EP3188859A1

EP3188859A1 - Low strain high ductility alloy

Info

Publication number: EP3188859A1
Application number: EP15757561.4A
Authority: EP
Inventors: Dominique Flahaut
Original assignee: Paralloy Ltd
Current assignee: Paralloy Ltd
Priority date: 2014-09-04
Filing date: 2015-08-28
Publication date: 2017-07-12
Anticipated expiration: 2035-08-28
Also published as: PL3188859T3; WO2016034865A1; PT3188859T; GB201415624D0; DK3188859T3; US20170306463A1; GB2544441A; TR201816177T4; GB201703804D0; EP3188859B1; MY191542A; ES2694052T3; CA2959837A1

Abstract

The present invention relates to alloys used to prepare a low strain steel pipes 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).

Description

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% 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 1 100°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 1 100°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 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. 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 (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 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. 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 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 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 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 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 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 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 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 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 1.0. 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- 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 %. 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 80g to 120g, with a rotation providingl OOg 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. 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 mean and minimum stress value as table below. In other words, following the Larson-Miller model in our predictive 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 1 100°C.

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:

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. Alloys having the following compositions were produced in accordance with the invention.

- general requirements

LC+ (Trial A)

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 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 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 Cr23C6 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. 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 950C 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+.

Density

7.94 Gm/cc (0.288 lb/in³)

Thermal conductivity (w/mK)

100°C 13.4 800°C 25.9

1000°C 30.5

Further steel alloys having the following compositions were produced.

Claims

1. 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

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 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, 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.

2. 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, and

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, and

(c) the amount (nitrogen / 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.

3. A steel as claimed in claim 1 or 2, wherein carbon is present in the range of from 0.4 to 0.7 atomic %.

4. A steel as claimed in claim 1 , 2 or 3, wherein nitrogen is present in an amount in the range of from 0.30 to 0.60 atomic %.

5. A steel as claimed in claim 1 , 2, 3 or 4, wherein the ratio (nitrogen/carbon) is in the range of from 0.70 to 0.80 atomic %.

6. A steel as claimed in any preceding claim, wherein nickel is present in the range of from 25.0 to 35.0 atomic %.

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

8. A steel as claimed in any preceding claim, wherein chromium is present in the range of from 20.0 to 30.0 atomic %.

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

10. A steel as claimed in any preceding claim, wherein manganese is present in an amount in the range of from 0.5 to 3.0 atomic %.

11. A steel as claimed in any preceding claim when not dependent on claim 2, wherein titanium is present in the range of from 0.01 to 0.20 atomic %.

12. A steel pipe made from an alloy according to any of claims 1 to 1 1.

13. A steel pipe as claimed in claim 12, wherein the time to rupture is at least 215 hrs when tested at 1050°C with 30MPa Stress.

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