minutes at 1000 C, a time to break of only 10.4 hours
is reached at an elongation A of 79.5%. After 20 minutes at
600 C to 950 C, a time to break of longer than 100 hours, at
least 7 times longer, is achieved at an elongation A of greater
than 22.7%. The longest time to break for annealing steps of 20
minutes is achieved at 850 C, with 564 hours. The longest time
to break for annealing steps of 240 minutes is achieved at
800 C, with 396 hours. After 1440 minutes at 700 C, a time to
break of 645 hours is even reached. Figure 3 shows the
microstructure after the various annealing steps for 20 minutes.
The microstructures in Figure 3 are not globularly
recrystallized. Up to 850 C (the maximum of the time to break),
the microstructure has the typical appearance of a deformed
microstructure. Starting from approximately 900 C, recovery can
be clearly recognized, but this means that the dislocation
density is still Increased compared with a globularly
recrystallized microstructure. In a recovered microstructure,
the dislocations have become partly reordered at small-angle
grain boundaries. This has an effect similar to that of
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preworking. In the microsections etched by means of electrolytic
etching with oxalic acid, it can be clearly recognized that the
Laves phase is precipitated in microscopically visible form
starting from approximately 750 C, wherein it is precipitated
increasingly more densely and more homogeneously up to 850 C
(the maximum of the time to break). From approximately 900 C on,
small-angle grain boundaries or grain boundaries can also be
recognized markedly besides the precipitates in the grain, thus
assuring jaggedness of the small-angle grain boundaries or grain
boundaries, which corresponds to measure 2 for increasing the
creep strength (see above). At 1000 C, very large grains have
recognizably formed due to advancing recovery, such that the
dislocation density is greatly reduced and so no further
increase of the time to break occurs. The maximum of the time to
break occurs in the deformed microstructure with dense
homogeneously precipitated Laves phases.
At room temperature, the sheet annealed for 20 minutes at all
temperatures between 600 C and 950 C has an elongation of at
least 13%, which is still to be regarded as satisfactory for a
ferritic alloy and makes the material processable. The
elongation is smallest in the range of 700 C to 800 C and is
improved at the lower or higher annealing temperatures
respectively, because at the lower temperatures Laves phase is
certainly already precipitated but is not yet microscopically
visible and therefore has a smaller proportion by volume, but in
return is very finely dispersed. At the higher temperatures, a
larger proportion by volume is precipitated, but in return is
somewhat coarser and recognizable at the small-angle grain
boundaries and grain boundaries.
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The annealing at 1000 C for an annealing time of 20 minutes
exceeds trnax = 19.6 minutes. Thus it is not in the range of the
invention and is used as reference. Also, the time to break is
only 10.4 hours. The annealing time of 20 minutes at the
temperatures between 600 C and 950 C lies in the inventive range
between train and trõ Accordingly, the time to break was clearly
prolonged according to the invention by more than a factor of 7
compared with the coarse-grained, globularly recrystallized
condition from Example 1, which is obtained after annealing at
1075 C/20 minutes followed by quenching in stationary air.
Example 3
In this example, material from the batch 161061 listed in Table
1 was hot-rolled to 12 mm thick sheet after solution annealing
above 1070 C for a period of longer than 7 minutes followed by
quenching in stationary air, wherein the working was begun with
a start temperature > 1070 C and the last 60% of mechanical
working was applied by rolling between 1000 C and 500 C.
If the sheet worked in this way is then annealed industrially in
the continuous furnace at 920 C for 28 minutes in air and
quenched in stationary air, a tension specimen made from this
material has a time to break of 391 hours at an elongation A of
38% (Table 4) in the creep test with an initial stress of 35 MPa
at a temperature of 750 C. The microstructure is not globularly
recrystallized but instead is recovered. It has precipitates in
the grain and at the small-angle grain boundaries or grain
boundaries (Figure 4). The time to break is 30 times the time
achieved in Example 1 after annealing at 1075 C for 20 minutes
with a globularly recrystallized coarse-grained microstructure
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with a grain size of 137 m. The annealing at 920 C for an
annealing time of 28 minutes lies in the inventive range between
train = 0.32 minutes and tmax = 162.6 minutes.
At room temperature, the sheet treated in this way has a very
good elongation of 18%, an offset yield strength of 475 MPa and
a tensile strength of 655 M2a (see Table 4), which makes the
material readily workable.
Example 4
In this example, material from batch 161061 and batch 161995 was
cold-rolled to 1.5 mm thick sheet after solution annealing at
above 1070 C for a period longer than 7 minutes followed by
quenching in blown air and hot rolling as well as removal of the
oxide layer, wherein cold working of 53% was applied.
Subsequently annealing at 1050 C was carried out for 3.4 minutes
under protective gas in the continuous furnace with subsequent
quenching in the cold stream of protective gas. Thereafter both
batch 161061 (Figure 5) and batch 161995 exhibit a recovered
microstructure with elongated grains (Figure 7) and
precipitation of Laves phase, although much less than
recognizable in Figure 4. Thereafter part of the material was
annealed once again at 1050 C for 20 minutes under air with
subsequent quenching in stationary air. After this, both batches
were globularly recrystallized, batch 161061 with a grain size
of 134 m (Figure 6) and batch 161995 with a grain size of 139
m. Only slight precipitated Laves phase can still be found.
Tables 5a and 5b show the results of the creep tests and of the
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33
tension tests at .room temperature. After the annealing at 1050 C
for 3.4 minutes, batch 161061 has a time to break of 25.9 hours
at an elongation A of 50% in a creep test at 750 C with an
initial load of 35 MPa, and after additional annealing at 1050 C
for 20 minutes, which produces very coarse grain, a time to
break of only one third, 7.9 hours, at an elongation A of 83%.
Similarly, batch 161995 has a time to break of 33.5 hours
at an elongation A of 89% in
a creep test at 750 C With an initial load of 35 MPa, and after
additional annealing at 1075 C for 20 minutes, which produces
very coarse grain, a time to break of only one third, 7.9 hours,
at an elongation A of 92%. The elongation of 28% in the tension
= test at room temperature for batch 161061 for 1050 C and 3.4
minutes of annealing time and of 26% for batch 161995 is very
good for a ferrite, which makes the material very readily
workable. For the coarse-grained structure, it is even higher,
with 31% for batch 161061 and 29% for batch 161995.
= =
This shows the influence of the annealifig time at temperatures
around 1050 C. In short-time annealing steps of a few minutes,
dislocations (deformation) and adequate Laves phase are present
in the material, which in this example has the consequence of a
3 to 4 times longer time to break in the creep test. For longer
annealing steps, the Laves phase is sufficiently dissolved, as
batch 161061 shows, and the microstructure recrystallizes
globularly with correspondingly short times to break in the
creep test.
=
The annealing at 1050 C for 20 minutes lies with an annealing
time of 20 minutes above t.. - 6.0 minutes. Thus it does not
fall within the range of the invention and is used as reference,
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just as the annealing at 1075 C for 20 minutes. The annealing at
1050 C for 3.4 minutes lies with an annealing time of 3.4
minutes in the inventive range between tmin = 0.32 minutes and
tmax = 6.0 minutes and according to the invention exhibits a
clearly improved time to break in the creep test.
Example 5
In this example, material from the batch 161061 was hot-rolled
to 12 mm thick sheet after solution annealing above 1070 C for a
period of longer than 7 minutes followed by quenching in
stationary air, wherein the working was begun with a start
temperature > 1070 C and the last 70% of mechanical deformation
was applied by rolling between 1000 C and 500 C.
When the material worked in this way is then subjected to
solution annealing at 1075 C for 22 minutes with quenching in
stationary air, a very coarse-grained microstructure is obtained
with only few precipitates of Laves phase and a grain size of
approximately 134 to 162 m (Figure 9). When a creep test is
performed on this material with an initial stress of 40 MPa at a
temperature of 700 C, the specimen breaks after 228 hours at an
elongation A of 51%. (Table 6) When the creep test is performed
at 60 MPa, the specimen breaks after 8.1 hours, at an elongation
A of 43%. At room temperature, the material annealed in this way
has an elongation of 35%, which is a very good value for a
ferrite.
If the material solution annealed at 1075 C for 22 minutes is
additionally subjected to annealing for 4 hours at 700 C with
subsequent quenching in stationary air, Laves phase dispersed in
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the microstructure is precipitated. (Figure 10). When the creep
test is then performed at 700 C with an initial stress of 40
MPa, the specimen already breaks after 104 hours, at an
elongation A of 72.6%, therefore a much shorter time than after
the solution annealing at 1075 C for 22 minutes. When the creep
test is performed at 60 MPa, the specimen breaks after 6.3
hours, at an elongation A of 63%, therefore also after
substantially shorter time than after the solution annealing at
1075 C for 22 minutes.
This is the proof that the precipitation of the Laves phase(s)
must take place in a microstructure with elevated dislocation
density, therefore in a worked or recovered microstructure, in
order to achieve prolongation of the time to break.
Precipitation in a solution annealed microstructure has exactly
the opposite effect, namely shortening of the time to break. The
cause of this is the more homogeneous dispersion of very fine
precipitates in the case of precipitation in a microstructure
with elevated dislocation density, or in other words a deformed
or recovered microstructure, in comparison with precipitation in
a dislocation-poor coarse-grained microstructure.
Example 6
In this example, as in Example 2, annealing steps were performed
on the hot-rolled material from Example 1 for 20 minutes each
between 750 C and 1000 C or for some temperatures also for 120
minutes, 240 minutes, 480 minutes, 960 minutes, 1440 minutes or
5760 minutes (see Table 7 for train and tmax according to Equation
1 and 2) in air, followed by quenching in stationary air. After
the heat treatment, specimens were made from the sheets and then
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the creep test was performed with a stress of 40 MPa at 750 C as
described above. The higher stress in comparison with Example 2
was chosen for shortening of the test time. The objective was to
find heat-treatment times suitable for the annealing steps. The
results are compiled in Table 7.
After 20 minutes at 1000 C, a time to break of only 8.8 hours is
reached at an elongation A of 78.7%. In Example 2, after 20
minutes at 1000 C and a creep test at 750 C and 35 MPa, a time
to break comparable with that after solution annealing at 1075 C
for 20 minutes with quenching in stationary air was reached, and
so this value can be taken as reference for the time to break of
the solution annealed condition. Inventive variants should also
exceed this break time once again by a factor of at least 1.5.
After 20 minutes at 750 C to 900 C, a time to break of longer
than 100 hours, at least 10 times longer, is achieved at an
elongation A of greater than 27%. The longest time to break for
annealing steps of 20 minutes is achieved at 850 C with 296
hours. The longest time to break for annealing steps of 120
minutes is achieved at 800 C with 227 hours. The longest time to
break for annealing steps of 240 minutes is achieved at 750 C
with 182 hours, but in this connection no value exists for
700 C. The longest time to break for annealing steps of 480
minutes is achieved at 800 C with 169 hours. For 960 minutes,
only one time to break was determined, for 750 C, with a value
of 139 hours at an elongation of 24.2%. After 1440 minutes and
5760 minutes at 750 C and 800 C, only times to break clearly
shorter than the maximum times to break achieved at these
temperatures are still achieved. At 800 C, for example, the time
to break after a treatment time of 480 minutes drops from 169
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hours to a value of 46 hours after a treatment time of 1440
minutes, to a value of 17.5 hours after a treatment time of 5760
minutes, although this is still in the inventive range. A
further prolongation of the treatment time should shorten the
time to break further, so that tmax of 7059 minutes is logically
somewhat longer than the time of 5750 minutes. All elongations
for the heat treatment temperatures from 750 to 900 C and times
from 20 minutes to 5760 minutes lie between 24.2% and 43% and
therefore are larger than 18%, as required, in order to avoid
brittle failure. For the microstructure after 20 minutes of
annealing, what was said in Example 2 is applicable, since the
annealing steps are the same. Even at the higher stress of 40
MPa in the creep test, the maximum of the time to break occurs
in the deformed microstructure, with dense homogeneously
precipitated Laves phase.
At room temperature, the sheet annealed for 20 minutes at all
temperatures between 600 C and 900 C in Example 2 has an
elongation of at least 13%, which is still to be regarded as
satisfactory for a ferritic alloy and makes the material
proces sable.
Example 7
In this example the 1.5 mm thick material of batch 161995, which
was annealed after cold working of 53% at 1050 C for 3.4 minutes
under protective gas in the continuous furnace with subsequent
quenching in the stream of cold protective gas, was used once
again. In the same manner, 2.5 mm thick material from batch
161995 was produced by annealing it, after cold working of 40%,
at 1050 C for 2.8 minutes under protective gas in the continuous
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38
furnace with subsequent quenching in the stream of cold
protective gas. Even the 2.5 mm thick material then exhibits a
recovered microstructure with elongated grains '(Figure 11), just
as the material from Example 4 in Figure 7, and precipitation of
Laves phase, albeit clearly less than recognizable in Figure 4.
Part of the material was then annealed once more at 1050 C for
minutes under air with subsequent quenching in stationary
air. After this the material had globularly recrystallized
structure, with a grain size of 108 Km. Only little precipitated
Laves phase is still to be found. The material with 1050 C/2.8
minutes and the material with 1075 C/10 minutes as the last heat
treatment was then rolled with degrees of working between 2.8
and 40%. Thereafter creep tests were performed at 750 C and 35
MPa and tension tests were performed at room temperature. The
results are summarized in Table 8.
After the annealing at 1050 C for 3.4 minutes, batch 161995, in
a creep test at 750 C with an initial load of 35 MPa, had a time
to break of 33.5 hours at an elongation A of 89% and, after the
additional annealing at 1050 C for 10 minutes, which produces
very coarse grain, it had a time to break amounting to only one
third, 10.8 hours, at an elongation A of 50.4%.
If, from the material worked after 1050 C/2.8 minutes, a
specimen for a creep test is'made as simulation for a component
and this is then heated at approximately 60 C/minute to an
application temperature of 750 C and then a creep test is
performed with an initial stress of 35 MPa at a temperature of
750 C, the time to break for degrees of working between 5
and 40% drops to values around the 10 hours with elongations at
break greater than 45%.
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39
If, in contrast, from the material formed after 1050 C/10
minutes, a specimen for a creep test is made as simulation for a
component and this is then heated at approximately 60 C/minute
to an application temperature of 750 C and then a creep test is
performed with an initial stress of 35 MPa at a temperature of
750 C, the time to break for degrees of working between
2.9 and 40% increases to values between 49 and 137 hours, which
means an increase by more than a factor of 4 in the time to
break compared with the material worked after 1050 C/2.8
minutes, wherein a maximum occurs at 10% and the elongations at
break lie between 18.9 and 60%.
From 10% degree of working on, however, the elongation at break
in the tension test at room temperature becomes smaller than 8%,
so that the material is increasingly more poorly processable. In
other words, preferred degrees of forming for cold shaping lie
between 0.05 and 10%. Thus this example shows that the increase
of the elongation at break after working does not occur if the
annealing before working was carried out at too low temperatures
or for too short times. (Here at 1050 C .for 2.8 minutes) The
increase of the elongation at break occurred after annealing was
carried out at > 1050 C for > 6 minutes (here at 1075 C for 7
minutes).
The titles/descriptions of the tables/figures are reproduced as
follows:
Table 1 Composition of the investigated alloy (all values in %
by weight)
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Table 2 Results of the creep tests at 750 C with 35 MPa and of
the tension tests at room temperature for the hot
rolling and the heat treatments in Example 1 for a 12
mm thick sheet. (R: reference according to the state
of the art, I: according to the invention)
Table 3 Results of the creep tests at 750 C with 35 MPa and of
the tension tests at room temperature for the hot
rolling from Example 1 and the heat treatment from
Example 2 for a 12 mm thick sheet. (R: reference
according to the state of the art, I: according to the
invention)
Table 4 Results of the creep tests at 750 C with 35 MPa and of
the tension tests at room temperature for Example 3
for a 12 mm thick sheet. (R: reference according to
the state of the art, I: according to the invention)
Table 5 Results of the creep tests at 750 C with 35 MPa and of
the tension tests at room temperature for Example 4
for a 1.5 mm thick strip. (R: reference according to
the state of the art, I: according to the invention)
Table 6 Results of the creep tests at 700 C and of the tension
tests at room temperature for Example 5 on 12 mm thick
sheet (R: reference according to the state of the art,
I: according to the invention)
Table 7 Results of the creep tests at 750 C with 40 MPa and of
the tension tests at room temperature for the hot
rolling and the heat treatments in Example 6 for a 12
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41
mm thick sheet. (R: reference according to the state
of the art, I: according to the invention)
Table 8 Results of the creep tests at 750 C with 35 MPa and of
the tension tests at room temperature for Example 7
for 1.5 mm to 2.5 mm thick strip from batch 161995.
(R: reference according to the state of the art, I:
according to the invention)
Figure 1 Microstructure of the hot-worked material in Example 1
Figure 2 Microstructure of the hot-worked material in Example 1
after annealing at 1075 C for 20 minutes and quenching
in stationary air, grain size 137 m.
Figure 3 Microstructure of the material in Example 2 after
annealing between 600 C and 1000 C for 20 minutes in
each case and quenching in stationary air.
Figure 4 Microstructure of the material in Example 3 after
annealing at 920 C in the continuous furnace in air
with subsequent quenching in stationary air for 20
minutes in each case and quenching in stationary air.
(etching with V2A pickling fluid)
Figure 5 Microstructure of batch 161061 in Example 4 after
annealing at 1050 C/3.4 minutes under protective gas
in the continuous furnace with quenching in the stream
of cold protective gas.
Figure 6 Microstructure of batch 161061 in Example 4 after
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annealing at 1050 C/3.4 minutes under protective gas
in the continuous furnace with quenching in the stream
of cold protective gas and annealing at 1050 C/20
minutes under air with subsequent quenching in
stationary air, grain size 134 m (etching with V2A
pickling fluid)
Figure 7 Microstructure of batch 161995 in Example 5 after
annealing at 1050 C/3.4 minutes under protective gas
in the continuous furnace with quenching in the stream
of cold protective gas.
Figure 8 Microstructure of batch 161995 in Example 5 after
annealing at 105000/3.4 minutes under protective gas
in the continuous furnace with quenching in the stream
of cold protective gas and annealing at 1075 C/20
minutes under air with subsequent quenching in
stationary air, grain size 139 m
Figure 9 Microstructure of the hot-worked material in Example 5
after annealing at 1075 C for 22 minutes and quenching
in stationary air, grain size 134 m to 162 m
Figure 10 Microstructure of the hot-worked material in Example 5
after annealing at 1075 C for 22 minutes followed by
quenching in stationary air and subsequent annealing
at 700 C for 4 hours followed by quenching in
stationary air. Grain size 136 m.
Figure 11 Microstructure of batch 161995 in Example 7 after
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annealing at 1050 C for 2.8 minutes under protective
gas in the continuous furnace with quenching in the
stream of cold protective gas.
Figure 12 Microstructure of batch 161995 in Example 7 after
annealing at 1050 C for 2.8 minutes under protective
gas in the continuous furnace with quenching in the
stream of cold protective gas, followed by annealing
at 1075 C for 10 minutes under air with subsequent
quenching in stationary air. Grain size 108 m.
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Table 1
Batch Batch
Element 161061 161995
0.007 0.009
<0.002 <0.002
0.015 0.018
Cr 22.9 22.6
Ni 0.30 0.22
Mn 0.43 0.43
Si 0.21 0.24
Mo 0.02 0.02
Ti 0.07 0.06
Nb 0.51 0.49
Cu 0.02 0.02
Fe Remainder Remainder
0.014 0.017
Al 0.02% 0.019
Mg 0.0006 <0.01
Pb <0.001 <0.001
Sn <0.01 <0.01
Ca 0.0002 0.01
V 0.05 0.02
Zr <0.01 <0.01
1.94 1.97
Co 0.04 0.02
La 0.08 0.05
Ce <0.01
0 0.004
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Table 2
Creep test
Heat treatment with 35 MPa at 750 C
Tension test at room temperature
Time tB_ to break Elongation A in % Offset
yield Tensile strength Elongation A 5 in
in hours strength
Rp0.2 RM in MPa yo
0
1075 C/ 20 minutes 12.8 69.8 359
494 35
Hot-worked 511
604 19
0
CO
Component + heating at
0
60 C/minute to application 255 29.0
0
temperature 750 C
46
tv
--]
--.)
Table 3
LD
1
u.)
w
Limit times .
according to Creep test with 35 MPa at
Heat treatment
Tension test at room temperature
equations 1 and 2 in 750 C ,
minutes
. _
Offset yield
Tensile
Temperature Time in Time tB to Elongation A
Elongation A 5
_
tmin tmax strength strength RM in
in C minutes break in hours in %
in % ' n
Rp0.2
MPa
0
I.)
Hot-worked
-1
-1
u.),
.
600 20 27.9 39.2*105 ' 278 37.6
511 604 19 -1
0
.
0
650 29 11.26 31.8*105 286 35.2
527 ' 622 17 I.)
0
700 20 5.00 33.6'104 - 260; 264
22.7; 39.1 537 676 13 H
IV
I
.
0
750 20 2.41 44300 263 430.6
514 707 16 I.)
.
1
800 20 1.25 7059 344 25.0
505 699 13 u.)
.
_
850 20 0.69 1328 564 32.2
484 672 17
900 20 0.40 289 337 25.3
467 ' 638 18
_ .
950 20 0.24 71.2 121:72 31.1;
28.8 451 614 17
1000 20 0.15 19.6 10.4 79.5
n. m. n. m. n. m.
650 240 11.26 31.8*105 - 293 36.2
n. m. n. m. n. m.
700 240 ' 5.00 33.6 104 233 32.8
n. m. n. m. n. m.
750 240 2.41 44300 224 23.2
n. m. n. m. n. m.
_
800 240 1.25 7059 396 43.2
n. m. n. m. n. m.
850 240 0.69 1328 181 35.2
n. m. n. m. n. m.
= ___
900 240 0.40 289 45.6 55.7
n. m. n. m. n. m.
..
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47
950 240 0.24 71.2 10.8 78.7 n.
m. n. m. n. m.
700 1440 5.00 33.6*104 645 30.9 n.
m. n. m. n. m.
For comparison from Example 1
1075 20 12.8 69.8
359 494 35
n. m. = not measured
0
0
CO
I\)
0
0
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Table 4
Limit times
according to Creep test with 35 MPa at
Heat treatment
Tension test at room temperature
equations 1 and 2 in 750 C
minutes
Offset yield
Tensile
Temperature Time in Time tB_ to Elongation A
Elongation A 5
tmin tmax
strength strength RM in
in C minutes break in hours in %
in %
Rp0.2
MPa
920 28 0.32 162.6 391 38
475 655 18
0
0
CO
0
0
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Table 5a
Limit times
according to Creep test with 35 MPa at
Batch 161061
Tension test at room temperature
equations 1 and 2 in 750 C
minutes
Offset yield
Tensile
Time tB_ to Elongation A
Elongation A 5
Heat treatment train tmax
strength strength RM in
break in hours in %in %
Rp0.2
MPa
1050 C/
385 541 27
0.1 6.0 25.9 50
3.4 minutes
380 537 28 0
1050 C/
0.1 6.0 7.9 83
344 494 31
0
20 minutes
co
0
0
Table 5b
Limit times
according to Creep test with 35 MPa at
Batch 161995
Tension test at room temperature
equations 1 and 2 in 750 C
minutes
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'
Offset yield
Tensile
Time tB_ to Elongation A
Elongation A 5
Heat treatment tmin tmax
strength strength RM in
break in hours in %
in %
Rp0.2
MPa
1050 C/
I 0.1 6.0 33.5 89.0 399 538 26
3.4 minutes
1050 C/
R 7.7 92.1 331 475 29
20 minutes
0
0
I.)
-.1
-.1
LO
-.1
0
CO
I\)
0
H
IV
I
0
V
I
I
I\)
LO
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Table 6
Limit times 1
according to
Batch 161061Creep test at 700 C
Tension test at room temperature
equations 1 and 2
in hours
Time tB to
Offset yield Tensile
Stress in - Elongation A
Elongation A 5
Heat treatment trnin tmax break in
strength strength RM in
MPA in %
in %
hours Rp0.2 MPa
n
R 1075 C/22 minutes 40 228 51 367
502 35
0
R 1075 C/22 minutes 60 8.1 43.1 367
502 35 "
-.1
-.1
UJ
1075 C/ 22 minutes +
R 0.083 5601 40 104 72.6
0
co
700 C/ 4 hours
I.)
0
1075 C/ 22 minutes +
H
I \ )
R 0.083 5601 60 6.3 63.5
1
700 C/ 4 hours
0
I.)
1
I.)
UJ
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Table 7a
Limit times
according to Creep test with 40 MPa at
Heat treatmentTension test at room temperature
equations 1 and 2 in 750 C
minutes
Offset yield
Tensile
Temperature Time in Time tE3 to Elongation A
Elongation A 5
tmin tmax
strength strength RM in
in C minutes break in hours in %
in %
Rp0.2
MPa
0
750 20 2.41 44300 128 34.7
514 707 16
0
I.)
800 20 1.25 7059 189 27.2
505 699 13 -A
-A
LO
850 20 0.69 1328 296 32.1
484 672 17 -A
0
CO
900 20 0.40 289 174 35.4
467 638 18 I.)
0
H
1000 20 0.15 19.6 8.8 78.7 n.
m. n. m. n. m. "
1
0
750 120 2.41 44300 150 31.7 n.
nn. n. m. n. m. I.)
1
I.)
u.)
800 120 1.25 7059 227 26.2 n.
m. n. m. n. m.
850 120 0.69 1328 133 29.7 n.
m. n. m. n. m.
900 120 0.40 289 32.7 40 n.
m. n. m. n. m.
n. m. = not measured
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Table 7b
Limit times
according to Creep test with 35 MPa at
Heat treatment
Tension test at room temperature
equations 1 and 2 in 750 C
minutes
Offset yield
Tensile
Temperature Time in Time tB_ to Elongation A
Elongation A 5
tmin tmax
strength strength RM in
in C minutes break in hours in %
in %
Rp0.2
MPa
0
750 240 2.41 44300 182 32.2 n.
m. n. m. n. m.
0
I.)
800 240 1.25 7059 163; 135 27.5; 26.5
n. m. n. m. n. m.
-.1
LO
850 240 0.69 1328 57 33.9 n.
m. n. m. n. m.
0
CO
750 480 2.41 44300 152 43.4 n.
m. n. m. n. m. I.)
0
H
800 480 1.25 7059 169 26.5 n.
m. n. m. n. m. I.)
1
0
850 480 0.69 1328 35 33 n.
m. n. m. n. m. "
1
I.)
u.)
750 960 2.41 44300 139 24.2 n.
m. n. m. n. m.
750 1440 2.41 44300 82 25.5 n.
m. n. m. n. m.
800 1440 1.25 7059 46 46.1 n.
m. n. m. n. m.
750 5760 2.41 44300 54 52.9 n.
m. n. m. n. m.
800 5760 1.25 7059 17.5 50.3 n.
m. n. m. n. m.
n. m. = not measured
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Table 8a
Limit times
Degree of
according to Creep test with 40 MPa at
Heat treatment working in
Tension test at room temperature
equations 1 and 2 in 750 C
%
minutes
Time tB to
Offset yield Tensile
_
Temperature Time in Elongation
A Elongation
tmin tmax break in
strength strength RM
in C minutes in %
A 5 in A
hours
Rp0.2 in MPa
I 1050 3.4 0.1 6.0 0.0 33.5 89
399 538 26 P
0
R 1050 2.8 0.1 6.0 5.0 10.6 45.8
581 616 16 N)
-.1
-.1
R 1050 2.8 0.1 6.0 10.0 10.9 68.6
663 685 8 LO
-.1
0
CO
R 1050 2.8 0.1 6.0 20.0 9.4 71.8
725 745 6 I.)
0
R 1050 2.8 0.1 6.0 40.0 12.0 85.9
811 832 4 H
I.)
1
0
n. m. = not measured
I.)
1
I.)
u.)
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,
Table 8b
Degree of Deformation in Creep test with
40 MPa at
Heat treatment
Tension test at room temperature
working in % % 750 C
Offset yield
Tensile
Temperature Time in Time tB_ to
Elongation A Elongation
strength
strength RM
in C minutes break in hours in
% A 5 in %
Rp0.2
in MPa
I 0.0 1075 10 0.0 10.8 50.4
338 479 28
I 2.8 1075 10 2.8 86 23.3
506 545 17
n
I 2.8 1075 10 2.8 53 19.8
506 545 17 0
I.)
I 5.0 1075 10 5.0 75 18.9
548 571 15
-.1
LO
I 5.0 1075 10 5.0 110 49.5
548 571 15 0
co
_
R 10.0 1075 10 10.0 137 31.7
650 555 8 K)
0
H
R 20.0 1075 10 20.0 108 23.9
739 750 5 I.)
1
0
I.)
R 20.0 1075 10 20.0 73 60.6
739 750 5 1
I.)
u.)
R 40.0 1075 10 40.0 61 24.5
831 837 3
R 40.0 1075 10 40.0 49 34
831 837 3
n. m. = not measured
