Title of Invention


Abstract (EN) The invention concerns steel for high temperature use containing by weight: 0.06 to 0.20% of C, 0.10 to 1.00% of Si, 0.10 to 1.00% of Mn, not more than 0.010% of S, 10.00 to 13.00% of Cr, not more than 1.00% of Ni, 1.00 to 1.80% of W, Mo such that (W/2+Mo) is not more than 1.50%, 0.50 to 2.00% of Co, 0.15 to 0.35% of V, 0.040 to 0.150% of Nb, 0.030 to 0.12% of N, 0.0010 to 0.0100% of B and optionally up to 0.0100% of Ca, the rest of the chemical composition consisting of iron and impurities or residues resulting from or required for preparation processes or steel casting. The chemical constituent contents preferably verify a relationship such that the steel after normalizing heat treatment between 1050 and 1080°C and tempering has a tempered martensite structure free or practically free of $g(d) ferrite.
Full Text

The invention relates to steels for use under stress at high temperatures of about 600°C to 650°C, more particularly steels known as ferritic steels with a high chromium content with a tempered martensitic structure both at ambient temperature and at service temperatures.
The invention is applicable to tubular metal products such as superheater tubes, reheater tubes, headers or pipings for superheated or reheated steam for boilers, or tubes for furnaces for chemistry or petrochemistry. Background technique
Such products are usually seamless tubes obtained after a severe hot plastic deformation operation carried out on solid bars of highly specialized steel.
Apart from ferritic steels with 2.25% Cr-1% Mo according to ASTM A213 type T22, austenitic stainless steel tubes according to ASTM A213 (ASTM = American Society for Testing and Materials) type TP321H, TP347H have long been known, containing about 0.05% C, 18% Cr, 11% Ni and stabilized with Ti or Nb respectively.
Such steels are highly resistant to corrosion by steam because of their chromium content and have high creep rupture strengths of up to 700°C due to their austenitic structure.
In contrast, they suffer from major drawbacks due to their austenitic structure, which renders them incompatible with steels with a ferritic or martensitic structure, which are of necessity used in other parts of the boiler that are less exposed to high temperatures; hence, the search for materials with a ferritic or martensitic structure is of great importance.
For high temperature uses, then, tubes of ASTM A213 T91 steel (generadly used for small superheater tubes) or ASTM A335 P91 (generally used for the largest pipes for header or superheated steam piping) are known. These grades contain 0.1% C, 9% Cr, 1% Mo, 0.2% V,

0.08% Nb and 0.05% N and have a creep rupture strength at 105 hrs at 600°C (aR105600°C) of 98
ASTM A213 T92 steel (or ASTM A335 P92 steel) has a chemical composition close to T91/P91 except that the Mo content is greatly reduced and it contains 1.8% W and a tiny amount of boron; the creep rupture strength at 105 hrs at 600°C (aR105h600°C) for that steel is of the order of 120 MPa.
Said steels T91, P91, T92, P92 contain 9% Cr and some of their users believe that such a Cr content is insufficient to resist hot oxidation and/or corrosion by steam beyond 600°C, in particular at 650°C because of the metal temperature envisaged for the tubes of the superheaters in future power stations.
Certainly, the presence of an oxide layer on the inner surface of the tubes of superheaters, which layer derives from corrosion of the steel by the steam moving in the tubes, creates a thermal resistance which increases with the thickness of said layer and, at constant thermal flux, entrains an increase in the mean temperature of the tubes and thus a large reduction in their service life.
Further, flaking of said layer when it is too large may lead to accumulation of debris in the bends in the superheaters, impeding the movement of steam with a supplemental risk of overheating the tubes. Flaking can also result in debris being entrained into the turbine and can thus damage its blades.
German DIN 17175 X20CrMoV12-l (abbreviated to X20) steel is also known, containing 0,20% C, 11% to 12% Cr, 1% Mo and 0.2% V.
That steel is claimed to be more resistant to hot oxidation than T91 or T92 because of its Cr content, but it is far less resistant to creep rupture than T91/P91 and it is difficult to weld, in particular when very thick.

It would thus be of advantage to modify the T92/P92 steel for which the creep strength is satisfactory but for which the hot oxidation resistance is insufficient by increasing its Cr content to 12% Cr, but such an increase would come up against the problem of the appearance of 5 ferrite in the structure, which is deleterious to the transformation of steel (forgeability), for its toughness and for its creep strength.
The increase in the Cr content in X20 steel is compensated for by a higher C content (0.20% as opposed to 0.10%) and by addition of a moderate amount of Ni (between 0.5% and 1%).
A C content of 0.20% or more appears to be not much desirable as regards weldability. Adding a large amount of Ni, though, has the disadvantage of greatly reducing the Acl point and thus limiting the maximum tempering temperature of the tubes; it also appears to be deleterious to the creep rupture strength.
United States patent US-A-5 069 870 discloses the addition of Cu (austenite forming element) in amounts of 0.4% to 3% to a 12% Cr steel to compensate for the increase in Cr content. However, adding Cu causes problems as regards forgeability when fabricating tubes for superheaters by hot rolling.
A grade with 11% Cr, 1.8% W, 1% Cu and micro-alloyed with V, Nb and N with the same disadvantages is defined in ASTM A213 and A335 and termed T122, P122.
Japanese patent application JP-A-4 371 551 discloses adding between 1% and 5% (and generally more than 2%) of Co (also austenite forming) to a steel containing 0.1% C, 8% to 13% Cr, 1% to 4% W, 0.5% to 1.5% Mo, less than 0.20% Si (and in fact less than 0.11% Si) and micro-alloyed with V, Nb, N and B to obtain a creep rupture resistance that is very high and a Charpy V-notch impact test strength that is sufficient after ageing. Such a steel is expensive to produce, however.

The same is true for the steels described in European patents EP-A-0 759 499, EP-A-0 828 010, JP-A-9 184 048 and JP-A-8 333 657, which contain more than 2% Co and preferably at least 3%.
European patent application EP-A-0 892 079 also proposes adding Co in amounts of 0.2% to 5% but in a steel containing less than 10% Cr, which does not solve the problem described above.
Japanese patent application JP-A-11 061 342 and European patent application EP-A-0 867 523 also propose adding Co, but jointly with the addition of Cu for the first document and at least 1% Ni for the second document. However, we described the unacceptable disadvantages of such additions above.
European patent application EP-A-0 758 025 also proposes adding Co, generally in very large amounts; for that reason, to prevent the formation of intermetallic precipitates based on Cr, Mo, Co, W, C and Fe, that document jointly proposes adding (Ti or Zr) and alkaline-earths (Ca, Mg, Ba) or rare earths (Y, Ce, La).
Adding Ti or Zr, however, suffers from the major drawback of forming coarse nitrides with the nitrogen in the steel and preventing the formation of ultrafine carbonitrides of V and Nb responsible for the high creep strength.
JP-A-8 187 592 also proposes adding Co with a particular relationship between the (Mo+W) and (Ni+Co+Cu) contents, but said additions and relationships are proposed for optimizing the composition of the added materials for welding, and are not proposed to tolerate forming such as that carried out when fabricating seamless tubes (forgeability properties).
JP-A-8 225 833 also proposes adding Co, but concerns a heat treatment to reduce the amount of residual austenite and not a chemical composition; the chemical composition ranges are thus broad and a teaching for the envisaged use cannot be deduced therefrom. Disclosure of the invention

The present invention proposes the production of a steel;
• with a creep strength at 600°C and 650°C at least equivalent to that of T92/P92 steel;
• with a hot oxidation resistance and steam corrosion resistance that is at least that of X20CrMoV12-l steel;
• which results in a lower production cost for seamless tubes compared with the improved grades cited above, the production cost being affected not only by that of the addition elements but also by that for transformation into seamless tubes.
We have also strived to produce a steel of the invention that allows the fabrication of small or large diameter seamless tubes using a variety of known hot rolling processes such as the Stiefel plug mill, MPM, pilger mill, push bench, continuous rolling mill with stretch reducing-mill. Axel rolling mill or planetary rolling mill processes.
In accordance with the invention, the steel under consideration contains, by weight:
C 0.06% to 0.20%
Si 0.10% to 1.00%
Mn 0.10% to 1.00%
S 0.010% or less
Cr 10.00% to 13.00%
Ni 1.00% or less
W 1.00% to 1.80%
Mo such that (W/2 + Mo) is 1.50% or less
Co 0.50% to 2.00%
V 0.15% to 0.35%
Nb 0.030% to 0.150%
N 0.030% to 0.120%

0.0010% to 0.0100%
and optionally, at most 0.050% by weight of Al and at most 0.0100% by weight of Ca.
The remainder of the chemical composition of said steel is constituted by iron and impurities or residual elements resulting from or necessary to steelmaking and casting.
Preferably, the amounts of the constituents of the chemical composition are linked so that after normalization heat treatment between 1050°C and 1080°C and tempering, the steel has a tempered martensitic structure that is free of or almost free of 5 ferrite.
The elements in the chemical composition of the steel have the following influence on the properties: CARBON
At high temperatures, in particular during the process for hot fabrication of metal products or during austenitization in the final heat treatment, said element stabilizes the austenite and as a result, tends to reduce the formation of 5 ferrite.
At ambient temperatures or at service temperatures, the carbon is in the form of carbides or carbonitrides the initial distribution and the change in said distribution of which with time act on the mechanical characteristics at ambient temperature and at the service temperature.
A C content of less than 0.06% would render obtaining a structure free of 5 ferrite and the production of the desired creep characteristics difficult.
A C content of more than 0.20% is deleterious to the weldability of the steel.
A content range of 0.10-0.15% is preferred. SILICON
This element is an element that deoxidizes liquid steel and also limits the kinetics of hot oxidation by air or steam in particular, according to the inventors, acting in synergy with the chromium content,
A content of less than 0.10% of Si is insufficient for producing said effects.

In contrast, Si is a ferrite forming element which has to be limited to avoid the formation of 5 ferrite and it also tends to encourage precipitation of embrittling phases in service. For this reason, its content is limited to 1.00%.
A content range of 0.20% to 0.60% is preferred. MANGANESE
This element encourages deoxidation and fixes the sulphur. It also reduces the formation of 5 ferrite.
In an amount of over 1.00%, however, it reduces the resistance to creep rupture.
A content range of 0.15% to 0.50% is preferred. SULPHUR
This element essentially forms sulphides which reduce the impact properties in the transverse direction and forgeability.
An S content limited to 0.010% prevents the formation of defects when hot piercing billets during the fabrication of seamless tubes.
A content that is as low as possible, for example 0.005% or less, or even 0.003% or less, is preferred. CHROMIUM
This element is found both dissolved in the steel matrix and precipitated in the form of carbides.
A minimum Cr content of 10% and preferably 11% is necessary for the hot oxidation behaviour.
Because of the ferrite forming nature of chromium, a content of more than 13% makes avoiding the presence of 5 ferrite difficult. NICKEL

This encourages impact strength and prevents the formation of 6 ferrite, but substantially reduces the Acl temperature and thus reduces the maximum tempering temperature of the steel.
Thus, a content of more than 1% is undesirable; moreover, nickel tends to reduce the creep rupture strength.
Preferably, the maximum Ni content is limited to 0.50%. TUNGSTEN
This element, which is both dissolved and precipitated in the form of carbides and intermetallic phases, is fundamental to the creep behaviour at 600°C and above, hence the minimum content of 1.00%.
However, this element is expensive, highly segretative and ferrite forming, and tends to form embrittling intermetallic phases.
The inventors have discovered that it is not advisable to increase the W content beyond 1.80%. MOLYBDENUM
This element has an effect similar to tungsten even though it appears to be less effective as regards creep strength.
Its effects add to that of tungsten and so the (W/2 + Mo) content is advantageously limited to 1.50%.
The molybdenum content is preferably 0.50% or less. COBALT
This element stabilizes austenite and thus enables more than 10% Cr to be tolerated; it also improves the creep strength properties; a minimum content of 0.50% is thus desirable.
In contrast, this element contributes to forming embrittling intermetallic compounds that can precipitate at the service temperature; further, it is very expensive.

Until now, this element has been used in contents of more than 2% in materials for use at high temperatures to improve their creep rupture strength.
The inventors of the present invention have surprisingly established that a range of cobalt contents of 0.50% to 2.00% and preferably 1.00% to 1.50% can satisfy the aims for said steel and in particular provide an optimum compromise between the various, possibly contradictory characteristics (for example oxidation resistance, creep strength and forgeability), using a relatively simple metallurgy and a limited manufacturing cost for metal products.
This is not the case with steels containing more than 2% of Co, which until now have not been used. VANADIUM
This element forms nitrides and carbonitrides that are very fine and stable and thus very important for the creep rupture strength.
A content of less than 0.15% is insufficient for producing the desired result.
A content of more than 0.35% is deleterious as regards the risk of the appearance of 6 ferrite.
A preferred range is from 0.20% to 0.30%. NIOBIUM
Like vanadium, this element forms stable carbonitrides and its addition reinforces the stability of vanadium compounds.
A Nb content of less than 0.030% is insufficient.
A Nb content of more than 0.15% is not favorable as the Nb carbonitrides may become too large and reduce the creep resistance.
A preferred range is from 0.050% to 0.100%. NITROGEN
This austenite forming element can reduce the appearance of 6 ferrite.

It can also, and especially, form very fine nitrides and carbonitrides which are much more stable than the corresponding carbides.
A minimum nitrogen content of 0.030% is therefore stipulated.
A nitrogen content of more than 0.120% results in blow holes in ingots, billets or slabs in the steels under consideration and as a result to defects in the metal products. The same risk exists on welding when processing said products.
A nitrogen content range of 0.040% to 0.100% is preferred. BORON
This element contributes to stabilizing carbides when added in an amount in excess of 0.0010%.
A content of more than 0.0100% can, however, substantially reduces the burning temperature of products, in particular of as cast products, and thus is detrimental. ALUMINIUM
This element is not necessary per se to produce the desired metallurgical characteristics and it is considered here as a residual; its addition is thus optional.
It is a powerful metal and slag deoxidant and can thus allow rapid, effective desulphurization of the steel by metal-slag exchange.
This element is also ferrite forming and scavenges nitrogen; thus, Al contents of more than 0.050% are discouraged.
Depending on requirements, if necessary, aluminium can be added to obtain a final content of up to 0.050%. CALCIUM
A Ca or Mg content of less than 0.0010% results from exchanges between liquid steel and slag containing lime or magnesia in a highly deoxidized medium: they are thus inevitable
steelmaking residuals.

However, calcium can optionally be added in amounts of a little over 0.0010% to improve castability and/or control the form of oxides and sulphides.
A Ca content of more than 0.0100% denotes an oxygen-rich and therefore dirty steel and is thus discouraged. OTHER ELEMENTS
Apart from iron, which is the base constituent of steel, and the elements indicated above, the steel of the invention only contains other elements as impurities; examples are phosphorus and oxygen, and residuals deriving mainly from the iron added to the furnace to produce the steel or from exchange with the slag or refractories or necessary to the steelmaking and casting processes.
Ti or Zr contents of less than 0.010% result thus from the fumaced scrap and not from any deliberate addition; such low contents actually have no substantial effect on the steel for the use tmder consideration.
Preferably, as regards forgeability care is taken that the copper content (resulting from fumaced scrap and not from deliberate addition) remains less than 0.25% and optionally less than 0.10%. Contents of more than said contents may proscribe certain hot rolling processes for seamless tube rolling and require the use of more expensive glass extrusion processes. CHEMICAL COMPOSITION RELATIONSHIP AND 8 FERRITE CONTENT
Steelmakers know how to equilibrate the chemical composition of a steel containing about 12% Cr, aiming at an absence or near absence of 6 ferrite after heat treatment from a relationship between the contents of the elements in the chemical composition. The term "structure almost free of 5 ferrite" means a structure containing no more than 2% of 5 ferrite and preferably no more than 1% of 5 ferrite (measured with an absolute precision of ± 1%).
One example of such a relationship is given below, but any relationship that is in the public domain or otherwise can be used, providing it has the desired effect.

An example is the Shaeffler diagram or diagrams derived therefrom which in particular incorporate the influence of nitrogen (De Long diagram) and the parameter Md derived from electronic orbital studies mentioned by Ezaki et al (Tetsu-to-Hagane, 78 (1992), 594). BRIEF DESCRIPTION OF THE DRAWINGS
The accompanying drawings illustrate a non-limiting example of an implementation of the invention.
Figure 1 shows a diagram of 5 ferrite content against equivalent chromium content for different specimens of heat treated steels containing 8% to 13% of Cr.
Figure 2 shows a diagram of the results of forgeability tests on steel F in accordance with the invention compared with other steels.
Figure 3 shows, for the same steel F compared with other steels, a diagram of hot tensile tests. Figure 3a) relating to the yield point and Figure 3b) to the tensile strength.
Figure 4 shows, for the same steel F compared with other steels, a transition curve for the Charpy V-notch impact strength test.
Figure 5 shows, for the same steel F compared with other steels, a graph of results of creep rupture strength tests under a constant unit load.
Figure 6 shows, for the same steel F compared with other steels, a master curve for the results of creep rupture strength tests under different unit loads as a function of the Larson-Miller parameter.
1ST EXAMPLE: tests on experimental heat
A 100 kg laboratory heat formed from the steel of the invention was produced under vacuum (F).

Figure 1 shows the relationship between an equivalent chromium parameter (Crequ) derived from the chemical composition and the 6 ferrite content: Crequ = Cr + 6Si + 4Mo + 1.5 W + H V + 5Nb + 8Ti - 40C - 30N - 2Mn - 4Ni - 2Co - Cu
The parameter Crequ derives from studies by Patriarca et al (Nuclear Technology, 28 (1976), p 516).
In Figure 1, we show the 5 ferrite content measured by image analysis in the optical microscope for a certain number of heats of T91, P91, T92 and X20 as a function of the parameter Cfequ-
Figure 1 provides analytical evidence that the amounts of elements in heat F lie within the ranges given in the chemical composition defined in claim 1. We aimed to obtain a Crequ content of 10.5% or less and if possible 10.0% or less to seek to remain substantially free of 6 ferrite (less than 2% and preferably less than 1%) after heat treatment.

Table 1 shows the chemical composition of this heat F and the mean chemical composition of known prior art grades (weight %) as well as the corresponding value of the parameter Crequ.
Said heat F contains no added Ca and its Al content is less than 0.010% (Al and Ca as residuals).

The ingots obtained were heated to 1250°C then hot rolled to a 20 mm thick sheet which then underwent stress-relieving tempering.
The specimens for the tests and examinations described below were produced from this sheet.
Firstly, a metallographic specimen taken in the longitudinal direction from said sheet was examined under the optical microscope after metallographic attack using Villela's reagent.
The presence of 6 ferrite was observed in the form of short white filaments in zones segregated into ferrite forming elements (Cr, W, Mo...). Its content was determined using automatic image analysis as 0.50%, i.e., an amount of almost zero.
Specimens were then taken from the transverse direction to carry out hot tensile forging tests at a mean deformation speed of 1 s-1
The forging tests were carried out comparatively on these specimens of heat F and on specimens from a rolled 310 mm diameter bar in P91 steel and from a rolled 230 mm diameter bar in P92 steel.
Figure 2 shows the reduction in area results.
It can be seen that the reduction in area remained over 70% from 1200°C to 1320°C and was comparable to that of P92.
Such behaviour was attributed to the low sulphur content of heat F and a relatively low 5 ferrite content at said temperatures.
The influence of temperature on the 5 ferrite content was also verified by metallographic tests: see Table 2.

The values for the 5 ferrite content obtained were comparable with those measured under the same conditions for comparative steels P91, P92.

The 5 ferrite content was less than 15% up to 1250°C and less than 20% up to 1280°C.
The limited 5 ferrite content in heat F at high temperature probably resulted from the deliberate absence of 5 ferrite at ambient temperatures.
The burning temperature was over 1320°C.
Thus, satisfactory behaviour can be expected for material F during hot piercing of round bars (termed rounds for tubes) between rolls using the Mannesmann process if heating of the rounds is limited to less than 1300°C and if possible to 1250°C.
Thus, it should be possible to produce seamless tubes by a number of hot rolling processes and thus it should be possible to produce them at relatively low cost. This is not the case for austenitic grades or grades containing 12% Cr and 1% Cu which, at least for small diameter tubes of the superheater tube type, have to be produced using the less productive glass extrusion process.
Dilatometric specimens were then taken from steel F of the invention and the steel transformation points on heating (Acl, Ac3) and cooling (Ms, Mf) were determined by dilatometry.
Table 3 shows the results obtained compared with typical results for known steels.

Temperature Acl of 830°C for steel F is comparable with that of P91 and P92 and much higher than that of P122 containing copper which does not allow a tempering temperature of more than 780°C. In contrast, a tempering temperature of 800°C is entirely possible with steel F of the invention.

Temperatures Ms and Mf at the beginning and end of the martensitic transformation remained sufficiently high for the transformation of austenite to martensite to be on cooling to ambient temperature.
The microstructure and hardness were measured after a normalizing heat treatment of 20 minutes at 1060°C (treatmentNl) or 1080°C (treatment N2); the results are shown in Table 4.

The microstructure and hardness were also measured after normalizing heat treatment Nl and tempering for 1 hour at 780°C (Tl), 30 minutes at 800°C (T2) or 1 hour at 800°C (T3): see the results shown in Table 5,

Note the fine austenitic grain size the dimensions of which did not exceed 0.030 mm. The tensile characteristics were then determined at ambient temperature and at 500°C and at 600°C - see the results in Table 6 and Figures 3a and 3b.

The Charpy V-notch impact strength characteristics were then measured in the longitudinal direction at test temperatures of -60°C to +40°C after heat treatments Nl+Tl, Nl-M2 or Nl+T3.
The results obtained and those on a tube with an outer diameter of 356 mm and wall thickness 40 mm in P92 are illustrated in Figure 4. The transition temperature for the Charpy V-notch impact strength was about 0°C for heat F, as for tubes P92.

The creep rupture strength characteristics were then determined using different tests at different temperatures under a constant unit load (140 and 120 MPa) compared with steel F of the present invention (heat treatments N1+T2 or N2+T2) and on a P92 tube.
The results of the stress rupture test at 120 MPa are shown in Figure 5 as a function of the parameter 1000/T (in °K-1), as is conventional for this type of grade. The temperatures were selected so that the maximum duration of the test was close to 4000 h. Figure 5 allows the temperature corresponding to a test duration of 105 h to be extrapolated for a unit load. It can be seen that for steel F, this temperature at least equals if not exceeds that of steel P92.
Other creep rupture strength tests at constant temperature were also carried out or are still running at 600°C, 625°C, 650°C.
The results of these tests (and those under a constant unit load) are shown in Figure 6 in the form of a diagram (master curve) showing log GR as a function of the Larson-Miller parameter (LMP) which combines the duration and temperature of the test: LMP = 10-3,T,(c+log IR) where c = 36 and T and tR are respectively expressed in °K and hours.

The ruptured tests reached a duration of 7800 h at 600°C, 10000 h at 610°C, 7800 h at 625°C and 7200 h at 650°C; the arrow on the diagram indicates a test at 600°C that had still not been ruptured after 11000 h.
Figure 6 shows that the tests are favourable compared with the mean master curve (solid line) and the lower scatter band (dotted line) for steels T92 and P92 defined by ASME.
Hot oxidation tests in steam were undertaken for product F in the N1+T2 temper at 600°C and 650°C for periods of up to 5000 hours compared with different steels for high temperature use according to ASTM A213 or DIN 17175:
• T22, T23 at low Cr contents (2.25%);
• T91,T92 at 9%Cr;
• X20, T122 at about 11 % Cr;
• TP347H (austenitic grade, 18% Cr-10% Ni-Nb).
Intermediate weight gain results, measured by weighing after 1344 h (8 weeks), are shown in Table 7.
The results are coded as follows:
• 1: weight gain of 2 mg/cm or less;
• 2: weight gain in the range 2 to 5 mg/cm2;
• 3: weight gain in the range 5 to 10 mg/cm2;
• 4: weight gain in the range 10 to 50 mg/cm ;
• 5: weight gain over 50 mg/cm2.
The X20 specimens could not be used for measurements due to major exfoliation of the oxide layers when leaving the furnace or during weighing (results shown in the Table as NA). In contrast, specimens of heat F and TP347H showed an absence of flaking of oxide layers. The fine crystallization of the oxidation products on heat F should also be noted.

These intermediate results allow it to be predicted, in particular at 650°C, that the steam oxidation behaviour of heat F of the invention will satisfy expectations, namely better than that for P91, P92 and at least equivalent to that of X20, or even close to that of TP347H.

The same specimens were removed after 5376 h and the loss of mass was measured after stripping off the oxides formed; this type of measurement is more accurate than weight gain measurements without stripping, but can only be carried out at the end of the test.
The table below summarizes the corrosion rates for the steel in mm/year, deduced from these measurements.
A test result order similar to that of Table 7 was found.
The corrosion rates for X20 and T122 (which contain 11% Cr) are not substantially different from those for T91 and T92, which contain only 9%.
In contrast, highly surprisingly, the corrosion rates for grade F of the invention were extremely low, lower even than for the austenitic steel specimen 347H containing 18% Cr and almost as low as for the 347 GF steel specimen (also austenitic, 18% Cr) which is a reference for hot oxidation behaviour.
The steel of the invention allows thus to produce boilers with a steam temperature of more than 600°C completely from ferritic steel, including the hottest parts of the boiler.

It should also be noted that the corrosion rates obtained for grade F were extremely low despite the very low sulphur contents, while certain prior art documents disclose moderate sulphur contents to combat hot oxidation, of the order of 0.005% or even 0.010%, and sulphur fixing by adding rare earths and/or alkaline-earths.
In contrast, grade F of the invention perfectly fits in with sulphur contents of 0.005% or less or even 0.003% or less, and does not necessitate the addition of rare earths and/or alkaline-earths which are difficult to implement. 2ND EXAMPLE: tests on industrial heat
An industrial heat labeled 53059 formed from grade F of the invention was produced (mass = 201) and cast into ingots.
The analysis for the heat was as follows.

Ingots were forged into solid bars with a diameter of 180 mm, which were then transformed into seamless tubes with an outer diameter of 60.3 mm and a thickness of 8.8 mm

using continuous rolling over a retained mandrel with diameter reduction on a stretch reducing-
This transformation into tubes was carried out without problems (no defects resulting from the presence of 5 ferrite) and the resulting tubes were of satisfactory quality according to non-destructive testing using ultrasonic waves.
Other ingots were transformed into large pipes with an outer diameter of 406 mm and a wall thickness of 35 mm using the hot pilger mill rolling process.
Here again, rolling was carried out without problems and no defects were observed during the inspection procedure.
These results confirm the expectations derived from the forgeability test results on the experimental heat (see Figure 2 and Table 2 above).
Table 10 shows the results of tensile tests at ambient temperature on tubes treated by normalization at 1060°C and tempering for 2 h at 780°C.
Table 11 shows the results of Charpy V-notch impact strength tests on tubes that underwent the same heat treatment as that for the tensile tests.

The mechanical traction and resilience characteristics for the tube were in line with the results for the bars from the experimental heat.

1. A steel for seamless tubular products intended for high temperature use, characterized
in that it contains, by weight:
C 0.06% to 0.20%
Si 0.10% to 1.00%
Mn 0.10% to 1.00%
S 0.010% or less
Cr 10.00% to 13.00%
Ni 1.00% or less
W 1.00% to 1.80%
Mo such that (W/2 + Mo) is 1.50% or less
CO 0.50% to 2.00%
V 0.15% to 0.35%
Nb 0.030% to 0.150%
N 0.030% to 0.120%
B 0.0010% to 0.0100%
and optionally, at most 0.050% by weight of Al and at most 0.0100% by weight of Ca;
the remainder of the chemical composition being constituted by iron and impurities or residual elements resulting from or necessary to steelmaking or casting.
2. A steel according to claim 1, characterized in that the amounts of the constituents of
the chemical composition are linked by a relationship such that after normalization
heat treatment between 1050°C and 1080°C and tempering, the steel has a tempered
martensitic structure that is free of or almost free of 5 ferrite.

3. A steel according to claim 1 or claim 2, characterized in that its Cr content is in the range 11.00% to 13%.
4. A steel according to any one of claims 1 to 3, characterized in that its Si content is in the range 0.20% to 0.60%.
5. A steel according to any one of claims 1 to 4, characterized in that its C content is in the range 0.10% to 0.15%.
6. A steel according to any one of claims 1 to 5, characterized in that its Co content is in the range 1.00% to 1,50%.
7. A steel according to any one of claims 1 to 6, characterized in that its Mo content is 0.50% or less.
8. A steel according to any one of claims 1 to 7, characterized in that its Mn content is in the range 0.10% to 0.40%.
9. A steel according to any one of claims 1 to 8, characterized in that its Ni content is 0.50% or less,
10- A steel according to any one of claims 1 to 9, characterized in that the residual elements are controlled so that the Cu content in the steel is 0.25% or less and preferably 0.10% or less.
11. A steel according to any one of claims 1 to 10, characterized in that its S content is 0.005% or less, and preferably 0.003% or less.

12. A steel for seamless tubular products substantially as herein described with reference to the accompanying drawings.



1561-chenp-2003-claims filed.pdf

1561-chenp-2003-claims granted.pdf






1561-chenp-2003-form 1.pdf

1561-chenp-2003-form 26.pdf

1561-chenp-2003-form 3.pdf

1561-chenp-2003-form 5.pdf

1561-chenp-2003-other documents.pdf


Patent Number 212223
Indian Patent Application Number 1561/CHENP/2003
PG Journal Number 07/2008
Publication Date 15-Feb-2008
Grant Date 26-Nov-2007
Date of Filing 01-Oct-2003
Name of Patentee V & M FRANCE
Applicant Address 130, rue de Silly, F-92100 Boulogne-Billancourt
# Inventor's Name Inventor's Address
1 ARBAB, Alireza 23, Avenue des Androuin, F-59300 Valenciennes
PCT International Classification Number C22C 38/00
PCT International Application Number PCT/FR2002/001151
PCT International Filing date 2002-04-03
PCT Conventions:
# PCT Application Number Date of Convention Priority Country
1 01/04551 2001-04-04 France