US20170130301A1

US20170130301A1 - Low ductility alloy

Info

Publication number: US20170130301A1
Application number: US15/317,600
Authority: US
Inventors: Dominique Flahaut
Original assignee: DONCASTERS PARALLOY
Current assignee: DONCASTERS PARALLOY
Priority date: 2014-07-10
Filing date: 2015-06-12
Publication date: 2017-05-11
Also published as: CA2953240A1; CA2953240C; GB201621670D0; GB2542519A; EP3167090A1; GB2542519B; WO2016005724A1; GB2542519A8

Abstract

The present invention relates to a low ductility steel tube for use in chemical engineering applications. In particular, the invention relates to a high strength steel tube which has low ductility at elevated temperatures. Such tubes are typically used in chemical plants for transporting reactants and products. One such application includes the use in plants for producing hydrogen and methanol. The tubes could also be used when producing ethylene and other hydrocarbons.

Description

The present invention relates to a low ductility steel tube for use in chemical engineering applications. In particular, the invention relates to a high strength steel tube which has low ductility at elevated temperatures. Such tubes are typically used in chemical plants for transporting reactants and products. One such application includes the use in plants for producing hydrogen and methanol. The tubes could also be used when producing ethylene and other hydrocarbons.
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 20 kg/cm²and up to 40 kg/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 substoichiometric combustion and catalytic steam reforming. It is recognised that the annual production of methanol exceeds 40 million tons and continues to grow by 4% 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 blendstock for motor fuel. Consequently, the limited lifespan of conventional tubes used in reforming reactors represents a problem for the industry.
Another area in which the tubes could be of potential value is in ethylene production since this process is also conducted at elevated temperatures and pressures. Ethylene is a basic chemical which is widely used in the production of a number of common items such as plastic packaging that are prepared by polymerisation of ethylene and ethylene derivatives. The feedstock for an ethylene plant is usually ethane or other natural gases and gas liquids derived from traditional oil and gas sources. The gas, principally or entirely ethane, is heated in a steam mix in order to break down the ethane into ethylene, hydrogen and other by-products. This process is known as cracking.
Steam cracking is a petro-chemical process which, in general terms, saturated hydrocarbons are broken down into smaller, frequently unsaturated, hydrocarbons and hydrogen. This process represents the principal industrial method for producing lighter alkenes for subsequent use in a variety of chemical processes. Naphtha cracking represents the dominant source of ethylene globally, although gas cracking has recently become more important. Irrespective of the exact nature of the feedstock and of the chemical process used to produce ethylene and other olefins, the cracking process is invariably conducted at elevated temperatures and these are typically in the region of 800 to 1100° C.
Sudden cooling then stops the reaction and the subsequent mixture of gaseous products can be compressed, chilled and then separated in a serious of distillations towers. Product such as ethene and other alkenes are individually recovered at this stage. Any un-cracked ethane or other feedstock material is recycled and the remaining by-product gases are further processed for other uses.
Pipes of this type are generally 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. The internal surface has slightly more impurities and inclusions but may be subject to machining to modify the surface roughness and/or geometry.
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.
WO2012/121389 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 ductility 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. With conventional pipes, there is an issue with pipe creep and pipe sag due to the high temperatures used during the cracking process and these effects become more problematic over time. In the case of reforming reactors, the possible failure of such pipes could have catastrophic effects. Equally, the necessary maintenance schedule for reactors of this type means that the “down-time” can present a significant economic burden.
Another problem resides in the generation of carbon as a by-product of the cracking reaction which can contaminate the product and also build up on the inside of the pipework. This causes constrictions and necessitates more frequent maintenance and consequent plant downtime.
The present invention aims to provide pipes which are creep-resistant i.e. pipes which have a low ductility compared to other steel alloys at elevated temperatures. A further aim is to provide low ductility alloy steels which can be produced in normal atmospheric conditions i.e. in air and to be prepared either in a reduced pressure atmosphere or in an inert atmosphere. This represents a considerable processing and economic advantage.
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.
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.
It is another aim to produce pipes who metallurgical properties mean that they tend to be less prone to carbon accumulation. This is particularly important for use in cracking reactors.
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 the present invention, there is provided a steel pipe 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 0.5 to 2.5 atomic % carbon,
from 0.01 to 1.0 atomic % nitrogen
from 0.01 to 0.90 atomic % niobium,
from 1.0 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;
(b) the total amount of carbon plus nitrogen is in the range of 1.2 to 3.0 atomic %, preferably in the range 1.5 to 2.5 atomic %;
(c) the amount (nitrogen/carbon) is in the range 0.20 to 1.0, and
(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;
with the balance of the composition being iron and incidental impurities.
The alloy compositions of the present invention have a reduced propensity to suffer from stress fractures. The occurrence of stress fractures and the strength and ductility 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 is 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 and show enhanced creep resistance as well as improved strength.
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 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 0.91 atomic weight %. Thus the total amount of niobium together with one or more of titanium, hafnium, zirconium, vanadium, tungsten and molybdenum 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. In some circumstances, however, the niobium may be present in an amount of less than 50 atomic % of the total.
The steel alloy of the invention consequently has 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. A further benefit of the steel pipes of the invention is that they require no subsequent treating.
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 composition plays an important role in the creep resistant steel of the present invention. The combination of elements gives rise to the very low ductility i.e. high creep resistance that is observed in the case of the present invention. Furthermore, the combination of elements also contributes to the high-temperature strength of the steel tube. This high strength and high creep resistance is manifested over an extended period of time relative to conventional alloys.
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 1.2 to 2.5 atomic %. Preferably, it is in the range of from 1.5 to 2.5 atomic %, and more preferably it is present in an amount from 1.75 to 2.25 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 above a minimum level to allow the precipitation of minimum quantity of fine primary chromium carbides and fine secondary carbides/carbo-nitrides and also needs to be below a maximum level to avoid over saturation of the austenitic matrix, and as a consequence, loss of the beneficial effect of fast solidification and loss of the control of precipitation before and during solidification. Hence, the total amount of carbon plus nitrogen is in the range of from 1.2 to 3.0 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 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 (M7C3) 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 0.20 to 1.00, and preferably is in the range 0.20 to 0.50 atomic %.
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 0.50 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 %. Preferably, it is in the range of from 1.45 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 800 C. 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 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. 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 %.
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 (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 primary carbides are of the form M73C. 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 %. Preferably, it 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 %. 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 50 nm) 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.
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. 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 %.
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-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 1.0 to 3.0 atomic %. Preferably, it is in the range of from 1.0 to 2.0 atomic %.
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 mould is then rotated at high speed to obtain usually the range of 80 g to 120 g, with a rotation providing 100 g 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 microaddition 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.
Alloys of the invention can be assessed by the Larson-Miller relation. The Larson-Miller relation, also widely known as the 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.
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 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 a minimum stress value of 40 MPa when measured at 850° C. for 100,000 hours rupture time. In other words, following the Larson Miller model in our predictive test for an extended rupture time of 100,000 hours, the alloy of the present invention has a minimum stress value of 40 MPa at 850° C. In a further embodiment, the minimum stress value is 30 MPa at 900° C. In another embodiment, the minimum stress value is 20 MPa at 950° C. In yet another embodiment, the minimum stress value is 15 MPa at 1000° C. In a still further embodiment, the minimum stress value is 10 MPa 1050° C. In further embodiments, the alloy will exhibit creep properties such that it satisfies to or more of the above minimum stress values taken in any combination. In a particularly preferred embodiment, the alloy of the present invention has creep properties in which the minimum stress is the same as or higher in the range from 900° C. to 1050° C. than the line represented by H39WM in FIG. 3.
Creep strength can be measured in accordance with the standard industrial test ASTM E139-1.
Alloys having the following compositions were produced in accordance with the invention.

H39WM+ at % - general requirements

Ni at %

Cr at %

Nb at %

Si at %

M + Ti

N/(M + Ti)

N/C

N + C at %

N at %

C at %

Fe

30 min	25.5 min	0.78	2	0.7	0.5	0.2	2.25	2.45	0.350	0.550	1.900	2.100	balance
		max	Max	Optimum	Min	Min	Min	Max	Min	Max	Min	Max

H39WM++ at % - general requirements

Ni at %

Cr at %

Nb at. %

Si at %

M + Ti

N/(M + Ti)

N/C

N + C at %

N at %

C at %

Fe

30 min	25.5 min	0.7	1	0.75	0.4	0.22	1.9	2	0.300	0.400	1.600	1.700	balance
32.5	26.0	max	Max	Optimum	Min	Min	Min	Max	Min	Max	Min	Max
max	max

H39WM+ (Trial A)

Fe

Ni

Cr

Si

C

Nb

Mn

N

Ti

C + N

N/C

Nb + Ti

N/(Ti + Nb)

A	wt %	37.18	34.94	24.53	0.88	0.44	1.12	0.78	0.10	0.04
	at %	36.28	32.44	25.72	1.71	2.00	0.66	0.77	0.38	0.04	2.37	0.19	0.70	0.539

H39WM++ (Trial B)

Trial

C

Si

Mn

Ni

Cr

Mo

Nb

W

Ti

Zr

N

Fe

C + N

N/C

Nb + Ti

N/(Ti + Nb)

B	wt %	0.35	0.76	0.72	36.07	24.36	0.01	1.24	0.04	0.04	0.01	0.08	36.32
	at %	1.61	1.48	0.71	33.71	25.69	0.01	0.73	0.01	0.05	0.01	0.31	35.67	1.92	0.20	0.79	0.401

H39WM+ (further production examples)

Tubes

production

Fe

Ni

Cr

Si

C

Nb

Mn

N

Ti

C + N

N/C

Nb + Ti

N/(Ti + Nb)

1	wt %	37.51	34.71	24.19	0.96	0.44	1.22	0.82	0.10	0.05
	at %	36.61	32.23	25.36	1.86	2.00	0.72	0.81	0.40	0.05	2.39	0.20	077	0.517
2	wt %	37.59	34.95	24.03	0.88	0.43	1.22	0.76	0.11	0.03
	at %	36.68	32.45	25.19	1.71	1.95	0.72	0.75	0.44	0.04	2.40	0.23	0.75	0.591
3	wt %	37.61	34.85	24.08	0.88	0.43	1.23	0.77	0.11	0.03
	at %	36.71	32.36	25.25	1.71	1.95	0.72	0.76	0.44	0.04	2.39	0.22	0.76	0.573
4	wt %	37.66	34.79	24.02	0.97	0.43	1.19	0.81	0.10	0.04
	at %	36.75	32.30	25.18	1.88	1.95	0.70	0.80	0.38	0.04	2.33	0.20	0.74	0.516
5	wt %	37.46	35.15	23.99	0.87	0.43	1.19	0.76	0.11	0.04
	at %	36.56	32.64	25.15	1.69	1.95	0.70	0.75	0.42	0.05	2.37	0.22	0.74	0.565

The steel tubes of the present invention show excelled high-temperature strength and low ductility 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 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 Cr₃C₇) 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 30%, and more preferably 35% of niobium carbide as a proportion of the total amount of carbide present is preferred. Similarly, a maximum proportion of 70%, and more preferably a maximum of 65% of the total 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 smaller and more dispersed carbides in the present invention improve the creep resistance of the steel because the growth of secondary carbides 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.
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 Cr₃C₇is about 4% whereas in the alloys according to the invention it is at least 6% and more preferably 8% based on a fraction analysis of a photo at 200 times magnification.
The properties of a steel according to the invention having the composition H39WM+ (Trial A) and H39WM++ (Trial B) were investigated and the results are shown in the following tables.
FIGS. 1 to 4 show the properties of the steels H39WM (a conventional steel alloy) and H39WM+ (a steel alloy according to the invention). FIG. 1 shows the carburisation properties over a carburisation test cycle of 100 hours, FIG. 2 shows the improvements in MSW thickness, FIG. 3 shows the creep properties and rupture strength after 1000 hours at a variety of temperatures, and FIG. 4 shows a constant stress creep test. 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 H39WM+.


Room temperature tensile properties (Minimum values) N/mm²)

	UTS	460 (66.7 ksi)
	0.2% PS	225 (32.6 ksi)
	Elongation	10%


Coefficient of linear expansion mm/mm° C. (1/K)

	20-100° C.	14.4 × 10⁻⁶
	20-800° C.	16.8 × 10⁻⁶
	20-1000° C.	17.6 × 10⁻⁶
	20-1100° C.	18.1 × 10⁻⁶


Density

	7.97 Gm/cc (0.288 lb/in³)

Hot tensile properties N/mm²(Typical value)

	800° C.	900° C.	1000° C.	1100° C.

Uts	250 (36.3 ksi)	160 (23.2 ksi)	98 (14.2 ksi)	74 (10.7 ksi)
0.2% PS	162 (23.5 ksi)	107 (15.5 ksi)	72 (10.4 ksi)	59 (8.6 ksi)
Elongation	26%	28.5%	32.5%	21%


Thermal conductivity (w/mK)

100° C.	13.0 (0.031 cal/cm sec ° C.)
800° C.	24.3 (0.058 cal/cm sec ° C.)
1000° C.	27.7 (0.066 cal/cm sec ° C.)
1100° C.	29.7 (0.071 cal/cm sec ° C.)

FIG. 5 shows the constant stress creep test at 950° C.
The following tables show the creep rupture test results and the composition of the tube is shown below (H39WM+(Trial A)). The carbon content was greater than 1 at %.


Fe	Ni	Cr	Si	C	Nb	Mn	N	Ti

A	wt %	37.18	34.94	24.53	0.88	0.44	1.12	0.78	0.10	0.04
	at %	36.28	32.44	25.72	1.71	2.00	0.66	0.77	0.38	0.04


	Temp (° C.)	Stress (Mpa)

	1050	30
	1075	30
	1100	30


Creep test life (hrs)

240 & 289
108 & 130
33 & 34


	Temperature	Minimum Strain rate
	(° C.)	(%/hr)

	1050	1.10E−03 & 1.10E−03
	1075	1.44E−03 & 1.04E−03
	1100	1.88E−04 & 4.14E−04

Further steel alloys having the following compositions were produced.


	Elements
	(wt %)

TUBE	Fe	Ni	Cr	Si	C	Nb	Mn	N

1	37.9	34.9	24.0	0.88	0.43	1.16	0.76	0.11
2	37.7	35.0	24.0	0.88	0.43	1.22	0.76	0.11
3	37.1	35.2	24.5	0.85	0.42	1.10	0.78	0.10
4	37.8	34.9	24.1	0.88	0.43	1.23	0.77	0.11
5	37.7	34.7	24.2	0.96	0.43	1.22	0.82	0.10
6	37.8	34.8	24.0	0.97	0.43	1.19	0.81	0.10

The creep test results for these production examples are shown in the table below.


		986 C./42.3 Mpa
		Creep test life
	Tubes	(hrs)

	1	217
	2	220
	3	232
	4	242
	5	264
	6	267

The tables below show a comparison between the creep test results for the conventional alloy H39WM, and alloys of the invention H39WM+ and H39WM++. All of the alloys were produced under identical casting conditions which were as follows:


Mould	Mould wash (Silica based)	1.0-1.3 mm
	thickness
	Mould temperature	240-250 C. (stick)
Tapping	Temperature	1710 C.-1720 C.
Casting	Temperature	1615 C.-1620 C.
	weight	225 kg/278 kg
Tube	Outside diameter	127.4 mm
	Minimum Sound Wall	13.5 mm/16.7 mm

Creep Test Results for H39WM, H39WM+ and H39WM++

Life (hrs)

Temp	Stress	H39WM	H39WM+	H39WM++

900	58.88	415	701	717
950	42.17	546	734	863
980	35.42	487	*	802
1000	31.24	494	*	813
1050	22.05	520	*	683
1100	16	230	379	329

* data not available

It is evident that both H39WM+ and H39WM++exhibited superior lifetimes relative to H39WM across the temperature range tested. The temperatures at which the alloys of the invention and the conventional alloy was tested are representative of temperatures which might typically be experienced by the tubes in use. Thus, it can be seen that the tubes produced from alloys according to the invention exhibited extended, lifetimes both at low temperatures and also at high temperatures, relative to tubes made of the known H39MW alloy using the same process. Indeed, across all the temperature ranges where data was available, it can be seen that the lifetimes have been extended by about 50% or more relative to tubes produced from conventional alloys. This represents a significant advantage in engineering applications.

Claims

1. A steel pipe 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 0.5 to 2.5 atomic % carbon,

from 0.01 to 1.0 atomic % nitrogen,

from 0.01 to 0.90 atomic % niobium,

from 1.0 to 3.0 atomic % manganese, and

from 0.01 to 0.90 atomic % of one or more second carbide forming elements selected from: titanium, hafnium, zirconium, vanadium, tungsten, or molybdenum,

wherein:

(a) a total amount of niobium and the one or more second carbide forming elements is from 0.50 to 0.91 atomic %;

(b) a total amount of carbon plus nitrogen is in the range of 1.2 to 3.0 atomic %;

(c) an amount of (nitrogen/carbon) is in the range 0.20 to 1.0, and

(d) an amount of [nitrogen/(the one or more second carbide forming elements plus niobium)] is in the range 0.2 to 1.1;

with the balance of the composition being iron and incidental impurities.

2. A steel pipe as claimed in claim 1, wherein carbon is present in the range of from 1.5 to 2.5 atomic %.

3. A steel pipe as claimed in claim 1, wherein nitrogen is present in an amount in the range of from 0.20 to 0.70 atomic %.

4. A steel pipe as claimed in claim 1, wherein the amount of (nitrogen/carbon) is in the range of from 0.20 to 0.50 atomic %.

5. A steel pipe as claimed in claim 1, wherein nickel is present in the range of from 25.0 to 35.0 atomic %.

6. A steel pipe as claimed in claim 5, wherein nickel is present in an amount from 30.0 to 33.0 atomic %.

7. A steel pipe as claimed in claim 1, wherein chromium is present in the range of from 20.0 to 30.0 atomic %.

8. A steel pipe as claimed in claim 7, wherein chromium is present in an amount from 22.5 to 27.5 atomic %.

9. A steel pipe as claimed in claim 1, wherein manganese is present in an amount in the range of from 1.0 to 2.0 atomic %.

10. A steel pipe as claimed in claim 1, wherein titanium is present in the range of from 0.01 to 0.20 atomic %.

11. A steel pipe as claimed in claim 1, wherein a time to rupture is at least 215 hours when tested at 1050° C. with 30 MPa Stress.

12. A steel pipe as claimed in claim 1, wherein a time to rupture is at least 210 hours when tested at 1000° C. with 40 MPa Stress.

13. A steel pipe as claimed in claim 1, wherein the total amount of niobium and the one or more second carbide forming elements is from 0.60 to 0.75 atomic %.

14. A steel pipe as claimed in claim 1, wherein the total amount of carbon plus nitrogen is in the range of 1.5 to 2.5 atomic %.

15. A steel pipe as claimed in claim 1, wherein the amount of [nitrogen/(the one or more second carbide forming elements plus niobium)] is in the range 0.4 to 1.0.