Note : Les descriptions sont présentées dans la langue officielle dans laquelle elles ont été soumises.
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LOW STRAIN HIGH DUCTILITY ALLOY
The present invention relates to alloys used to prepare a low strain steel
pipes i.e.
tubes for use in chemical engineering applications. In particular, the
invention relates
to low carbon high strength steel alloys and pipes made from such alloys which
have
low strain but high ductility at elevated temperatures. Such pipes are
typically used in
chemical plants for transporting reactants and products, as the inlet
manifolds in plants
for producing hydrogen and methanol. They may be used in plant such as steam
reformer furnaces which include inlet manifolds that need to possess good
creep
resistance (low strain) as well as having good thermo-mechanical fatigue
resistance
(high ductility).
Steam reforming is the most widespread process for the generation of hydrogen-
rich
synthesis gas from light carbohydrates. The feed material is natural gas,
mostly in the
form of methane, which is ultimately converted into methanol and hydrogen
using
water. This is an endothermic reaction of the gas with water in the form of
steam which
takes place at high temperatures in catalytic tube reactors. The natural gas
feed is
mixed with superheated steam with the appropriate ratio of steam/carbon to
allow
efficient conduct of the reforming process. The mixture then is distributed
via manifold
in vertical rows of catalyst-filled reformer tubes. The mixture flows from top
to bottom of
the tubes and is heated from the outside through the catalyst and reacts
endothermically to produce hydrogen and carbon monoxide which are collected by
outlet manifold.
It is necessary to heat the exposed the tubes to very high temperatures (above
900 C)
to allow the endothermic reaction to take place continuously. This places
stringent
design requirements on the reactor. In addition, the reaction generally takes
place at
elevated pressures and since the reaction occurs at relatively high pressures
(above
20kg/cm2 and up to 40kg/cm2) in the tube, creep damage of the tube is the
usual
parameter limits the working lifetime of the tube.
The increase in the availability of shale gas has driven the development of
the use of
reforming reactors, particularly in the USA and as China. In addition, there
is an
increasing demand for the production of methanol from a combination of sub-
stoichiometric combustion and catalytic steam reforming. It is recognised that
the
annual production of methanol exceeds 40 million tons and continues to grow by
4%
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per year. Methanol has traditionally been used as feed for production of a
range of
chemicals including acetic acid and formaldehyde. In recent years methanol has
also
been used for production of dimethylether, and olefins by so called methanol-
to olefins
process or as blend-stock for motor fuel. Consequently, the limited lifespan
of
conventional tubes used in reforming reactors represents a problem for the
industry.
Pipes for Manifolds can be prepared by a centrifugal casting process.
Centrifugal
casting is a well-established process that is used to cast thin-walled
cylinders, pipes
and other axially symmetric objects. One benefit of this process is that it
allows precise
control of the metallurgy and crystal structure of the alloy product. It is
generally used
for casting iron, steel, stainless steels and alloys of aluminium, copper and
nickel. The
centrifugal casting process employs a permanent mould which is rotated about
its axis
at high speeds of typically 300 to 3000 rpm as the molten metal is poured. The
molten
metal is centrifugally thrown towards the inside mould wall where it is able
to solidify
after cooling. The resulting cast cylinder i.e. tube, has a fine grain and the
surface
roughness of the outer surface of the cylinder is relatively low-Wrought pipes
very
ductile possess in other hand limited creep resistance properties because of
their
relatively low carbon content to allow manufacture. Centrifugally cast pipe
process
allowing casting of alloy with higher carbon content possess very low strain/
higher
creep resistance but lower ductility. The invention proposes the centrifugal
casting of a
pipe with low strain/high creep resistance and high ductility/high thermo-
mechanical
fatigue resistance.
JP64-031931 describes the production of a curved tube made of heat-resistant
alloy.
The tube is prepared by centrifugal casting and the alloy of JP64-031931 is
made from
high strength and heat-resistant cast steel containing 15 to 30% chromium, 20
to 40%
nickel as well as the optional inclusion of smaller quantities of manganese
and
molybdenum. Small quantities of niobium and titanium are also added to the
alloy.
The cast tube is then subjected to the further step of an aging treatment at a
temperature of from 700 to 1100 C to deposit secondary carbide within the
grain
structure. This patent does not attempt to control the primary carbide
formation or to
control the relative amounts of niobium and titanium, or carbon and nitrogen.
Subsequently it is subjected to another processing step involving high
frequency
bending or die-bending at a temperature in the range of 550 to 1100 C.
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W02012/121389 discloses an alloy intended for use in nuclear applications such
as in
heat exchanges in pressurised water reactors. The material is said to have
excellent
thin workability and corrosion resistance. This material is based on a nickel-
chromium-
iron alloy and contains small amounts of manganese, titanium, and optionally
aluminium as alloying elements.
EP1679387 discloses a heat-resistant cast steel which has good high-
temperature
strength, aged strain and creep rupture strength for use as a material in
steam
reforming reaction tubes in fuel cell hydrogen generation systems. The cast
steel
contains chromium and nickel, together with manganese, niobium, titanium and
cerium
as alloying elements.
In addition to all of the usual technical issues associated with preparing a
steel pipe for
use in chemical plant, there are two particular problems which need to be
addressed
when fabricating pipes for this type of application. These issues arise
because of the
harsh working environment that the steel tubes will be exposed to and the fact
that any
'downtime' in plant operation is very costly in terms of lost production. The
pipes need
to be both strong enough to withstand the condition and be resistant to creep
i.e.
deformation over time when exposed to elevated temperature and also its needs
to be
sufficiently ductile to endure to thermo-mechanical fatigue i.e. repeated
variation of
temperatures and/or mechanical stresses and strains.
The present invention aims to provide pipes which are creep-resistant i.e.
pipes which
have a low strain compared to other steel alloys at elevated temperatures. In
addition
the invention aims for the same pipes to be ductile enough to resist damage
from
thermo-mechanical fatigue.
It is also an aim of the present invention to prepare a pipe which can be
produced in a
process which is convenient to run, so that the manufacturing process is
relatively
straight forward. It is also an aim to provide a process which is applicable
to the large
scale production of steel alloy pipes. The invention aims to provide a more
economic
production method and/or which is also more economic when the whole of life
use and
maintenance interruptions are considered. It is also an aim of the present
invention to
provide a steel alloy which is economical to manufacture and which avoids or
reduces
the need for expensive alloying components.
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It is also an aim to have pipes which can be prepared without the need for
further
subsequent processing steps.
It is a further aim to provide pipes which have low internal and external
surface
roughness.
A further aim is to produce pipes which have high strength and / or are high
in
toughness. Another aim is that the pipes should have a good "shelf-life". Long
term
exposure under high temperature conditions can be quite detrimental to
conventional
steel alloys used in such applications. A further aim is to produce steel
alloys which
have good high-temperature strength over an extended period of time. A further
aim is
to provide steel alloys that have good corrosion resistance, particularly at
elevated
temperatures such as those found in a chemical plant. Another aim is to
produce pipes
in which the corrosion resistance is maintained over an extended period of
time in use.
The invention satisfies some or all of the above aims.
According to a first aspect of the present invention, there is provided a
steel
comprising:
from 20.0 to 40.0 atomic % nickel,
from 20.0 to 40.0 atomic % chromium,
from 1.0 to 3.0 atomic % silicon,
from 0.2 to 1.0 atomic % carbon,
from 0.01 to 1.0 atomic % nitrogen
from 0.01 to 0.90 atomic % niobium,
from 0.5 to 3.0 atomic % manganese, and
from 0.01 to 0.90 atomic % of one or more of: titanium, hafnium, zirconium,
vanadium,
tungsten, and molybdenum,
wherein:
(a) the total amount of niobium and one or more of a second carbide forming
element
selected from: titanium, hafnium, zirconium, vanadium, tungsten and molybdenum
is
from 0.50 to 0.91 atomic %, preferably 0.60 to 0.75; most preferably about
0.65
(b) the total amount of carbon plus nitrogen is in the range of 0.2 to 1.2
atomic %,
preferably in the range 0.4 to 1.0 atomic %;
(c) the amount (nitrogen / carbon) is in the range 0.60 to 0.90, preferably
0.80, and
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(d) the amount [nitrogen / (the second carbide forming element(s) plus
niobium)] is in
the range 0.2 to 1.1, preferably 0.4 to 1.0; most preferably about 0.5.
with the balance of the composition being iron and incidental impurities.
In certain cases, the carbide-forming function may be achieved by niobium
alone. In
5 this case, more niobium is required in such alloys than in the case when
the second
carbide-forming elements are present in addition to niobium. The second
carbide-
forming elements need not be present in this group of alloys i.e. the lower
limit of one
or more of: titanium, hafnium, zirconium, vanadium, tungsten, and molybdenum
may
effectively be 0 atomic % in the sense they are not required to perform any
functional
role in such alloys because the effect is achieved by having a higher content
of niobium
alone.
According to a second aspect of the present invention, there is provided a
steel
comprising:
from 20.0 to 40.0 atomic % nickel,
from 20.0 to 40.0 atomic % chromium,
from 1.0 to 3.0 atomic % silicon,
from 0.2 to 1.0 atomic % carbon,
from 0.01 to 1.0 atomic % nitrogen
from 0.50 to 0.91 atomic % niobium, and
from 0.5 to 3.0 atomic % manganese
wherein:
(a) the total amount of carbon plus nitrogen is in the range of 0.2 to 1.2
atomic %,
preferably in the range 0.4 to 1.0 atomic %
(b) the amount (nitrogen / carbon) is in the range 0.60 to 0.90, preferably to
0.80, and
(c) the amount (nitrogen / niobium) is in the range 0.2 to 1.1, preferably 0.4
to 1.0; most
preferably about 0.5.
with the balance of the composition being iron and incidental impurities.
The requirements discussed below for the alloys of the invention, unless
explicitly
stated to the contrary, apply independently to the alloys of both the first
and second
aspects of the invention.
The alloys of the first and second aspects may be used to fabricate steel
pipes.
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The alloy compositions of both aspects of the present invention have a reduced
propensity to suffer from stress fractures. The occurrence of stress fractures
and the
strength and strain of an alloy composition are generally dictated by the
occurrence of
dislocations and their distribution throughout the bulk material. Good high-
temperature
strength and creep-resistance are properties which are mainly due to the
precipitation
strengthening of the grain interiors by alloy carbides. Precipitation
strengthening and
ductility are governed by the precipitate size, shape, distribution and
crystallographic
orientation within the surrounding matrix. The steel alloys of the present
invention have
excellent mechanical properties, show enhanced creep resistance as well as
improved
strength and appropriate ductility to avoid Thermo-Mechanical Fatigue damages.
The
quantities of the various elements in the present invention are quoted in
atomic % since
one of the important features of the invention is the relative quantities of
the various
components. This ensures the correct carbide formation.
Metal carbides that normally provide the strengthening effect in steels are
derived from
niobium, vanadium, molybdenum and tungsten. Hafnium, zirconium and titanium
are
also known carbide formers. All of these elements can be classified as carbide-
forming
elements.
The principal carbide forming component in the alloys of the invention is
usually
niobium and the remainder of the abovementioned carbide-forming elements may
be
used (alone or in a combination of one or more of them) as a second carbide-
forming
component.
One important feature that relates only to the first aspect of the invention
resides in the
careful control of the total amount of the niobium and the one or more second
carbide-
forming elements. This total of the carbide-forming elements mentioned above
is
deliberately controlled to a maximum of from 0.5 to 0.91, preferably from 0.60
to 0.75
atomic weight %, most preferably 0.65 atomic weight %. Thus, the total amount
of
niobium together with one or more of titanium, hafnium, zirconium, vanadium,
tungsten
and molybdenum in the alloys of the first aspect of the invention is never
greater than
0.91 atomic weight %. Usually, the niobium will be the principal part (in
terms of the
number of elemental atoms) of this total. Thus niobium will account for 50
atomic % of
this total of the carbide-forming elements and is more usually at least 80
atomic %, and
may be at least 90 atomic % or even at least 95 atomic % of the total of the
carbide
forming elements. In some circumstances, however, the niobium may be present
in an
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amount of less than 50 atomic % of the total of the carbide forming elements,
with one
or more of the other carbide forming elements being present is a relatively
high
amount.
In the case of only the second aspect of the invention, niobium must account
for a
minimum of 0.50 atomic % of the alloy and is more usually at least 0.60 or
0.70 atomic
%, or even at least 0.80 atomic % of the alloy. The maximum amount of the
niobium is
0.91 atomic %.
The steel alloys of both aspects of the invention consequently have a
relatively small
and dispersed carbide formation compared with known steels for similar
applications.
It is this feature, arising from a careful control of the metallurgical
composition, which
gives contributes to the improved mechanical properties of these steel alloys.
A further
benefit of the steel pipes of the invention (i.e. pipes made using steel
alloys of the
invention) is that they require no subsequent treating to enhance the
mechanical
properties.
According to a third aspect of the invention, there is provided a steel pipe
made from a
steel alloy according to the first or second aspect of the invention.
Classical precipitation strengthening of alloys due to carbide formation
varies as a
function of time at a given temperature. Initially, clusters of solute atoms
form and
then, eventually, precipitate forms which is largely coherent with the matrix.
The
precipitates strengthen the matrix because they prevent dislocation movement
which in
turn inhibits plastic deformation. The steel alloy composition of the
invention is
designed to control primary carbide formation and ensure a substantially
homogeneous
distribution of carbide throughout the matrix. It is also designed to ensure
smaller,
more regular, carbide growth.
Each of the elemental components described in the above steel compositions
plays an
important role in the creep resistant steel and ductility properties of the
alloys described
in the first and second aspects of the present invention. The combination of
elements
gives rise to the very low strain i.e. high creep resistance and maintains the
ductile high
enough i.e. Thermo-Mechanical Fatigue resistance that is observed in the case
of the
alloys of both aspects of the present invention. Furthermore, the combination
of
elements also contributes to the high-temperature strength of the steel tube.
This
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combination of high strength, high creep resistance and high ductility is
manifested
over an extended period of time relative to conventional alloys. The comments
below
in relation to each element apply to all aspects of the invention.
Carbon is an important component of the steel for providing tensile strength
and
resistance to creep rupture. Carbon is an essential component in the carbide
formation
which provides the steel of the present invention with its unique properties.
The carbon
improves the strengthening of the alloy by precipitation of the primary and
secondary
carbides as follows: chromium based carbides (M7C3) and niobium carbides
during
solidification (primary carbides), and chromium based carbides (M23C6) and
niobium
carbides, niobium carbido-nitrides, niobium nitrides during ageing (secondary
carbides). However, too high a quantity of carbon can result in grain boundary
corrosion resistance due to excessive carbide formation and can also result in
reduced
strength to excessive carbide formation. Consequently, carbon must be present
in an
amount in the range of from 0.2 to 1.0 atomic %. Preferably, it is in the
range of from
0.3 to 0.9 atomic %, and more preferably it is present in an amount from 0.40
to 0.70
atomic %.
Not only is it important to control the absolute amount of carbon in the alloy
composition but it is also important to control the amount of carbon relative
to the
amount of nitrogen. The total of (carbon + nitrogen) needs to be below a
maximum
level to allow the precipitation of minimum quantity of fine primary chromium
carbides
and fine secondary carbides/carbo-nitrides but to avoid an excess
precipitation of
primary carbides which will decrease the steel ductility. Hence, the total
amount of
carbon plus nitrogen is a maximum of 1.2 atomic%. Preferably, the total amount
of
(carbon + nitrogen) is in the range 0.4 to 1.0 atomic %. Similarly, the lower
limit of
(carbon + nitrogen) is 0.2 atomic %.
Nitrogen is required because it forms austenite together with carbon and it
contributes
to high-temperature strength. Nitrogen allows the dilution, dispersion, and
the
homogenisation of the carbon. The control of the amount of nitrogen is
important
because it slows the precipitation of primary chromium carbides when it is
added in a
suitable quantity. In effect, the nitrogen helps to control the 'behaviour' of
the carbon
so to control its several precipitations. The nitrogen participates in the
precipitation of
secondary niobium carbides, niobium carbido-nitrides, and niobium nitrides
during
ageing. However, if the quantity of nitrogen is too large then an excessive
amount of
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nitrides are produced which reduces the toughness of the alloy over an
extended
period of time. Both the absolute quantity of nitrogen, and the quantity
relative to
carbon are important to ensure high strength. The nitrogen allows control of
the carbide
precipitations and hence nitrogen must be added in a quantity that is
controlled relative
to that of carbon. Therefore, as more carbon is added so more N is needed;
similarly
as less carbon is added then less nitrogen is needed. The nitrogen disperses
the
carbon in the austenitic matrix which, with fast solidification, slows down
the
precipitation of primary chromium carbides and limits the segregation of the
carbon
close to the primary carbides. Accordingly, in addition, the ratio
(nitrogen/carbon) must
be in the range of from 0.60 to 0.90, preferably 0.70 to 0.80, and more
preferably is
about 0.80. In some embodiments, it is preferable to have more nitrogen
present than
carbon as measured by their contents in atomic %, this helps ensure effective
dispersion of the carbides.
Consequently, nitrogen must be present in an amount in the range of from 0.01
to 1.0
atomic %. Preferably, it is in the range of from 0.20 to 0.70 atomic %, and
more
preferably it is present in an amount from 0.30 to 0.60 atomic %.
Silicon provides the function of a deoxidiser and is usually an essential
component in
an austenite stainless steel. Silicon may also contribute to increasing the
stability of
any surface oxide film. On the other hand, if the content of silicon is too
high the
workability of the steel is reduced. A high Si content can also cause the
formation of a
detrimental phase known as the G phase which is composed of nickel, silicon
and
niobium (Ni16Nb6Si7). Consequently, silicon must be present in an amount in
the
range of from 1.0 to 3.0 atomic%. More often, the amount of silicon is in the
range 1.0
to 2.5 atomic %, or more usually 1.0 to 2.0 atomic %. Preferably, it is in the
range of
from1.45 to 1.75 atomic%, and more preferably it is present in an amount from
1.65 to
1.75 atomic %.
Nickel is an element which is essential in order to obtain a stable austenite
structure
and improves the stability of austenite and supresses the generation of the
sigma
phase. Nickel is the austenitic stabiliser element, allowing the alloy to be
generally
strong at above 800C. Therefore it forms a stable matrix with the iron which
allows the
possible precipitation of the carbides/nitrides. The lower limit of the nickel
content is
chosen simply for the reason that this is a sufficient amount for improving
the stability
of austenite with respect to the lower limits of the other elements. The lower
limit is
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20.0 atomic %. The upper limit is chosen on the grounds of economy and also
with
respect to the upper limits of the other alloy components. Furthermore, nickel
when
present in conjunction with chromium forms a stabilised austenitic structure
which
imparts additional strength and resistance to oxidation at elevated
temperatures.
5 There are diminishing returns as the content of nickel rises hence the
practical upper
limit is around 40.0 atomic%. Preferably, nickel is present is in the range of
from 25.0
to 35.0 atomic %, and more preferably it is present in an amount from 30.0 to
33.0
atomic %.
10 Chromium provides a well-documented and effective corrosion resistance
and
oxidation resistance effect. Chromium also acts as a carbide-former, ensuring
the
creep strengthening precipitations in the alloy. Chromium-based carbide of
general
formula M7C3 is formed during solidification (primary carbide formation) and
chromium-based carbide of general formula M23C6 is formed during ageing
(secondary carbide formation).The lower limit of 20.0 atomic weight % of
chromium is
required in order to ensure sufficient oxidation resistance and the upper
limit of 40.0
atomic weight% is determined by the fact that above this level it is difficult
to obtain a
stable austenite phase. In addition, a high level of chromium renders the
steel
unworkable. Preferably, chromium is present in the range of from 20.0 to 30.0
atomic
%, and more preferably it is present in an amount from 22.5 to 27.5 atomic %.
The principal function of niobium in the alloy is to act as a carbide forming
element.
Niobium allows formation of stable carbides, and even more stable carbo-
nitride and
nitrides in the alloy. Niobium carbides form during solidification which also
may contain
some nitrogen (primary carbides), and niobium carbides, niobium carbido-
nitrides, and
niobium nitrides form during ageing (secondary carbides). Similarly, the
presence of
titanium, or one or more of the second carbide-forming elements, is to form
carbides.
Although titanium is mainly intended for the formation of carbides it is also
engaged in
the formation of nitrides and carbo-nitrides to some degree.
The addition of niobium needs to be carefully controlled in order to ensure
sufficient,
but not too much, carbide formation. The niobium carbide which is formed gives
an
enhanced creep rupture strength and also contributes to maintenance of the
properties
of the high strength and high creep resistance steel alloy over an extended
period of
time. Consequently, niobium must be present in an amount in the range of from
0.01
to 0.90 atomic % when other carbide forming elements are present (i.e. the
first aspect
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of the invention). Usually the lower limit of niobium is 0.30 atomic % though
it may be
as high as 0.40 atomic %. In some embodiments, the lower limit of niobium is
0.50
atomic %. Preferably, the amount of niobium is in the range of from 0.60 to
0.80
atomic %, and more preferably it is present in an amount from 0.65 to 0.75
atomic %.
In certain other cases, the desired range may be 0.40 to 0.60 atomic %. In the
case of
the second aspect, the range of niobium is 0.50 atomic % to 0.91 atomic %, and
more
preferably is 0.60 to 0.91 atomic %, or even more preferably 0.70 to 0.91
atomic %
The ratio N/(Nb+ second carbide forming element) is also important. The
quantity
needs to be such that it allows the beneficial precipitation of very small
secondary
niobium nitrides (MN) (less than 50nm) during ageing of the alloy. Hence the
amount
the amount [nitrogen / (the second carbide forming element(s) plus niobium)]
is in the
range 0.20 to 1.10, preferably in the range 0.40 to 1Ø
The alloy of the first aspect of the invention requires at least one further
carbide-
forming element (the second carbide-forming element) to be present on order to
achieve the desired technical effect. In addition to controlling the upper
limit in order to
avoid excessive carbide formation, the present of excess niobium may also
reduce
corrosion resistance and / or oxidation resistance. The second carbide-forming
element serves the purpose of providing the required degree of carbide
formation
whilst avoiding the problem of a possible reduction in corrosion and/or
oxidation
resistance. These second carbide-forming elements include: titanium, hafnium,
zirconium, vanadium, tungsten and molybdenum. The total amount of carbide
forming
elements must be carefully controlled because if the carbide content is too
high the
mechanical properties once again deteriorate. This is due to a continuous
carbide
network being formed at higher concentrations which weakens the matrix. Hence
the
total amount of niobium and one or more of: titanium, hafnium, zirconium,
vanadium,
tungsten and molybdenum is from 0.50 to 0.91 atomic weight %, preferably 0.60
to
0.91 atomic weight %, and is more preferably from 0.65 to 0.80 atomic %, and
most
preferably is from 0.70 to 0.80 atomic %. The or each second carbide-forming
element
in the alloys of the first aspect of the invention may be present in an amount
of 0.01 to
0.40 atomic %, subject to the total content of the or each second carbide
forming
element and the niobium not exceeding 0.91 atomic %.
Titanium is added to the alloy as a deoxidiser. Furthermore, titanium as a
carbide
forming element not only forms titanium carbides but is also able to form a
titanium-
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niobium double carbide precipitate which improves creep strength. The addition
of too
high an amount of titanium can lead to undesirable oxide formation thereby
reducing
strength. Consequently, titanium when present must be present in an amount in
the
range of from 0.01 to 0.90 atomic %. Preferably, it is in the range of from
0.01 to 0.20
atomic %, and more preferably it is present in an amount from 0.01 to 0.10
atomic %.
Similar restrictions apply to the other carbide forming elements: hafnium,
zirconium,
vanadium, tungsten and molybdenum which taken individually and independently
when
present must be present in an amount in the range of from 0.01 to 0.90 atomic
%.
Preferably, any of those elements when present is present in an amount in the
range of
from 0.01 to 0.20 atomic %, and more preferably is present in an amount from
0.01 to
0.10 atomic %. Only titanium need be present as the second carbide forming
element.
Thus in one embodiment the alloy contains only nickel, chromium, silicon,
carbon,
nitrogen, niobium, manganese, and titanium with the balance being iron and
incidental
impurities. Equally, any one of those other elements may be present as the
sole
second carbide-forming element and the composition would then contain only
nickel,
chromium, silicon, carbon, nitrogen, niobium, manganese, and one of the other
second
carbide-forming elements described above with the balance being iron and
incidental
impurities.
The double carbide formation of niobium and titanium is the reason for the
careful
control of the total amount of niobium and titanium and / or one of the other
carbide-
forming elements hafnium, zirconium, vanadium, tungsten and molybdenum. Each
of
those other carbide-forming elements is able to function in a similar way to
titanium in
forming carbides which contribute to enhanced creep rupture strength. Similar
considerations apply, in terms of the need to avoid excess carbide formation,
when
using these elements hence the requirement that the upper limit of these
elements is
controlled to 0.90 atomic weight % either when present alone as a sole
component
(other than niobium) or when present in combination with one another.
Manganese is a required component of the steels of the present invention
because it
can improve the workability of the alloy. It is also an effective de-oxidant
and
contributes to austenite formation in the steel. The addition of too much
manganese
can result in a reduction in high-temperature strength and also toughness over
an
extended period of time. Consequently, manganese must be present in an amount
in
the range of from 0.5 to 3.0 atomic %. Preferably, it is in the range of from
1.0 to 2.0
atomic %.
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Alloys according to the present invention are produced in a conventional
furnace and
without the need for a special atmosphere. The first stage of preparing the
alloy
involves working out the relative proportions by weight of the various
component
minerals (which are the source of the various elements required in the final
alloy) in
order to achieve the desired amounts of the various elements which are
required in the
final alloy. The solid minerals are added to the hot furnace. Heating is
continued in
order to melt all of the mineral components together and ensure a thorough
mixing of
the minerals in the furnace so that the elements are properly distributed
within the
matrix.
Once melting and mixing has been achieved, any slag is decanted from the
furnace in
order to remove impurities and clean the bath of liquid alloy in the furnace.
A sample of
the molten alloy is then removed from the furnace, allowed to cool and
analysed by x-
ray fluorescence in order to determine its elemental composition. An
adjustment to the
composition may or may not be required at this stage to accommodate for any
elemental mass loss due to volatility. The composition is adjusted by the
addition of
further minerals as necessary, and optionally re-analysed to ensure that the
desired
composition has been achieved.
After the desired composition has been achieved, the temperature is further
raised
above the melting temperature to a tapping temperature in order to ensure easy
pouring of the melt. At the same time, the mould is prepared for centrifugal
casting.
The mould is a conventional centrifugal casting mould and this type of mould
is well
known to the skilled person. The process of preparing the mould involves
washing the
mould with water/steam to clean it and to remove any old mould wash or coating
that
might have been used in a previous casting process. The washed mould is then
coated
with an insulating/release agent which is required to prevent the alloy from
sticking to
the mould after casting. A typical insulating/release agent is silica.
A disc of ceramic is then added to the centrifugal casting mould in the manner
known in
the art in order to ensure that the mould is liquid tight and ready for
casting. This
prevents any alloy leakage during the casting process. The mould temperature
is
adjusted in preparation for the casting and may be in the range of 200 to 300
C. The
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mould is then rotated at high speed to obtain usually the range of 80g to
120g, with a
rotation providing100g being typical for a centrifugal casting speed.
A ladle is then brought to the furnace and a desired weight of alloy is tapped
off for the
purposes of casting. The ladle itself is preheated to a temperature in the
region of 800
to 1000 C in order to minimise cooling of the alloy after pouring. Alloy is
then
transferred to the hot ladle. At this stage, a further analysis of the alloy
may be
performed and any micro-addition of elemental components may also, optionally,
be
performed in order to adjust the final chemistry of the alloy if this is
necessary.
The molten alloy in the ladle is then transferred to a pouring cup. The nose
of the
pouring cup has previously been adjusted to ensure that it mates with and
properly fits
the size of the input tube for the centrifugal casting mould. The level of
molten alloy in
the pouring cup is maintained in order to maintain adequate flow of alloy into
the mould
which is in effect fed by gravity. This provides a continuous flow of alloy
into the mould
until all of the weight of the alloy has been poured into the mould. The mould
is rotated
at high speed i.e. maintained at the centrifugal casting speed during the
process and
whilst the alloy is molten. The length of time the casting process takes
depends
ultimately on the desired thickness of the tube required and the skilled
person is able to
determine a suitable rotation time for a particular thickness of tube and
weight of alloy.
The mould is gradually slowed down as the alloy cools from its solidification
point.
Generally speaking, a "fast" solidification process is one in which the alloy
is cast and
then cools at a rate of more than about 100 C per minute and a "slow"
solidification
process is one in which the alloy is cast and then cools at a rate of about 50
C or
greater per minute. The casting process is usually completed in less than
about 10
minutes. The tube is extracted after the mould stops and the process may be
repeated
again.
The Larson-Miller relation, also widely known for its Larson-Miller Parameter
is a
parametric relation used to extrapolate experimental data on creep and rupture
life of
engineering materials. Larson and Miller (Larson, Frank R. and Miller, James:
A Time-
Temperature Relationship for Rupture and Creep Stresses. Trans. ASME, vol. 74,
pp.
765-775) proposed that creep rate could adequately be described by an
Arrhenius
type rate equation which correlates the creep process rate with the absolute
temperature. They established also that creep rate is inversely proportional
to time.
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Using the assumption that activation energy for the creep process is
independent of
applied stress, it is possible to relate the difference in rupture life to
differences in
temperature for a given stress. The Larson-Miller model is used for
experimental tests
so that results at certain temperatures and stresses can predict rupture lives
of time
5 spans that would be impractical to reproduce in the laboratory. In our
invention we use
a time span of 100,000 hours.
In an embodiment, the alloy of the present invention has mean and minimum
stress
value as table below. In other words, following the Larson-Miller model in our
predictive
10 test for extended rupture time of 10,000 hours, 50,000 hours and 100,000
hours the
alloy has mean and minimum stress values as listed in the Table below at the
listed
temperature from 750 C to 1100 C.
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STRESS RUPTURE VALUES OF PARALLOY LC+
Based on (20/32/1/0.1 - Cr/Ni/Nb/C) "ASTM A351 CT15C"
METRIC UNITS
CREEP RUPTURE STRESS (N/mm2) FOR LIVES OF:
TEMP 10,000 Hour Life 50,000 Hour Life 100,000 Hour Life
C MEAN Minimum MEAN Minimum MEAN
Minimum
750 82.76 78.57 75.53 71.70 72.28 68.62
775 75.66 71.83 67.86 64.42 64.45 61.18
800 68.19 64.73 60.06 57.01 56.58 53.71
825 60.57 57.50 52.36 49.70 48.91 46.43
850 53.04 50.35 44.96 42.68 41.64 39.53
875 45.79 43.46 38.03 36.10 34.91 33.14
900 38.96 36.98 31.69 30.08 28.82 27.35
925 32.68 31.02 26.01 24.69 23.42 22.24
950 27.02 25.65 21.02 19.96 11.47 10.89
975 22.02 20.90 16.74 15.89 14.78 14.03
1,000 17.69 16.79 13.13 12.47 11.47 10.89
1,025 14.01 13.38 10.15 9.63 8.77 8.32
1,050 10.94 10.38 7.72 7.33 6.60 6.26
1,075 8.42 7.99 5.79 5.50 4.89 4.64
1,100 6.38 6.06 4.28 4.06 3.57 3.39
Notes:
1 Values obtained by Larson-Miller extrapolation
2 Regression
Analysis / Method of Least Squares
3 Minimum Value = MEAN - (1.65 x Standard Deviation)
4 Data study using "Trijay " computer program
Creep strength can be measured in accordance with the standard industrial test
ASTM
E139-1.
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Alloys having the following compositions were produced in accordance with the
invention.
LC+ at% - general requirements
Ni at% Cr at% Nb at% Si at% WTI N/(M+Ti) N/C N+C at%
N at% N at% C at% C at% Fe
30 min 25.5min 0.78 2 0.65 0.5 0.8 1.2
0.30 0.60 0.40
0.70
balance
max Max Optimum Min Min Max Min Max Min Max
LC+ (Trial A)
Fe Ni Cr Si C Nb Mn N Ti C+N N/C Nb+Ti N/(Ti+Nb)
A wt% 42.98 33.71 20.15 0.97 0.131 1.122 0.77 0.121
0.03
at% 42.45 31.67 21.38 1.90 0.602 0.666 0.77 0.477 0.03 1.079 0.79 0.696
0.69
The steel tubes of the present invention show excelled high-temperature
strength and
low strain i.e. high creep resistance. The tubes also display excellent
corrosion
resistance at elevated temperatures over an extended period of time.
Consequently,
these steels are particularly suited to use in chemical plant under demanding
environments such as a reformer. In addition, it is expected that steel tubes
according
to the invention may be used in other applications such as ethylene crackers
and in
nuclear applications in heat exchanges and the like, such as those found in
pressurised
water reactors.
Without wishing to be bound by theory, it is believed that the beneficial
properties of the
steel alloys of the present invention arise due to the improved primary
carbide
precipitation and subsequent secondary carbide formation that occurs due to
the
carefully controlled relationships between the carbon, nitrogen and carbide
forming
elements in the alloys of the present invention. The alloys of the present
invention
benefit from particularly small carbide formation and the carbides formed in
the steels
of the present invention are longer and thinner than those in comparable
nickel
chromium steels.
We consider that careful control of the niobium carbide formation relative to
other
carbides so that relatively a greater proportion of niobium carbide is formed
in the
alloys according to the invention. For example, the standard H39WM alloy
contains
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25% by weight chromium, 35% by weight nickel, 1% by weight niobium and 0.4% by
weight carbon together with micro additions of other alloying elements and
this alloy
has a chromium carbide (present as Cr307) present in an amount of 74%, based
on a
fraction analysis of a photo at 200 times magnification. In the alloys of the
present
invention this is found at levels around 61%.
Similarly, the niobium carbide content in the traditional H39WM alloy is
typically about
26% whereas in the alloys according to the present invention it is around 39%,
based
on a fraction analysis of a photo at 200 times magnification. An important
feature of
the alloys of the present invention is that they have a more homogeneous
carbide
formation. In other words, the carbides that are formed are more similar in
size to one
another and are smaller than in the conventional alloys. Thus, not only do the
alloys of
the present invention contain smaller carbides in otherwise apparently similar
alloy
compositions but also contain a greater proportion of niobium carbide. A lower
limit of
85%, and more preferably 90% of niobium carbide as a proportion of the total
amount
of as-cast carbide present is preferred. Similarly, a maximum proportion of
15%, and
more preferably a maximum of 10% of the total as-cast carbides present is
represented
by chromium carbide. Again, these figures refer to a fraction analysis of a
photo at 200
times magnification. The presence of finer, longer but discontinued niobium
carbides in
the present invention improve the creep resistance of the steel as improvement
of the
control of the growth of secondary carbides which over time reduces the
ability to stop
movement of dislocations. This in turn means that the steel would become
weakened
over time.
The fast precipitation of niobium from the melt during the centrifugal casting
process
allows the alloy compositions of the present invention to be cast with the
homogeneous
carbide formation and relatively larger proportion of niobium carbides to
chromium
carbides as compared with convention steel alloys, 6 to 9.5 time more in the
present
invention based on a fraction analysis of a photo at 200 times magnification..
A further important feature of the alloys of the present invention relates to
the amount
of secondary chromium carbides on the surface. In the conventional alloys, the
surface
fraction of Cr2306 is about 4% whereas in the alloys according to the
invention it is at
least 30% and more preferably 50% based on a fraction analysis of a photo at
200
times magnification.
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The properties of a steel according to the invention having the composition
LC+ (Trial
A) was investigated and the results are shown in the following tables.
Figures 1 to 3 show the properties of the steels H39WM, CR32W (a conventional
steel
alloy) and LC+ (a steel alloy according to the invention). Figure 1 shows the
improvements in MSW thickness, Figure 2 shows the Larson-Miller curves of
H39WM
and LC+, and Figure 3 shows a constant stress creep test at 9500 of CR32W and
LC+.
The superior properties of the steels of the present invention are also
evident in each
case relative to the conventional steel from the following data for LC+.
Room temperature tensile properties (Minimum values) N/mm2)
UTS 450
0.2% PS 250
Elongation 15%
Coefficient of linear expansion mm/mm C (1/K)
20-100 C 14.5 x 106
20-750 C 17.5 x 10-6
20-1000 C 18.5 x 10-6
Density
7.94 Gm/cc (0.288 lb/in3)
Hot tensile properties N/mm2 (Typical value)
800 C 900 C
Uts 238 145
0.2% PS 150 90
Elongation 40% 47%
Thermal conductivity (w/mK)
100 C 13.4
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800 C 25.9
1000 C 30.5
5
Temp ( C) Stress (Mpa) Creep test life (hrs)
800 83.44 544
850 62.77 889
870 58.8 1,243 (sample including a weld)
950 35.9 1,414
950 41.52 416
983 27.58 1,427
1000 30.9 442
1050 20.82 414
1075 18.72 215
1100 16 117
Further steel alloys having the following compositions were produced.
